HIGH TEMPERATURE SHAPE MEMORY POLYMERS
&
IONOMER MODIFIED ASPHALTS
A Dissertation
Presented to
The Graduate Faculty of The University of Akron
In Partial Fulfillment
of the Requirements for the Degree
Doctor of Philosophy
Ying Shi
August, 2013
HIGH TEMPERATURE SHAPE MEMORY POLYMERS
&
IONOMER MODIFIED ASPHALTS
Ying Shi
Dissertation
Accepted: Approved:
______Advisor Department Chair Dr. Robert A. Weiss Dr. Robert A. Weiss
______Committee Member Dean of the College Dr. Sadhan C. Jana Dr. Stephen Z. D. Cheng
______Committee Member Dean of the Graduate School Dr. Kevin Cavicchi Dr. George R. Newkome
______Committee Member Date Dr. Matthew Becker
______Committee Member Dr. Yi Pang
ii
ABSTRACT
This dissertation consists of two research subjects: High Temperature Shape
Memory Polymers and Ionomer Modified Asphalts. Current development of thermally sensitive shape memory polymers (SMPs) has focused primarily on relatively low transition temperatures (Tc < 100°C) and elastomeric polymers, such as thermoplastic polyurethanes (TPU), crosslinked polyethylene, poly (ε-caprolactone), sulfonated EPDM and polynorbornene. Those materials are appropriate for applications such as biomedical and surgical materials, smart fabrics and heat shrinkable tubing. Materials used as aerospace or structural components often require higher modulus and switching temperatures for shape change and actuation. To the best of our knowledge, there have been no reports of thermoplastic SMPs with controllable switching temperatures above
100°C. There has been research on high temperature SMPs but based on thermoset polymer systems. High temperature thermoplastic shape memory polymers were developed from metal salts of sulfonated PEEK (M-SPEEK, M=Na+, Zn2+, Ba2+, Al3+,
Zr4+) ionomer and composites of the M-SPEEK ionomers with a fatty acid salt. M-
SPEEKs were prepared by neutralizing sulfonated PEEK acid to metal salts. The glass transition temperatures of M-SPEEK ionomers increased with increasing Coulomb energy of ion pairs and the ionomers were thermally stable to ~320°C. The M-SPEEK ionomers exhibited microphase separated morphologies and the average correlation length was determined by small angle X-ray scattering. Al-SPEEK and Zr-SPEEK showed crosslinked characteristics such as rubbery plateau above Tg and much reduced
iii water uptake. The M-SPEEK ionomers exhibited reasonable shape memory behavior, in which the permanent network was provided by ionic nanodomains formed by the interaction of ionic groups and glass transition temperatures served as the switching temperatures. The relative poor shape efficiency of Na-SPEEK and Zn-SPEEK (80-90%) can be improved by blending M-SPEEK with a low molar mass crystalline compound
NaOl. The composites were prepared from 70 wt% M-SPEEK (M = sodium or zinc) and
30 wt% sodium oleate (NaOl). Ionic nanodomains formed by the interactions of ionic groups provided a permanent physically crosslinked network and strong dipolar interactions between the ionomer and a dispersed phase of crystalline NaOl provided the temporary network. A temporary shape was achieved and fixed by deforming the material above the melting temperature (Tm) of NaOl and then cooling under stress to below Tm. The permanent shape was recovered by reheating the material above Tm without applying stress. Shape fixing efficiencies of 96% were achieved and shape recovery reached 100%. Triple shape memory behavior was also achieved for M-
SPEEK/NaOl compounds using the glass transition of the ionomer and the melting point of the NaOl as two separate switching temperatures.
Asphalt binders suffer from different kinds of distresses such as low temperature cracking, rutting and fatigue during “in life” service. In the regions with cold climate, thermal cracking of pavement occurs when low temperature shrinkage exceeds the ability of stress dissipation by the asphalt binders and leads to brittle fracture of the glassy pavement. The objective of the ionomer modified asphalt project is to improve the elasticity of the asphalt binder at low temperature in order to prevent the cracking. A
iv performance grade 64-28 asphalt and partially neutralized copolymers of ethylene and methacrylic acid ionomer were mixed at four concentration levels (0-9 wt%) to yield ionomer modified asphalt blends.
The thermal properties, morphology and viscoelastic behavior of ionomer modified asphalts were studied. The ionomer modified asphalt exhibited much better dispersion and smaller phase separation than did polyethylene modified asphalt. After establishing the linear viscoelastic range of response through strain sweep, frequency sweep tests at a temperature range of 30-80C were conducted to study the dynamic mechanic properties of the modified blends. The isothermal response curves were reduced to dynamic master curves of modulus and viscosity based on the time- temperature superposition principle. The effects of ionomer concentration and mixing time on the viscoelastic behavior were studied. The addition of ionomer improved the elasticity of the asphalt, but long times were needed to mix the ionomer into the asphalt and properties were very sensitive to mixing time. A series of SuperPave tests were conducted on both ionomer modified and neat asphalt, which simulating the real life temperature and traffic load condition. The performance grade of ionomer modified asphalt was transformed from 64-28 to 69.2-26.5.
v
ACKNOWLEDGEMENTS
The completion of this dissertation took place with the help and encouragement of many people. Like they say, “it takes a village to raise a child”, I believe it took just as many people for my dissertation, and I owe thanks to them. I would like to express my deepest appreciation to my advisor Dr. Robert Weiss., who has the attitude and a substance of a genius. He is the best advisor I could ever ask for. Thank you for training to be a good scientist.
I would like to thank Dr. Mitra Yoonessi for her support and guidance through the
High T SMP project. Thank you for giving me the opportunity to explore other high T polymer project from OAI and thank you for your suggestions and assistance for job search.
I also like to thank Dr. Montgomery T. Shaw and Dr. Richard Parnes at the
University of Connecticut. Thank you for helping me understanding rheology and discussing my IMA project.
I also like to give thanks to all members of my dissertation committee for their time and support: Dr. Kevin Cavicchi, Dr. Sadhan Jana, Dr. Matthew Becker, and Dr. Yi Pang, thank you.
My group members and colleagues in the Polymer Engineering Department (Univ.
Akron) and Institute of Materials Science (Univ. Conn.) have been a critical source of support and encouragement. I would like to thank my previous group member Dr. Jing
vi
Dong, Dr. Gerald Ling, Dr. Emmanuel Pitia, Dr. Xueyuan Wang, Dr. Hui Niu, and, Dr.
Jinkun Hao for discussing my research problems. I also like to thank the current group members: Sahil Gupta, Longhe Zhang, Xing Lu, Zhiyang Zhao, and Murat Bakan. It was a great pleasure to work with you.
Last but not the least I want to give much thanks to my mom, Lan Yang for her great support, and love.
vii
TABLE OF CONTENTS
Page LIST OF TABLES ...... xivi
LIST OF FIGURES ...... xv
CHAPTER
I. INTRODUCTION: HIGH TEMPERATURE SHAPE MEMORY POLYMERS ...... 1
1.1 Shape Memory Polymers...... 1
1.2 Shape Memory Polymers versus Shape Memory Alloys ...... 2
1.3 Thermal responsive shape memory polymers ...... 3
1.4 Research Objective ...... 5
II. BACKGROUND AND LITERATURE REVIEW HIGH TEMPERATURE SHAPE MEMORY POLYMERS ...... 6
2.1 Fundamental aspects of shape memory polymers ...... 6
2.1.1 Thermodynamic point of view ...... 7
2.1.2 Molecular mechanism of the shape memory effects ...... 8
2.2 Classification of Shape Memory Polymers ...... 10
2.2.1 Physically cross-linked glassy networks ...... 10
2.2.2 Physically cross-linked semi-crystalline networks ...... 13
2.2.3 Chemically cross-linked glassy networks ...... 15
2.2.4 Chemically cross-linked semi-crystalline networks ...... 17
2.2.5 Other thermoplastic networks ...... 18
2.3 Characterization Methods ...... 20
viii
2.3.1 Thermal properties characterization ...... 20
2.3.2 Viscoelastic properties characterization ...... 21
2.3.3 Morphology characterization ...... 22
2.3.4 Shape Memory Performance Characterization ...... 23
2.4 Triple shape memory effect and characterization ...... 26
2.5 Different stimuli responsive shape memory polymers ...... 30
2.5.1 Indirect actuation ...... 30
2.5.2 Light activated shape memory polymer ...... 32
2.6 Biomedical Applications ...... 33
2.7 Innovation of this research ...... 36
III.SYNTHESIS, CHARACTERIZATION, AND SHAPE MEMORY BEHAVIOR OF M-SPEEK IONOMERS...... 41 3.1 Introduction ...... 41
3.2 Experimental Details ...... 46
3.2.1 Materials ...... 46
3.2.2 Sulfonation ...... 46
3.2.3 Neutralization ...... 47
3.2.4 Water uptake measurements ...... 48
3.3 Materials characterization ...... 49
3.3.1 Titration...... 49
3.3.2 Structure Characterization ...... 50
3.3.3 Thermal Analysis ...... 50
3.3.4 Mechanical Properties and shape memory cycle ...... 51
3.3.5 Rheology ...... 52
viii
3.4 Results and Discussion ...... 52
3.4.1 Sulfonation ...... 52
3.4.2 Neutralization ...... 55
3.4.3 Morphology characterization of M-SPEEK...... 60
3.4.4 Thermal properties ...... 63
3.4.5 Linear Viscoelasticity ...... 68
3.4.6 Dynamic shear ...... 71
3.4.7 Shape memory behavior ...... 77
3.5 Conclusions ...... 84
IV. DUAL AND TRIPLE SHAPE MEMORY BAHAVIOR OF M-SPEEK/FATTY ACID SALT COMPOUNDS ...... 85
4.1 Introduction ...... 85
4.2 Experimental Details ...... 86
4.2.1 Preparation of M-SPEEK/NaOl compounds ...... 86
4.2.2 Thermal analysis ...... 86
4.2.3 Wide angle X-ray diffraction (WAXD) ...... 86
4.2.4 Mechanical tests ...... 87
4.2.5 Dual shape memory cycle test ...... 87
4.2.6 Triple shape memory cycle test ...... 88
4.3 Results and discussion ...... 89
4.3.1 Thermal properties ...... 89
4.3.2 Wide angle X-ray diffraction ...... 92
4.3.3 Mechanical and Viscoelastic Properties ...... 93
4.3.4 Shape memory behavior ...... 98
ix
4.3.5 Triple shape memory behavior ...... 104
4.4 Conclusions ...... 109
V. IONOMER MODIFIED ASPHALT ...... 111
5.1 Asphalt properties ...... 111
5.1.1 Chemical composition of asphalt ...... 111
5.1.2 Physical properties of asphalt ...... 113
5.2 Pavement Problems ...... 114
5.3 Polymer Modified Asphalt (PMA) ...... 116
5.3.1 Thermoplastic polymer modification ...... 118
5.3.2 Polyolefin modification ...... 121
5.3.3 Reactive polymers modification ...... 124
5.4 Research Objective ...... 125
5.5 Superpave Test ...... 126
5.5.1 Rotational Viscometer test ...... 128
5.5.2 Dynamic Shear Rheometer test ...... 129
5.5.3 Bending Beam Rheometer test...... 131
5.5.4 Rolling thin film oven aging (RTFO) ...... 132
5.5.5 Pressure Aging Vessel aging (PAV) ...... 133
5.6 Asphalt binder grade ...... 134
5.7 Innovation of this research ...... 135
VI. STRUCTURE AND PROPERTIES OF IONOMER MODIFIED ASPHALTS ..... 138
6.1 Introduction ...... 138
6.2 Experiment details ...... 141
6.2.1 Preparation of asphalt and ionomer blends ...... 141
x
6.2.2 Materials Characterization ...... 141
6.3 Results and discussion ...... 143
6.3.1 Thermal behavior of ionomer modified asphalt (IMA) ...... 143
6.3.2 Morphology of ionomer modified asphalt ...... 145
6.3.3 Viscoelastic behavior of IMAs ...... 147
6.3.4 Time Temperature Superposition ...... 153
6.3.5 SuperPave test ...... 163
6.4 Conclusions ...... 168
VII.SUMMARY AND FUTURE WORK ...... 170
BIBLIOGRAPHY ...... 176
xi
LIST OF TABLES
Table Page
3.1 q/a value for different cations ...... 45
3.2 Thermal Properties of SPEEK with different reaction conditions ...... 54
3.3 S-O stretching bands of -SO3¯ groups in M-SPEEK ...... 57
3.4 Wake uptake property of M-SPEEK...... 60
3.5 Electron density of PEEK, H-SPEEK and M-SPEEK ...... 63
3.6 Glass transition temperature of M-SPEEK ...... 65
3.7 Summary of the shear modulus ( G ′ ) at frequency=1 rad/s and molecular weight between crosslinks ( Mc ) at 270 and 300 °C...... 73
3.8 Shift factors of Na-SPEEK at 280°C ...... 75
3.9 Shape memory properties of M-SPEEKs ...... 84
4.1 Thermal characteristics of materials ...... 90
4.2 Engineering tensile propertiesa of M-SPEEK and M-SPEEK/NaOl(30) Compounds 94
4.3 Shape fixing and recovery efficiencies of M-SPEEKs ...... 99
4.4 Shape memory properties for M-SPEEK/NaOl(30) compounds ...... 103
4.5 Shape fixing and recovery efficiencies three consecutive triple shape memory cycles for Zn-SPEEK/NaOl(30)...... 107
4.6 Shape fixing and recovery efficiencies three consecutive triple shape memory cycles for Na-SPEEK/NaOl(30)...... 108
6.1 Thermal properties of IMAs ...... 144
xiv
6.2 Effect of mixing time at 180°C on η* [Pa⋅s] at 40°C and f = 1 rad/s ...... 147
6.4 Zero shear rate Viscosity ionomer modified asphalt ...... 153
6.5 WLF constants and slope of viscoelastic master curves ...... 162
6.6 Super Pave Test results for ionomer modified and neat asphalt* ...... 166
6.7 Performance Grade Determination of IMA* ...... 168
xiv
LIST OF FIGURES
Figure Page
1.1 Schematic illustration of dual shape memory behavior ...... 4
2.1 Molecular mechanism of thermally induced shape memory polymer. (a) A multiblock copolymer with Tc=Tm, (b) A covalent crosslinked polymer with Tc=Tm, (c) A chemically crosslinked polymer network with Tc=Tg. (Reprinted with permission from [1] )...... 9
2.2 Schematic illustration of DSC curve for SMPs (a) Amorphous polymer network Ttran=Tg; (b)Semi-crystalline polymer network Ttran=Tm.Reprined with permission from [73] ...... 21
2.3 viscoealstic behavior of different classes SMPs. (I) chemically crosslinked amorphous copolymer network Ttran=Tg; (II) Chemcially crosslinked semi-crystalline polymer network Ttran=Tm; (III) Physically crosslinked thermoplastic Ttran=Tg;(IV) Physically crosslinked thermoplastic Ttran=Tm. Reprinted with permission from [73] ...... 22
2.4 One ideal shape memory cycle. S and E denote the start and end of the cycle. The sample is heated without applying any stress along path 1. Certain stress is applied to stretch the sample above Tc. The sample is then cooled under stress along path 3 and the stress is removed along path 4. The sample recovers its permanent shape when reheated above Tc (path 5). The sample can then be cooled to end the cycle. Reprinted with permission from [74]...... 24
2.5 Schematic illustration of recovery curves of SMPs (a) recovered under constant strain conditions; (b) recovered under stress free conditions. Reprinted with permission from [73] ...... 25
2.6 Illustration of triple shape memory behavior Reprinted with permission from [76] .. 28
2.7 One shape memory cycle for the multiphase network composed of crystalline PCL segments and amorphous PCHMA segments. The solid line indicates strain, and the dashed line indicates temperature. The material has a permanent shape C and was first deformed to temporary shape B (εB) at 150°C (T> Tg, PCHMA) where the material is in the rubbery-elastic state. The samples was then cooled to Tm,PCL xv where the second temporary shape A recover to first temporary shape B. The recovery from temporary shape B to permanent shape C was obtained by further heating the temperature to 150°C. Reprinted with permission from [77] ...... 29 2.8 Illustration of shape recovery under infrared radiation. Radiation exposure is from theleft. Reprinted with permission from [5] ...... 30 2.9 Illustration of shape recovery of magnetically induced shape memory effect. Reprinted with permission from [80]...... 31 2.10 Water activated shape memory polymer recovery process Reprinted with permission from [81] ...... 32 2.11 Molecular mechanism illustration for light induced shape memory polymer. Photoreversible crosslinks (filled diamonds) are formed during fixation of the temporary shape by UV light. Reprinted with permission from [3] ...... 33 2.12 Illustration of degradable shape memory suture for wound closure. Reprinted with permission from [15]...... 34 2.13 Shape recovery of a transparent shape memory polymer stent to the permanent unfolded shape. Reprinted with permission from [92] ...... 36 3.1 Illustration of ionic interactions between metal ions and sulfonate group ...... 45 3.2 Schematic illustration of sulfonation and neutralization reaction...... 47 3.3 FTIR spectrum of M-SPEEK ...... 58 3.4 SAXS of H-SPEEK and M-SPEEK at room temperature ...... 62 3.5 DSC for H-SPEEK and M-SPEEK. Second heating curve are showing, heat rate was 10°C/min ...... 65 3.6 TGA of PEEK, HPEEK, and M-SPEEK. Samples were heated under nitrogen environment with a heating rate of 10°C/min ...... 67 3.7 TGA for Zn-SPEEK. Sample was held at 270°C for 60mins and less than 1% weight loss was observed (inset picture)...... 67 3.8 Storage modulus of M-SPEEK as a function of temperature ...... 68 3.9 Loss modulus of M-SPEEK as a function of temperature ...... 69 3.10 Tanδ of PEEK and M-SPEEK as a function of temperature ...... 70 xvi 3.11 Storage Modulus of M-SPEEKs as a function of frequency ...... 72 3.12 Loss Modulus of M-SPEEKs as a function of frequency ...... 74 3.13 Storage modulus of Na-SPEEK as a function of frequency at different temperatures ...... 75 3.14 Loss modulus of Na-SPEEK as a function of frequency at different temperatures .. 76 3.15 G' and G" Master curve for Na-SPEEK. Reference temperature is 280°C ...... 76 3.16 Schematic illustration of shape memory test programming ...... 77 3.17 Shape memory cycle for PEEK. Point S denotes the start of the cycle. Path 1: The sample was heated to 200°C (Tc ~150°C) and stretched with a stress of 3.8 MPa. Path 2: The deformed sample was cooled under a constant stress of 3.8 MPa. Path 3: The stress was to set the temporary deformed shape. Path 4: The sample was reheated to 200°C to recover the permanent shape. Point E denotes the end of the cycle...... 78 3.18 Shape memory cycle for H-SPEEK. The samples were stretched at 220°C...... 79 3.19 One shape memory cycle of Zn-SPEEK, sample was stretched at 270°C...... 81 3.20 Intermolecular inter-actions that provide physical crosslinks in ionomers...... 81 3.21 One shape memory cycle of Al-SPEEK, sample was stretched at 300°C ...... 82 3.22 One shape memory cycle of Zr-SPEEK, sample was stretched at 300°C ...... 83 3.23 Shape fixing and recovery ratio for M-SPEEKs (●) R; (▲) F ...... 83 4.1 TGA curves for PEEK, H-SPEEK, M-SPEEK and M-SPEEK/NaOl(30). Experiments were run using a nitrogen atmosphere ...... 91 4.2 WAXD for M-SPEEK and M-SPEEK/NaOl ...... 92 4.3 Engineering tensile stress versus strain curves at room temperature for neat PEEK, M-SPEEK and M-SPEEK/NaOl(30) compounds...... 93 4.4 Dynamic and loss tensile modulus versus temperature for (a) PEEK, ZnSPEEK, and Zn-SPEEK/NaOl(30); (b) NaSPEEK and Na-SPEEK/NaOl (30). The frequency was 1 Hz...... 95 4.5 Four consecutive shape memory cycles for Na-SPEEK. The samples were stretched at 270°C. (Tc = 250°C). The numbers denote the cycle number...... 99 xvii 4.6 Four consecutive shape memory cycles for Zn-SPEEK. The samples were stretched at 270°C. (Tc = 250°C). The numbers denote the cycle number...... 100 4.7 Illustration of molecular change after M-SPEEK blending with NaOl ...... 101 4.8 Four consecutive shape memory cycles for a Zn-SPEEK/NaOl composite film. The samples were stretched at 270°C. The numbers denote the cycle number...... 102 4.9 Four consecutive shape memory cycles for a Na-SPEEK/NaOl composite film. The samples were stretched at 270°C. The numbers denote the cycle number...... 103 4.10 Schematic illustration of a triple shape memory cycle programming, where Tg is from M-SPEEK and Tm is from NaOl ...... 105 4.11 Consecutive tripe shape memory cycles for Zn-SPEEK/NaOl(30%)...... 106 4.12 Consecutive tripe shape memory cycles for Na-SPEEK/NaOl(30%) ...... 108 5.1 (a)Low temperature cracking, (b)rutting and (c)fatigue cracking of pavements. Preprinted with perform from [184] ...... 116 5.2 Schematic illustration of a network in SBS modified asphalt Reprinted with permission from [198]...... 119 5.3 Penetration grading and viscosity grading. Reprinted with permission from [227] . 127 5.4 Schematic illustration of Superpave test types and temperature ranges. Reprinted with permission from [227]...... 128 5.5 Schematic illustration of rotational viscosity test. Reprinted with permission from [227]...... 129 5.6 Schematic illustration of DSR measurement. Reprinted with permission from [227] ...... 130 5.7 Schematic illustration of BBR test. Reprinted with permission from [227] ...... 132 5.8 Rolling thin film oven (left) and RTFO samples (right). Reprinted with permission from [227] ...... 133 6.1 Schematic of ionomer nano-structure The blue nanophase constitutes ion-rich domains dispersed in a continueshydrophobic matrix...... 139 6.2 DSC of base asphalt and ionomer modified asphalt ...... 144 6.3 TGA of base asphalt and ionomer modified asphalt ...... 145 xviii 6.4 Optical polarized micrographs of asphalt modified with (a) 5 wt% LDPE, (b,c) 5 wt% ionomer ...... 146 6.5 Dynamic viscosity of base asphalt and inomer modified asphalts at 40C and 60C :a 45minutes mixing ; b, 60minutes mixing ...... 149 6.6 Elastic modulus of neat and inomer modified asphalt at 40C and 60C:a 45minutes mixing ; b, 60minutes mixing ...... 150 6.7 viscosity data fitted to Carreau model (5% ionomer 60minutes mixing) ...... 152 6.8 Shear Modulus of neat asphalt versus frequency at different temperatures (30-80°C ) (a) Storage modulus and (b) loss modulus. Numbers denote temperatures (°C) ...... 154 6.9 Shear Modulus of 3 wt% IMA versus frequency at different temperatures (30-80°C ) (a) Storage modulus and (b) loss modulus. Numbers denote temperatures (°C) ...... 155 6.10 Shear Modulus of 5wt% IMA versus frequency at different temperatures (30-80°C). (a) Storage modulus and (b) loss modulus. Numbers denote temperatures (°C) ...... 156 6.11 Shear Modulus of 7wt% IMA versus frequency at different temperatures (30-80°C). (a) Storage modulus and (b) loss modulus. Numbers denote temperatures (°C) ...... 157 6.12 Shear Modulus of 9wt% IMA versus frequency at different temperatures (30-80°C). (a) Storage modulus and (b) loss modulus. Numbers denote temperatures (°C) ...... 158 6.13 Elastic modulus master curve for base asphalt and ionomer modified asphalt ...... 159 6.14 Loss modulus master curve for base asphalt and ionomer modified asphalt ...... 159 6.15 Shift factors (aT) vs. (T – To) used for the G’ and G” data shown in Fig. 12 for the asphalt and IMAs. The reference temperature, To, was 40°C. The curves are the WLF equation fits (see Table 6.5 for the WLF constants used)...... 161 6.16 Dynamic viscosity master curves at 40°C for base asphalt and the IMAs ...... 162 xix CHAPTER I INTRODUCTION: HIGH TEMPERATURE SHAPE MEMORY POLYMERS 1.1 Shape Memory Polymers Shape memory polymer is a class of smart materials that can change their shape in a predefined manner in response to an external stimulus [1]. A shape memory polymer can be deformed and fixed in to a temporary shape, which will remain stable until the material is exposed to a certain stimuli. This will trigger the material to recover to its original shape. A variety of external stimulus can be utilized, such as heat, light (UV and infrared light), chemical (solvent and pH change), electrical current/voltage and magnetic field [2-6]. This results in various applications, such as heat-shrinkable tubes for electronics or films for packaging, self-deployable structures and morphing wings and intelligent medical devices [7]. The “shape memory” effect in polymers was first mentioned in 1941 in a US patent, which claimed “elastic memory” property of a dental material made of methacrylic acid ester resin [8]. The well-known heat shrinkable tubing made from covalently cross-linked polyethylene in 1960s represents the first commercial application of shape memory polymers [9-13], even though the terminology had not been used. Despite the early discovery and application, not much fundamental inquiry was involved and scientific papers in this area were limited until 1990s. Significant efforts began since 1 the discovery of segmented polyurethane shape memory polymer by Mitsubishi Heavy Industries Ltd [14]. The research on shape memory polyurethanes remains quite active today. Recently, Lendlein et al [15] reported a self-tightening suture using shape memory polymers for minimum invasive surgery. This material is deformed in the form a suture as a temporary shape and applied loosely to the wound. The shape recovery is actuated by human body heat, leading to self-contract action and tightening the wound. This research has stimulated significant interests in the biomedical area application of shape memory polymers. 1.2 Shape Memory Polymers versus Shape Memory Alloys The shape memory effect was found in a gold-cadmium alloy [16] as early as 1932, but not much progress had been made until 1971, when a nikel-titanium alloy was observed to possess significant recoverable strain [17]. The shape memory effect in a shape memory alloy is based on phase transition from a high temperature cubic symmetry (austenitic phase) to a low temperature phase with lower symmetry (martensitic phase) [18]. The shape memory behavior can be achieved by thermal actuation. For the Ni-Ti alloy, the martensitic phase is not stable above a certain temperature without external force so that the material exhibits superelasticity and up to 8% strain recovery was observed. Today, Nitinol is still the most widely used shape memory alloy and a wide range of other shape memory alloys have been developed in the form of solid, film and even foams. There are three major classed of shape memory alloy that are commercial: NiTi-based, Cu-based (CuAlNi and CuZnAl) and Fe-based [19-21]. Due to the high performance and good biocompatibility of NiTi-based shape memory alloys, they are widely used in biomedical applications, for instance, stents and surgical devices [22]. The 2 major applications for shape memory alloys are in the field of medical device manufacturing. Other application areas are growing fast too, such as aircraft, automotive devices, and eyeglass frames [23]. Although shape memory alloys are more mature with regard to real engineering, they have several disadvantages. First, shape memory alloys have high manufacturing costs. Second, the recoverable deformations is limited, even though some of the shape memory alloys exhibit “superelasticity” which is 8%. Third, there is still difficulty in controlling micro-structural mechanism and physical behavior of shape memory alloys. Shape memory polymers have attracted increasing interest because of their intrinsic advantages. First, the cost of polymers, including materials cost and processing cost, is traditionally much lower. Second, shape memory polymer can achieve much larger strain, for instances, polymers with 800% attainable stain have been reported [24]. Third, tailoring the material properties by adjusting polymer structure is much easier than for metals and alloys. Recent research has shown that shape memory polymers can be triggered not only by heat, but by various stimuli and even multiple stimuli simultaneously [7]. There is more freedom to tailor shape memory polymers to meet the needs of certain application. 1.3 Thermal responsive shape memory polymers The most reported shape memory polymers so far are thermal responsive, e.g. use heat as a stimulus. This is a consequence of the intrinsic thermal phase transitions occurring in polymers, for example, the glass transition at which a material changes from a glassy state to a rubbery state, or the melting transition at which a crystalline material changes from a solid state to a melt liquid state. Thermally-actuated shape memory 3 polymers have either a covalently cross-linked or physical cross-linked structure as a permanent network and a low temperature reversible physical transition, e.g., vitrification, melting or microphase separation, as a temporary network [25]. The thermodynamic origin of shape memory is the shape change that accompanies a conformational entropy change at the transition temperature. Elastic strain energy is stored by the temporary network during shape fixing, while the shape recovery to original shape formed by the permanent network is due to the release of the stored energy. Figure 1.1 schematically illustrates a thermally induced shape memory behavior. The polymer has a permanent shape. During the programming process, the polymer is heated above a certain transition temperature, where it can be deformed to a temporary shape. This temporary can be fixed by cooling the material pass its transition temperature. When reheated, the polymer recovers to the permanent shape when the transition temperature is reached. Figure 1.1 Schematic illustration of dual shape memory behavior The development of thermally sensitive shape memory polymers has focused primarily on relatively low transition temperatures (Tc < 100°C) and elastomeric 4 polymers, such as thermoplastic polyurethane elastomers (TPU), crosslinked polyethylene, poly(ε-caprolactone), sulfonated EPDM and polynorbornene[26-30]. These materials are appropriate for applications such as biomedical and surgical materials, smart fabrics and heat shrinkable tubing [31]. 1.4 Research Objective Materials used as aerospace components often require higher switching temperatures for shape change and actuation. The state-of-the-art studies in the area of aerospace applications of shape memory polymers focus on shape memory polymer composites for deployable space structures [6]. The reinforced shape memory polymer composites possess enhanced improvement in material properties, especially modulus, strength and heat resistance. The polymer matrix investigated is usually polyurethane (PU), poly(ethyl methacrylate)(PEMA) or low density polyethylene (LDPE) [32]. Therefore, the switching temperature of these shape memory polymer composites is still relatively low. To the best of our knowledge, there have been no reports of thermoplastic SMPs with controllable switching temperatures above 200°C. There has been research on high temperature SMPs based on thermoset polymers systems [33-35]. The objective of this research is to develop a thermoplastic, high switching temperature shape memory polymer system with a switching temperature as high as 270°C that may be suitable for aerospace applications . 5 CHAPTER II BACKGROUND AND LITERATURE REVIEW HIGH TEMPERATURE SHAPE MEMORY POLYMERS 2.1 Fundamental aspects of shape memory polymers Shape memory polymers (SMPs) have obtained significant attention because of their scientific and industrial potentials. They have inherent advantages of being low cost and light weight. In addition, they provide much higher degree of deformation comparing to shape memory alloys and ceramics [16]. These give the material great potential for a variety of applications in smart medical devices, sensors, actuators and aircraft parts [36]. SMPs can be manipulated to a temporary shape by mechanical deformation at their low modulus state and they can fix this shape at their high modulus state. SMPs have the ability to memorize their permanent shape so that they can recover to their original shape under certain external stimulus. Heat activated SMPs have been widely investigated, because the intrinsic thermal transitions occurring in polymers, e.g. the glass transition at which a material can change reversibly from a glassy state to a rubbery state, or the melting transition at which a crystalline material loses its integrity and goes into a liquid state. During these thermal transitions, the material experiences low and high modulus states. Such a polymer system usually consists of two phases; one of them is a fixed phase and the other one is a switching phase. The permanent shape is provided by either physically (intermolecular 6 interactions) or chemically (covalent bonds) crosslinked network. This network is able to sustain during the deformation process where the material is in its low modulus state. The switching phase can either be amorphous segments, where the transition temperature (Tc) is glass transition (Tg), or crystalline segments, where Tc=Tm. 2.1.1 Thermodynamic point of view When polymer chains are in their amorphous sate (T> Tg), they favor the strongly coiled conformation and are randomly distributed in the polymer matrix. This is the state of maximum entropy (S=KBlnΩ, where KB is Boltzman’s constant and Ω is the number of configuration) [2]. In glassy state (T< Tg), the movement of polymer chains is restricted and large scale conformation changes are not possible but localized conformation changes are allowed. The polymer behaves as an elastic solid at small strains due to primary bond stretching, which causes a change in internal energy. When a shape memory polymer is heated above Tc, the thermal activation enables the polymer chains to move and occupy one of the possible conformations without significant disentanglement. Polymers are elastic and can be deformed to large strains under this condition whether they are chemically or physically crosslinked. Polymer network exhibits “superelasticity”, where the polymer chain segments between crosslinks are able to move freely to maintain the maximum entropy. This entropic state induced by the high temperature deformation can be stored at low temperature e.g. polymer chains are frozen at T 7 The modulus, which is related to the superelasticity, is determined by configurational entropy loss that occurs with the polymer chains deformation. This is so called “entropy elasticity”. The classic rubber elasticity theory predicts that the elastic shear modulus (G) is proportional to the crosslink density and temperature: G=υKBT= (1) where υ is the number density of network chains, ρ is the mass density, R is the universal gas constant, T is the absolute temperature, and Mc is the molecular weight between crosslinks [37]. Polymers usually have elastic modulus of several MPa and they are very flexible and easy to be deformed under external force. This is different from shape memory alloys whose modulus is approximately 200MPa [38]. 2.1.2 Molecular mechanism of the shape memory effects The molecular mechanisms of shape memory effects in SMPs using their thermal transitions as a switching function are illustrated in Figure 2.1 [37]. The shape memory polymer network consists of crosslinks and molecular switches. The crosslinks can either be covalent bonds (chemical) or intermolecular interactions (physical). The switches can either be the glass transition of amorphous phase or melting of crystals. Figure 2.1 (a) shows the molecular mechanism of a physically crosslinked multiblock copolymer with segregated domains during shape memory programming. The domain that is related to the highest thermal transition temperature (Tm of high melting point crystals) acts as physical crosslinks and defines permanent network. The domain that is related to the lower thermal transition temperature (Tm of low melting point crystals) provides the switching phase. In this case, Tc=Tm. Figure 2.1 (b) displays the mechanism for a 8 covalently crosslinked SMP network during programming. In this case, the switching phase is provided by segmental crystals and Tc=Tm. Figure 2.1 (c) indicates the mechanism of a chemically crosslinked SMP network during programming. The glass transition provides switching function and Tc=Tg. For all these different networks, at temperature above Tc, the chain segments are flexible and the polymer can be elastically deformed. Below Tc, the flexibility of chain segments are limited and the temporary shape can be stabilized either by the glassy state or the crystalline domains. The permanent shape can be achieved by reheating the material above Tc. Figure 2.1 Molecular mechanism of thermally induced shape memory polymer. (a) A multiblock copolymer network with Tc=Tm, (b) A chemically crosslinked SMP network with Tc=Tm, (c) A chemically crosslinked SMP network with Tc=Tg. (Reprinted with permission from [1] ) 9 2.2 Classification of Shape Memory Polymers Based on the nature of the polymer structure, thermally-induced shape memory polymers can be divided into the following categories. 2.2.1 Physically cross-linked glassy networks Linear block copolymers are the most important group in this category due to the phase segregated property. A typical example is polyurethane (PU), in which the soft domain shows a sharp glass transition [39]. Thermoplastic polyurethane elastomers are produced in industrial scale by a per-polymer method. Segmented PU was first synthesized by researchers at Mitsubishi Heavy Inc (Japan). They prepared PU by reacting 4,4-diphenylemthane diisocyanate with polyol, where 1,4 butanediol was used as a chain extender [40]. This material consists of hard and soft segments. The hard segments, which are formed by aggregation of high polar urethane and urea blocks, are embedded in an amorphous elastic matrix. The soft segments serve as a switching phase while the hard crystalline segments provide the permanent network. At Tg In certain PU systems, the switching segments are polyethers, such as poly(tetrahydrofuran), and polyester such as poly(ethylene adipate) [41] . The glass transition temperature is usually contributed by the soft segments. Therefore, Tg can be modified by changing the chain length of soft segments block. For example, Lin and coworkers [42-43] reported using poly(tetrahydrofuran) with a molecular weight of Mn=250g/mol as switching segments to prepare a SMP. Here, the Tg was in the range 10 between 16 and 54°C, depending on the content of hard segments (57~95wt%). If the molecular weight of soft segment was Mn=650g/mol, Tg were in the range of -13 to 38°C (hard segments content 32~87wt%). When the soft segment molecular weight increased to Mn=1000g/mol, Tg of the system were around -36 to 22°C (hard segments content 23~81wt%). Tg also depended on the hard segment content because the hard segments restricted the mobility of soft segment chains. Takahashi and coworkers reported [44] a thermoplastic polyurethane which contains a soft segment based on poly(ethylene adipate) showing excellent shape memory effect. The Tg of this polymer varied from -6 to 48°C with soft segment weight average molecular weight of 300, 600 and 1000g/mol at a constant hard segment content of 75mol%. Polyesters are another class of thermoplastic SMPs with Tc=Tg. Booth and coworkers [45] reported a copolyester SMP based on poly(ε-caprolactone) and poly(butylene terephthalate), in which poly(butylene terephthalate) segments served as physical crosslinks. Another example is a poly(ketone-co-alcohol) synthesized by Perez- Foullerat using a polymer analogous reaction (reduction of a ketone and alcohol) [46]. In this work a partially reduced poly(ethylene-co-propene-cocarbonoxide) was reported exhibiting promising shape memory behavior using the glass transition of amorphous polyketone ethylene-proene/CO rich matrix as the switching temperature. The crystalline domain of ethylene/CO rich segments provided the permanent network. The Tg of this polymer can be adjusted by changing the degree of reduction, suggesting a tunable switching temperature. Similar to phase separated block copolymers, melt-miscible blends can also form physically crosslinked glassy networks to produce shape memory polymers. The 11 crystalline or rigid amorphous domains in the blends can serve as permanent network and the glassy phase with a lower thermal transition temperature serves as the switching phase. An example is the blend of thermoplastic polyurethane with phenoxy resin, in which the soft segment of the polyurethane is poly (ε-caprolactone) (PCL) [47]. The switching temperature in this system is Tg of the blend, which can be tuned by changing the composition ratio between phenoxy resin and the PCL segments. The hard segments of polyurethane serve as physical crosslinks and set the permanent shape. Mather’s group [48] studied two melt miscible blend systems: poly(vinyl acetate) (PVAc) with poly(lactic acid) (PLA) and PVAc or PMMA with poly(vinyldiene fluoride) (PVDF). In these systems, PVAc and PMMA were amorphous and PLA and PVDF showed certain degree of crystallinity. The degree of crystallinity varied corresponding to the blend ratio. The crystalline PLA and PVDF served as the permanent network while amorphous PVAc and PMMA provided the temporary network. The transition temperature is Tg, which can be tuned by changing the composition of the blend. A polymer network can also be physically crosslinked by ionic rich domains, e.g. ionic clusters. The ionic clusters arise from the strong interaction of ionic groups in ionomers and serve as physical crosslinks. The relaxation time of ionic clusters in the polymer network is long enough so that these clusters can provide a robust permanent network and give elasticity similar to chemically crosslinked materials. Xie [49] reported the tunable multi-shape memory effect of a perfluorosulphonic acid ionomer (PFSA), e.g. Nafion . This ionomer had a broad glass transition from ~55 to ~130°C and exhibits dual, triple and multiple shape memory effects, where the molecular interaction between the ionic groups set the permanent network. The shape memory polymer described in this 12 thesis can be summarized in this category, which will be deliberated in detail in the following chapters. 2.2.2 Physically cross-linked semi-crystalline networks In some block copolymers, the soft domains can crystallize and this melting transition can be used as a switching phase and Tc=Tm. An example is the thermoplastic segmented polyurethanes with semi-crystalline flexible segments such as PCL. Kim and coworkers [50] synthesized the polyester-urethanes using pre-polymer method. In this polymer, the hard segments were methylene bis(4-Phenylisocyanate) block and soft segments were poly(ε-caprolactone). The hard segment had a higher melting temperature in the range of 200 and 240°C and set the permanent network. The crystalline soft segments melted at temperature between 44 and 55°C, serving as the switching phase. The transition temperature in this system can be tuned by changing the weigh fraction of the soft segment or its molecular weight. Besides the crystalline domains of hard segments forming physically crosslinked networks, other techniques such as hydrogen bonding and ionic interactions can also form physically crosslinked network and set the permanent network. The intermolecular interactions enhance the degree of phase separation and help the polymer form a robust permanent network. This kind of system has been investigated by incorporating an ionic group into polyurethanes to further stabilize the permanent network. Buhler and coworkers [51] investigated the polyester- urethanes which was synthesized by pre-polymer method with half of the chain extending 1,4-butanediol replacing by 2,2-bis (hydroxymethyl) propionic acid in the hard segments. The propionic acid was further neutralized with tri-ethylamine. With the same PCL as the soft segments, this material showed higher elastic modulus and higher mechanical 13 strength than the uncharged hard segments because of their additional Coulomb interactions. This material showed about 10% higher strain recovery than the uncharged ones and equal strain fixation. Another example of this class is a tri-block copolymer styrene-butadiene-styrene (SBS) [52]. Polystyrene (PS) serves as the minor component at each end of the polymer chains and the major component is the semi-crystalline poly (tran-butadiene) segments in the middle block. SBS is strongly segregated due to the immiscibility between PS and PB blocks. The semi-crystalline PB blocks have a Tm of 68°C and a Tg far below room temperature. The amorphous PS blocks have the highest thermal transition in this system, Tg=90°C. Therefore, the rigid PS blocks serve as the physical crosslinks and set the permanent network. The melting of PB blocks serve as the switching phase. When the material is heated above 68°C, PB crystals melt and the material becomes flexible and rubbery. However, it does not flow due to the restriction of PS hard segments. When the material is cooled to room temperature, PB blocks crystallize and fix the temporary shape. The permanent shape can be restored by heating the material above 68°C. One advantage of SBS shape memory polymer is that the material can be reprocessed by thermal processing above 100°C. However, the disadvantage is the relatively lower shape recovery ratio (<80%) especially upon application of a large strain. This is due to the formation of oriented crystals during the deformation, which are linked through amorphous PB domains. This structure is rather stable and does not melt even at temperature as high as 80°C [53]. Linear, phase separated block copolymer made of polyethylene terephthalate (PET) and polyethylene oxide (PEO) is a further example of physically crosslinked SMP 14 with Tc=Tm [54-55]. In this system, both thermal transitions are the melting of crystals. PET has a higher melting point of 260°C and serves as physical crosslinks. PEO has a much lower melting temperature in the range of 40 to 60°C and serves as the switching phase. The transition temperature of this system can be tuned by changing the molecular weight of PEO blocks and content of PET components. Studies showed improved strain recovery ratio with the increasing content of PET. This is due to the formation of more stable physical crosslinks in the better aggregated PET blocks. Similar to the SBS system, the PEO/PET system encounters the problem of strain induced crystallization. During deformation at T>Tm,PEO, PEO segments tend to orient in the direction of deformation. The newly formed structure is stable and melts at higher temperature, resulting in a hindered recovery process. Another disadvantage of this crystalline system is the creep of the crystals under stress when setting the temporary shape, which then limits the recovery ratio. 2.2.3 Chemically cross-linked glassy networks Chemically crosslinked SMPs have the advantage of a strong permanent network that eliminates the molecular slippage between chains during programming. Therefore, the degree of shape recovery is usually excellent due to the nature of nearly permanent cross-linking. However, this kind of material is difficult to reshape once processed since the primary shape is covalently fixed. A chemically crosslinked polymer with a sharp Tg and rubbery elasticity above Tg is an ideal type in this category. Mather group synthesized [56] a chemically crosslinked, castable, transparent shape memory polymer based on vinyldiene random copolymer. Two vinylene monomer: methyl methacrylate and butyl methacrylate were used to prepare the random copolymer. The homopolymers 15 of the two components had two distinguished glass transition temperatures at 20°C and 100°C. The random copolymer showed a sharp Tg, which can be adjusted by changing the composition of two components. The thermoset was obtained by reacting this random copolymer with a tetra-ethylene glycol dimethacrylate. This material showed complete shape recovery and excellent shape fixing. Thermosetting polyurethane is another example in this category. It is easy to design PUs with different structure for desired applications such as introducing hydrophilic or ionic groups in the soft segments. Chen and coworkers synthesized [57] a series of chemically crosslinked PU with poly (ethylene glycol) (PEG) as the soft segments. This material exhibited thermally induced shape memory property with the glass transition of PEG serving as switching phase. Beside the shape memory effect, this material also showed hydrogel properties, which have potential application in the area of medical treatments and health services. Lin and coworkers reported [58] a series of crosslinked ester-type PUs with poly (butylenes adipate) glycol (PBAG) as the soft segments and trimethylol propane (TMP) as the crosslinks. The permanent network was provided by the chemically crosslinked points (TMP) and the switching phase was the glass transition of the soft segment (PBAG). This study also showed that the less TMP, or longer length of soft segment, the better shape memory performance was achieved. Since the amorphous reversible phase (PBAG) was reinforced by the introduction of TMP, the glass transition was broadened and the chain movement was delayed in the high temperature region. Thus, the recovery was influenced in the heating process. 16 2.2.4 Chemically cross-linked semi-crystalline networks The melting transition usually gives a fast and sharp recovery compared to glass transition during the shape memory process. An example of this class is chemically crosslinked semi-crystalline rubber, trans-polyisoprene (TIP) [59]. TIP has a melting point of 67°C and a degree of crystallinity of 40%. TIP can be chemically crosslinked by peroxide and form a three dimensional network. Melting transition of TIP triggers the temporary shape change. As a semi-crystalline rubber, this material has the advantage of superelasticity, fast shape recovery, and flexible modulus. Other chemically crosslinked semi-crystalline SMP example is trans-polyocyclooctene (PCO) as reported by the Mather group [60]. This material has a Tm of 58°C and Tg at 270°C. When heated above Tm, PCO showed elasticity similar to rubber and were deformed easily to large strain. This temporary shape then was fixed by the crystallization when cooled below Tm. Complete shape recovery was achieved by reheating the material above Tm. Other than synthesis approaches, chemically crosslinked semi-crystalline structure can be obtained by ion radiation, which is of great importance in the economic point of view. Polyethylene (PE) crosslinked by ionizing radiation shows excellent shape memory effect as reported by Charlesby [61]. A chemically crosslinked network of PE was formed by low doses of irradiation and this network provided the permanent shape. The melting of PE crystals served as the switching mechanism. PE chains tend to orient during the application of external force above Tm, leading to a broad melting temperature in the range of 60 to 134°C. This technique has been used in the application of heat- shrinkable products. For example, heat-shrinkable tubing is usually made of chemically crosslinked polyolefins. The material is shaped by conventional processing methods at 17 first and then covalently crosslinked by ionizing radiation. The pre-shaped parts are heated above the melting temperature of crystals, at which the material exhibits rubber elastic properties. It can be deformed to a desired temporary shape and cooled down to fix this shape, which is called “expansion” in technical terms. The commercial product is made in this expansion state and will “shrink” to its original shape upon heating. Heat shrinkable products have applications in the area of packaging industry, electronic, civil and process engineering. The most widely used materials are HDPE, LDPE and copolymers of PE and poly (vinyl acetate) [62-66]. Lendlein [67] developed a chemically crosslinked shape memory polymer using oligomeric poly(ε-caprolactone) dimethacrylate and n-butyl acrylate. This material was crosslinked under UV radiation. The crystalline PCL segments with a sharp melting point fixed the temporary network below Tm. The permanent shape was provided the chemically crosslinked network. This material was biodegradable, and had the merits of fast shape recovery and easy fabricating structure. Chemically crosslinked SMPs have the advantage of possessing a relatively strong permanent network. However, several shortcomings accompany, such as broader melting temperature and difficulty of reshaping. Chemically crosslinked structure impedes crystal formation and causes a lower degree of crystallinity. The lower and broaden melting temperature leads to incomplete shape recovery and slows down the recovery speed. Recent research has been taking efforts on keeping a sharp transition by specifically crosslinking the amorphous portion but not the crystalline parts [68]. 2.2.5 Other thermoplastic networks If the material has very high molecular weight and a sharp glass transition, it can exhibit shape memory effect. The entanglement of the high molecular weight linear 18 chains leads to the formation of physically cross-linked network, which defines the permanent shape. The most widely known materials with these characteristics are polynorbornene (PN) and high molecular weight poly (methyl methacrylate) (PMMA). The material softens above Tg and can be deformed to a temporary shape. This shape is fixed when the polymer is cooled below Tg. The decrease in mobility of the polymer chains can neither slip over each other fast enough nor disentangle. The recovery can be achieved by reheating the material above its Tg. Complete shape fixing and fast recovery were reported due to the sharp glass transition and high entanglement density. PN was discovered to not be purely amorphous. Sakurai and coworkers [69-70] reported a high molecular PN showing a tendency of stain-induced crystallization, which was evident by wide-angle X-ray scattering. According to this research, it is possible that the PN crystals contribute to set the permanent shape. Mather and coworkers [71-72] reported a POSS modified polynorbornene by copolymerization. This approach increased PN glass transition from 52°C for pure PN to 81°C, which led to an improved heat and oxidation resistance. Better shape memory properties were obtained by increasing Tg due to slowing down the relaxation of PN chains after stretching above Tg. However, this class of material has several disadvantages: (1) At high temperature, the material tends to creep because of the finite lifetime of the entanglement (2) The glass transition is not easy to tune. (3) High molecular weight materials usually have high viscosities, which will cause difficulty in processing. 19 2.3 Characterization Methods Most shape memory polymers are composed of a crosslinked (either chemically or physically) network determining the permanent shape and a reversible network fixing the temporary shape. Physically crosslinked polymers are usually phase separated and are related to different thermal transitions. These transitions can be characterization by differential scanning calorimetry (DSC) or dynamic mechanical analysis (DMA). Regarding chemically crosslinked polymers, crosslinking density is related to the material’s mechanical and thermal properties, which can be characterized by swelling experiments or DMA. For semi-crystalline systems, polarized light microscopy (POM) or wide and small X-ray scattering (WAXS, SAXS) can be used to characterized the material’s structure. As far as the shape memory performance is concerned, the most conventional method is to conduct cyclic, thermomechanical tensile tests by DMA to quantify the extent to which a temporary shape can be fixed and the permanent shape can be recovered. 2.3.1 Thermal properties characterization DSC is the most common method to assess the various thermal transition temperatures (e.g melting transition and glass transition) for polymers. Glass transition is the discontinuing of the second derivative of molar Gibbs energy and appears as a step transition on a DSC curve. Tg is usually defined as the temperature at the half height of the heat capacity change. Melting corresponds to the first-order transition where the first derivative of the molar Gibbs energy is discontinuous. Melting transition appears as a peak on DSC curve and Tm is defined as the peak temperature at the maximum melting endoderm. Figure 2.2 illustrates thermal transitions in different classes of SMPs. The 20 crystallinity (Xc) of the semi-crystalline polymer can be determined by analyzing the heat of fusion ∆H, which is the area under the melting peak. Figure 2.2 Schematic illustration of DSC curve for SMPs (a) Amorphous polymer network Ttran=Tg; (b)Semi-crystalline polymer network Ttran=Tm.Reprined with permission from [73] 2.3.2 Viscoelastic properties characterization Dynamic mechanical analysis (DMA) is a technique used to characterize the viscoelastic nature of polymers. An oscillatory displacement is applied to a sample in tension. The information including storage modulus, loss modulus, and loss factor tanδ can be obtained. This test also helps to understand the mechanical strength of the material at different temperatures and also helps in defining processing conditions. Usually the sample was tested over a range of temperatures at a constant heating rate and frequency. This experiment provides information of thermal transitions such as Tg, around which the material changes from the glassy state to the rubbery state. Figure 2.3 shows the typical viscoelastic behavior of different classes of SMPs. For semi-crystalline SMPs, the second storage modulus drop at higher temperature corresponding to melting of crystals and can be characterized as Tm, as shown in Figure 2.3 (II,III,IV). 21 Figure 2.3 viscoealstic behavior of different classes SMPs. (I) chemically crosslinked amorphous copolymer network Ttran=Tg; (II) Chemcially crosslinked semi-crystalline polymer network Ttran=Tm; (III) Physically crosslinked thermoplastic Ttran=Tg;(IV) Physically crosslinked thermoplastic Ttran=Tm. Reprinted with permission from [73] 2.3.3 Morphology characterization Most SMPs are phase separated systems and the morphology studies help us to better understand the change in properties during shape memory test. Depending on different molecular structure, the investigation of morphology can be performed using polarized light microscopy (POM), atomic force microscopy (AFM) and transmission or scanning electron microscopy (TEM, SEM). In most SMPs, the phase separation is often 22 in the micro level and the domain size is in the nanometer range, which requires the technology of SEM and TEM. 2.3.4 Shape Memory Performance Characterization The shape fixing and recovery abilities are the most important properties to assess the performance of a shape memory polymer. They reflect the microstructural transformations and determine the extent to which SMPs can be practically used. Shape fixing and recovery are quantified by cyclic, thermomechanical tests which can be accomplished by DMA. A single cycle includes deforming the sample, fixing the temporary shape and recovering to its permanent shape. The programming protocols may differ from different SMP systems, for example, different control options (stress or strain control). Stress controlled tests measure the strain change at defined thermal conditions while strain controlled tests measure the stress on the sample. Generally, the sample is tested under tension. A typical test procedure includes: (1). The material is heated above the switching temperature (Tc), at which it exists in a rubbery, elastic state. (2). The material then is deformed to a certain strain. (3). The material is cooled down below that switching temperature (T (4). The stress is then removed and the material maintains its temporary shape. (5). The material is reheated above Tc and it recovers to the original shape under stress free condition. Polymer chains regain their mobility and the stored strain 23 energy is released. The material is then cool down to end one cycle. At each step, the material is held at least 10 minutes before loading or unloading to allow the sample to equilibrate. Figure 2.4 shows a three dimensional ideal shape memory cycle obtained from a DMA test. Figure 2.4 One ideal shape memory cycle. S and E denote the start and end of the cycle. The sample is heated without applying any stress along path 1. Certain stress is applied to stretch the sample above Tc. The sample is then cooled under stress along path 3 and the stress is removed along path 4. The sample recovers its permanent shape when reheated above Tc (path 5). The sample can then be cooled to end the cycle. Reprinted with permission from [74] Shape fixing Shape fixing ratio F(N) describes to what extent the material is capable to fix the mechanical deformation, which has been applied during the programming process. F(N) can be expressed as the strain ratio of the fixed deformation to the total deformation (Eq 2) [1]. (2) Where εu (N) is the tensile strain after unloading in the Nth cycle. εm is the strain after the stretching step (before cooling or unloading the sample). Several effects may occur 24 during cooling. For example, the expansion coefficient of the material specimen changes and volume changes arising from the crystallization in the case of Tc=Tm. [75] Shape recovery If the shape memory test is performed under stress controlled mode, a strain- temperature curve can be obtained. The characteristic switching temperature can be determined as the point where the strain starts to decrease. Under strain control, a stress- temperature recovery curve can be obtained. This curve gives the information of maximum stress generated during recovery and the corresponding switching temperature. Figure 2.5 shows the recovery curves under different programming protocols. Recovery under a constant strain leads to an increase of stress. After a maximum stress, the stress starts to decrease at higher temperature, where the softening of polymer chains occurs. Figure 2.5 Schematic illustration of recovery curves of SMPs (a) recovered under constant strain conditions; (b) recovered under stress free conditions. Reprinted with permission from [73] The shape recovery ratio R determines the ability of the material to memorize its permanent shape. R is a measure of how much strain, which was applied during the programming εm, is recovered. (Eq 3). 25 )N( R pm 100%x (3) pm 1)N( Here, εp (N) and εp (N-1) represent the strains of the sample in two successively passed cycles in the stress-free state before application of stress. The recovery interval ∆Trec, which is the temperature difference between the temperature at which the recovery started and where recovery is completed, is usually reported lower than 40K [73]. 2.4 Triple shape memory effect and characterization The shape memory behavior discussed above is the conventional shape memory effect stating that the material can fix one temporary shape and recover to its permanent shape when exposed to an appropriate external stimulus. There are only two shapes involved in this effect and this phenomenon is called dual-shape memory effect. As the multiphase polymer networks were explored, the multi-shape memory effects have been discovered. The multi-shape memory polymer is able to memorize multiple temporary shapes and subsequently recover them in one shape memory cycle. Generally, the triple shape memory polymer has at least two phase separated domains and each domain is associated with individual transition temperatures, e.g. glass transition or melting temperature. Upon cooling below the transition temperature of a specific domain, this domain solidifies and forms physical crosslinks. Therefore, tuning the triple shape memory effect can be accomplished by changing the transition temperatures or the ratio between the two shape memory transition phases. Studies have been reported on triple shape memory polymer with a single yet sufficiently broad thermo-mechanical transition that can exhibit multi-shape memory effects at chosen deformation and recovery temperatures [76]. Figure 2.6 shows a 26 crosslinked shape memory polymer with two distinctive glass transitions (38°C and 75°C) exhibiting triple shape memory effect. The permanent shape A was defined by covalent crosslinks, while the temporary shape B is provided by the higher glass transition and temporary shape C is related to the lower glass transition. As illustrated in Figure 2.6, triple shape memory polymer can change shape from a permanent shape (A) to a first temporary shape (B) and from there to a second temporary shape (C) corresponding to subsequent temperature increases. This effect can also be quantified using specific cyclic, thermo-mechanical tensile tests. Similar to the dual shape memory materials test, each cycle consisted of a shape fixing and a recovery step. One cycle starts with heating the polymer network to the temperature at which the material is in a rubber-elastic state (T>Tc,B). The sample was then deformed to a temporary shape B. After equilibration, the sample was cooled to a temperature in between the two thermal transitions in the system (Tc,C 27 Figure 2.6 Illustration of triple shape memory behavior. (Reprinted with permission from [76]) Shape fixing ratios F for two temporary shapes, and subsequent recovery ratios R, can be calculated according to the following equations, ,recxy (4) F xy 100% xy xy (5) R yx 100% ,loady x where x and y denote two different shapes, εy,load is the maximum strain after applying load, εy and εx are fixed strains after unloading, and εx,rec is the strain after recovery. An example of a triple shape memory polymer is a chemically crosslinked multiphase system consisting of crystalline poly (ε-carprolactone) (PCL) segments and amorphous poly (cyclohexyl methacrylate) (PCHMA) segments with 45wt% PCL content reported by Bellin et al. [77]. In this system, the permanent shape was denoted as 28 C (Figure 2.7) and two temporary shapes were B and A, respectively. Tc,A was the melting temperature of PCL segments, which was around 50°C. Tc,B was given by the glass transition of amorphous PCHMA segments (Tc, B=Tg,PCHMA=140°C). The co- existence of the two segregated phase was confirmed by the presence of two thermal transitions. Figure 2.7 shows one triple shape memory cycle test of this system similar to a conventional dual shape memory programming process under stress-controlled mode. Two switching temperatures were determined from the strain-temperature curve where the strain started to decrease at high temperature under a stress-free condition. F and R were reported to be higher than 98%. In order to exhibit a triple shape memory effect, the material has to have very pronounced physical crosslinks from both switching domains. For example, in the PCL-PCHMA system, the content of PCL has to be between 40 and 60wt%. Figure 2.7 One shape memory cycle for the multiphase network composed of crystalline PCL segments and amorphous PCHMA segments. The solid line indicates strain, and the dashed line indicates temperature. The material has a permanent shape C and was first deformed to temporary shape B (εB) at 150°C (T> Tg, PCHMA) where the material is in the rubbery-elastic state. The samples was then cooled to Tm,PCL 2.5 Different stimuli responsive shape memory polymers 2.5.1 Indirect actuation Direct heat is the main stimulus in shape memory polymer research. However, this approach may not always be practical for technological implementation of shape memory devices. Recent research has discovered alternative triggering mechanism such as non-remote but indirect actuation. Instead of increasing the environmental temperature, the shape memory mechanism can be triggered by infrared radiation or magnetic induction heating. One example of a thermally induced shape memory polymer triggered by illumination with infrared light is a laser-activated polyurethane medical device [78-79]. This device uses a laser to heat and subsequently actuate shape memory polymer. Heat transfer in this kind of SMP can be enhanced by incorporating conductive fillers, such as carbon black, carbon nanotubes or conductive ceramics. Figure 2.8 Illustration of shape recovery under infrared radiation. Radiation exposure is from theleft. (Reprinted with permission from [5]) Figure 2.8 shows the shape recovery of thermoplastic polyurethane incorporated with carbon nanotubes [5]. The material absorbed infrared light, resulting in local heating, and was constrained to the near-surface region of a stretched sample. Thus, the shape recovery of the near surface region led to curling of the sample. The mechanical 30 properties and shape recovery ratio were influenced by the incorporation of fillers. Polyurethanes that are incorporated with carbon nanotubes exhibit a certain level of conductivity. Therefore, when an electrical current is applied, the material’s temperature is increased due to the high ohmic resistance of the composite. This subsequently triggers the shape memory effect. Indirect heating can also be achieved by incorporating magnetic nanoparticles into the shape memory polymer. The sample temperature increases by applying an alternating magnetic field [80]. Figure 2.9 shows an example of the shape recovery process of a thermoplastic composite consisting of Fe (III) oxide nanoparticles in a silica matrix and polyurethane induced by magnetic field. . Figure 2.9 Illustration of shape recovery of magnetically induced shape memory effect. (Reprinted with permission from [80]). Another application of indirect actuation is to shift shape memory polymer’s switching temperature via molecular interaction with a surrounding solvent. Huang and coworkers [81] reported a thermally activated ether-based polyurethane exhibiting shape memory effect by immersion in water. The moisture diffused to the polymer sample, disrupting the intramolecular hydrogen bonding. This diffusing moisture therefore acted as a plasticizer. The glass transition was depressed by 35°C and allowed for room 31 temperature actuation (Figure 2.10). A similar approach was adopted in polyetherurethane-polysilesquisiloxane block copolymers, where PEG was used as a polyether segment in PU block. When immersed in water, PEG segments dissolved and the Tm disappeared, leading to shape recovery. Figure 2.10 Water activated shape memory polymer recovery process. (Reprinted with permission from [81]) 2.5.2 Light activated shape memory polymer Non-heating based light activated shape memory polymer systems have been realized by incorporating reversible photo-reactive molecular switches [82-84]. Lendlein and coworkers [3] developed a light induced shape memory polymer based on photo- reversible cinnamate. When the material was exposed to suitable wavenumber light, cyclo-addition reaction occurs between two of the light-sensitive moieties, resulting in a crosslink of the material. Irradiation to a different wavenumber light caused the cleavage of the newly formed bond and induced shape recovery. In this system, the permanent 32 shape was determined by the chemically crosslinked network. The temporary shape was fixed by exposing the stretched sample under UV light with a wavenumber larger than 260nm to create new covalent crosslinks. The shape recovery was achieved by exposing the sample to UV light smaller than 260nm to destroy the newly formed crosslinks (Figure 2.11). Lee and coworkers [85] studied a light induced shape memory system based on photoisomerization of azofunctionalities in a glassy liquid crystalline network. In this system, the temporary shape was fixed by exposure to linearly polarized 442nm light and the permanent shape was recovered by exposure to circularly polarized light with the same wavenumber. Figure 2.11 Molecular mechanism illustration for light induced shape memory polymer. Photoreversible crosslinks (filled diamonds) are formed during fixation of the temporary shape by UV light. (Reprinted with permission from [3]) 2.6 Biomedical Applications Besides traditional application such as heat shrinkable packing and tubing, the application of shape memory polymers in the area of biomedical treatment has been 33 explored. Recently, shape memory polymer research focuses primarily on the medical applications due to a combination of biocompatibility with their wide range of tunable stiffness, tailorable transition temperatures, fast actuation, large shape deformation, complete recover, and elastic properties of the materials [86]. In addition, shape memory polymers in the form of porous foams further broadens the range of potential volume- filling applications. Most importantly, the relatively inexpensive fabrication of shape memory polymer devices into complex shapes opens the possibility of customizing patient-specific device geometries. Lendlein and coworkers developed a self-tightenable biodegradable suture for wound closure from a biodegradable shape memory thermoplastic monofilament fiber [15]. The suture can be loosely connected and then heated above the critical temperature to trigger the shape recovery and tighten the suture. Figure 2.12 illustrates the process of wound closure by the intelligent shape memory suture. The extruded monofilament of oligo(ε-capropactone)diol-based shape memory polymer was stretched under heat and cooled to fix the temporary shape. When used, the suture responded to body heat (37°C) and started to recover to the permanent shape. It generated enough force during recovery, therefore tightening the wound by itself. Since the suture was biodegradable, removal of the suture was not necessary. Figure 2.12 Illustration of degradable shape memory suture for wound closure. (Reprinted with permission from [15]) 34 Another example is a laser-activated shape memory device for mechanical removal of blood vessel clot to treat ischemic stroke [78-79]. Maitland and coworkers demonstrated a photothermal actuation shape memory device manufactured into corkscrew-shape or umbrella-shape by coupling light from a diode laser operation at a wavenumber of 810nm. The device can be inserted by minimally invasive surgery into the blood vessel and penetrated the clot in a narrow form. The shape recovery occurred upon laser activation and the device turned into a clot-grabbing form, enabling the mechanical removal of the blood clot. Hybrid devices, consisting of a shape memory nickel-titanium alloy wire core encapsulated in a shape memory polymer shell [87], have also been developed in order to enhance the recovery force compared to a pure shape memory polymer device. Vascular stents are another example of shape memory polymer in biomedical applications. Vascular stents are small tubular scaffolds used to maintain the patency of an artery. [88-90]. Conventional stents are mostly made of stainless steel, shape memory alloy or other metal alloys. They are suffering several drawbacks such as being too stiff to navigate highly tortuous vessels, compliance mismatch due to the stiffness at the arterial wall, and high fabrication cost [91]. SMPs have the advantage of large flexibility, compliance, and low fabrication cost. Gall and coworkers described a thermally responsive shape memory stent based on acrylates [92-93]. This device was fabricated by injection molding to a temporary shape (Figure 2.13 t=0). The stent recovered to the permanent “expanded” shape upon heating to 37°C (body temperature), which occurred in only 100 seconds. The triple shape memory effects concept could be employed to fabrication of this device. The stents could be designed to implant with their first 35 temporary shape. Upon heating, they expand to their functional shape (secondary temporary shape). Upon further heating, the stents shrink to their permanent shape, which will be beneficial for easy removal. Figure 2.13 Shape recovery of a transparent shape memory polymer stent to the permanent unfolded shape. Black lines facilitate deployment visualization. Reprinted with permission from [92] 2.7 Innovation of this research As discussed above, the development of thermally induced shape memory polymers has focused primarily on relatively low transition temperature (Ttrans<100°C ) and targeting applications such as biomedical devices, smart fabrics, sensors, and actuators. Materials used as aerospace components often require higher switching temperatures for shape change and actuation. To the best of our knowledge, there have been no reports of thermoplastic SMPs with controllable switching temperatures above 100°C. There has been research on high temperature SMPs based on thermosetting polymers systems. For example, Mather and Liu [33] described a thermosetting shape memory polymer prepared by copolymerizing two monomers, each separately producing polymers characterized by different glass transition temperatures. In the presence of a 36 difunctional monomer, whereby the copolymer formed is cross-linked during the polymerization to produce a thermosetting network, the transition temperature of the final shape memory polymer can be tuned by the ratio of the monomers from 20°C to 110°C . This material can be processed as castable formulations in the form of coatings and films. It is optically transparent and has potential use as medical plastics. Tong [34] developed a similar thermosetting copolymer with a covalent crosslinked network. The first monomer is from the category of maleimide-based monomers and the second monomer is another vinyl monomer. The transition temperature of the final copolymer can be adjusted by the ratio of the monomers selected so that the resulting polymer has a Tg from about 150°C to 270°C . This material was also castable but the application is to make molds for casting composite parts. Shape memory polymer can be used to make molds and therefore create a gentle, automated and simple demolding process, which is due to the polymer’s ability to memorize the temporary shape and recover back to the original shape by simply heating over the transition temperature. For example, a liquid resin is injected into a mold thermoformed using a shape memory polymer sheet. After the resin cures, the mold is heated above the polymer Tg, which allows the mold to retract back to the sheet. The shape memory polymer discussed in this patent has a relatively higher glass transition temperature, which would allow resin to cure at high temperatures. Gao et al. [35] reported a thermosetting shape memory polymer made of epoxy resin modified by poly (ether ether ketone) (6F-PEEK). This blended material formed a phase-segregated network and showed two glass transition temperatures. The cured epoxy domain showed a higher Tg of 223°C and served as hard-segments forming phase to set the permanent shape while 6F-PEEK showed a lower glass transition and was responsible for the 37 switching phase. The transition temperature in this system was 150°C . Complete and fast shape recovery was reported. Koerner and coworkers [94] reported a shape memory polymer composite based on chemically crosslinked aromatic CP2 polyimide and associated single wall carbon nanotubes (SWNTs). The switching temperature of this system is 220°C and the material shows excellent fixity (>98%), good cyclability, and outstanding creep resistance. The above discussed shape memory polymer systems show relatively high transition temperature. However, all of these materials are thermosetting networks and have the disadvantage of difficulty in reshaping. The chemical crosslink net points restrict the polymer chain from moving freely, so these systems tend to have relatively broad glass transition which lead to a slow and imperfect shape recovery. The objective of this research is to develop a thermoplastic, high switching temperature shape memory polymer based on sulfonated poly(ether ether ketone) (SPEEK) with the transition temperature as high as 250°C that may be appropriate for aerospace applications. The shape memory polymers described in this dissertation are based on polyether ether ketone (PEEK), which is a semi-crystalline thermoplastic polymer with excellent mechanical properties. It is highly resistant to thermal degradation and has a maximum service temperature of 260C. PEEK has been developed as a shape memory device in biomedical research, but the shape memory effect was triggered by mechanical force. The product Morphix suture anchor invented by Medshape Inc., is made of PEEK [95]. The device is pre-compressed in a low-profile geometry. After implantation, it responds to cyclic loading to attain its permanent shape. Our objective is to use thermal transitions of the material to develop a high switching temperature shape memory polymer. 38 Neat PEEK has some thermally actuated shape memory behavior with a crystalline phase providing a permanent network and the glassy phase providing the temporary network. However, it exhibits rather poor shape memory behavior which is probably a consequence of creep of the “permanent” crystalline network. PEEK can be modified simply through sulfonation, which is an electrophilic substitution reaction. The degree of sulfonation can be controlled by varying the reaction time and temperature. SPEEK ionomer with high degree of sulfonation has been reported as a low-cost alternative membrane for hydrogen powered polymer electrolyte membrane fuel cells (PEMFC) and for methanol powered direct methanol fuel cells (DMFC) [96-98]. These membranes possess high proton conductivity and good thermal stability, but their mechanical properties are deteriorated by the high degree of sulfonation. In this work, we are interested in relatively lower sulfonation level where there is enough intermolecular interaction between sulfonic groups without compromising the mechanical properties. Neutralization of SPEEK to metal salts, such as Na+, Zn2+, Ba2+, Al3+, and Zr4+ provided a robust physical crosslinked network due to dipolar associations and microphase separation of ion rich domains. This physically crosslinked structure significantly changed PEEK chain’s packing and increased its molecular bulkiness. M-SPEEKs with reasonable high Tg that related to the Coulomb energy of ion pairs were obtained. Dual shape memory behavior can be achieved by using Tg as switching temperatures and the physically crosslinked network of M-SPEEK provided the permanent structure. A new approach to prepare a shape memory polymer by blending the polymer matrix with a low molar mass crystalline material was pursued in this research. The compounds of M- SPEEK and 30wt% NaOl exhibit excellent dual shape memory behavior, where the 39 strong dipolar interactions between the ionomer and a dispersed phase of crystalline NaOl provided the temporary network. The M-SPEEK/NaOl compounds also exhibited Triple shape memory using the two separate switching temperatures of the two temporary networks, Tg from the ionomer and Tm from the NaOl. The innovation behind this research is to develop a high temperature shape memory polymer from PEEK taking the advantage of the processing possibility due to its thermoplastic nature. PEEK can be melt processed above its melting temperature (340C). The ionomers are amorphous materials and in principle they should be processable above Tg by conventional polymer processing operations. The processability will be improved for the M-SPEEK/NaOl compound, which should be processable above Tm of the NaOl. The role of NaOl then would be a plasticizer. Chapter III and IV discuss details about preparation and characterization of the high temperature shape memory polymers based on M-SPEEK ionomers. 40 CHAPTER III SYNTHESIS, CHARACTERIZATION, AND SHAPE MEMORY BEHAVIOR OF M-SPEEK IONOMERS 3.1 Introduction This chapter presents the synthesis, characterization and shape memory behavior studies of metal ion neutralized sulfonated polyether ether ketone (M-SPEEK) ionomers. The starting material used in this research is polyether ether ketone (PEEK), which is a semi-crystalline thermoplastic polymer with excellent mechanical properties and thermal stability. It is highly resistant to thermal degradation and has a maximum service temperature of 260C [99]. Neat PEEK exhibits shape memory behavior with a crystalline phase providing a permanent network and the glass transition providing the reversible thermally actuated temporary network. However, PEEK exhibits rather poor shape memory behavior which is probably a consequence of creep of the “permanent” crystalline network. PEEK can be modified simply through sulfonation, which is an electrophilic substitution reaction. The degree of sulfonation can be controlled by varying the reaction time and temperature. SPEEK ionomer with high degree of sulfonation has been reported as a low-cost alternative membrane for hydrogen powered polymer electrolyte membrane fuel cells (PEMFC) and for methanol powered direct methanol fuel cells (DMFC) [100-101]. 41 Those membranes usually are sulfonated to ion exchange capacity (IEC) of greater than 1 meq/g (i.e., > 30 mol%) to achieve high proton conductivity. High sulfonation levels, however lead to high water absorption and poor mechanical properties. In the present work, lower sulfonate concentrations were used to prepare high temperature shape memory polymers to achieve sufficient intermolecular interactions and sustain good mechanical properties. Neutralization of SPEEK to metal salts (M-SPEEK), such as Na+, provided a physical crosslink due to dipolar associations and microphase separation of an ion-rich nanophase. These ionic clusters significantly changed PEEK polymer chain’s packing and increased molecular bulkiness. The limited chain mobility results in an increasing of Tg, which can be used as a switching temperature. Because of the long relaxation times of the strong ionic associations, the ionic nanodomains functioned as the “permanent network”. This structure is more robust than the PEEK crystallites and an improved shape memory behavior can be achieved. + The alkali metal cation, Na , can only associate with one –SO3H group, while one 2+ divalent cation such as Zn , can combine with two –SO3H groups to attain electron neutrality. When many –SO3H groups from different SPEEK chains are linked together by these divalent cations, a network of SPEEK chains can be formed. In this network, there are two possible crosslinking mechanisms that can take place. One is the ionic corsslinking, which occurs as a result of achieving electrical neutrality in the material. The other one is the physical crosslinking due to ion-dipole associations of the + + SO3¯ Zn SO3¯ groups, which produces ionic aggregation, e.g. ionic clusters. These ionic clusters provide multifunctional crosslinks produced by nano-phase separation of ion-rich domains of the order of 1~5 nm in size [102], which can be detected by small angle X-ray 42 scattering. When multivalent ions are used to neutralize an ionomer, the ionic crosslinking will take place due to the need for electronic neutrality in the material. Ionic crosslinking is usually achieved by exchange some of the protons in acidic form of the polymer with multivalent cations. The salts produced by the exchange reaction should have very low water solubility and form a stable crosslink network. Generally, this phenomenon occurs in solutions of anionic polyelectrolytes and cations due to the strong electrostatic forces surrounding the polymer chains. The cations are attracted to and constrained in the vicinity of polyanions. The most common example of ionic crosslinking is calcium-crosslinked alginate for in situ gelling applications [103- 105]. The sodium alginate, which is a sodium neutralized carboxylated polymer, is immersed in an aqueous solution of calcium chloride. Calcium electrostaically interacts with the anionic polycarboxylates and crosslinked the alginate gel rapidly. Ionic crosslinking has been explored in various applications such as gas separation, pervaporation, and facilitated transport membrane [106-109]. The strength and stability of the ionic bonds between the neutralizing cation and the polymer’s fixed charge group is relatively stronger than other noncovalent chemistries. Alumium and zircomium are trivalent and tetravalent cations, respectively. They can form salts that have very low solubility in water and are chosen to be the crosslinking species. Barium ion is a larger size divalent cation and is known to form salts that have extremely low solubility in water. The properties and shape memory behavior of barium neutralized SPEEK were also studied in this work to compare with other M-SPEEKs. Therefore, SPEEK has been neutralized to different cations with various charges: Na+, 43 Zn2+, Ba2+, Al3+, Zr4+. The strength of ionic interactions between the cations and sulfonate groups often scale with the Coulomb energy, Ec, of the ion-pair: (6) -19 + where Ke is Coulomb’s constant that equals 2.31 x 10 J ·nm, q is the charge of the mobile ion, q- is the charge of the fixed ion and a is the separation of the charges, which is essentially the ionic radius [110-111]. For the sulfonate ionomers, q- =1, so that + the Ec ~ q /a, written hereafter as q/a. Table 3.1 shows the characteristics of the metal ions, including the ionic radius, q/a values, and electronegativity. Aluminum-sulfonate bond is the strongest among other sulfonate bonds to metal ions, because aluminum has the highest q/a value, which indicates the highest coulomb energy. The stability of the metal sulfonate groups can be understood from the electronegativity values of the cations [112]. The oxygen atoms in the –SO3H group are highly electronegative, therefore, the stability of metal-sulfonate group increases as the electronegativity of the metal increases. More stable of metal-sulfonate group indicates more covalent nature of the bonds. Aluminum-sulfonate and Zinc-sulfonate bonds are more covalent, while sodium-sulfonate, barium-sulfonate and zirconium-sulfonate bonds are more ionic. As shown in Figure 3.1 Aluminum can associate with three –SO3H groups and zirconium can combine with four –SO3H groups. These –SO3H groups may come from different SPEEK chains and they are connected by the Al3+, Zr4+. These two multivalent metal ion act as crosslinking point and a network can be formed in the material. Generally, crosslinking affects several important properties of the material such as increase in transition temperatures, reduction in the sorption in swelling solvents, and increase in plateau modulus [113-121]. Crosslinking can also enhance the tensile strength 44 and toughness of the material in some cases. In the present work, the structure, thermal properties, and linear viscoelastic behavior of M-SPEEKs were studied. Dual shape memory behaviors of M-SPEEK with different switching temperature were observed and shape memory performance was assessed. Table 3.1 q/a value for different cations Cation a(Å-1) q/a(Å-1) Electronegativity Na+ 1.02 0.98 0.93 Ba2+ 1.35 1.48 0.89 Zn2+ 0.74 2.70 1.65 Al3+ 0.535 5.61 1.61 Zr4+ 0.72 5.55 1.33 Figure 3.1 Illustration of ionic interactions between metal ions and sulfonate group 45 3.2 Experimental Details 3.2.1 Materials PEEK powder with a Mw = 96,000 g/mol was purchased from Victrex (Grade 450PF). The PEEK powder was dried in a vacuum oven at 120C for 24 h prior to use. Concentrated sulfuric acid (96.3%), N-methyl-2-pyrolidone (NMP, 99%) was obtained from Fisher Scientific and used as received. Zinc acetate dihydrate (>98%), Barium acetate (99%), Aluminum acetylacetonate (98%), zirconium acetylacetone (99%) were obtained from Sigma Aldrich Chemical Co. and used as received. Sodium acetate (99%) was purchased from Baker & Adamson Chemical Co. and used as received. 3.2.2 Sulfonation Sulfonated PEEK (SPEEK) was synthesized by dissolving 10 g PEEK powder into 250 mL concentrated sulfuric acid (96.3%) to form 40g/L solutions, and vigorously stirred at either room or elevated temperatures for desired reaction time to achieve different sulfonation level [122]. The resulting SPEEK was precipitated by drop-wise addition of the solution into rapidly stirred de-ionized water at 0C. The product was filtered and washed repeatedly with de-ionized water until the pH was about 7. The polymer was then ground in a blender and washed thoroughly to completely remove acid. The final product was dried in a vacuum oven at 180C for 24 hours. The reaction is illustrated in Figure 3.2. For PEEK, sulfonation with concentrated sulfuric acid only places sulfonic acid groups on the phenyl ring attached to two ether oxygen atoms. The sulfonation level was determined by the titration of the sulfonic acid groups. 46 Figure 3.2 Schematic illustration of sulfonation and neutralization reaction 3.2.3 Neutralization The SPEEK samples prepared for this research have relatively lower sulfonation level and they are swollen by water but do not dissolve in water. The neutralization of sulfonic acid derivative of SPEEK to metal ion salts (M-SPEEK) was conducted under heterogonous conditions. The agents used to neutralize H-SPEEK to Na-SPEEK, Zn- SPEEK, and Ba-SPEEK are sodium acetate, zinc acetate, and barium acetate which are soluble in water. A slurry of SPEEK in distilled water containing 2-fold excess of one of the metal acetates was stirred for 24 hours at 100°C, then washed with water and methanol to remove the excess metal acetate. The M-SPEEK ionomers were then filtered 47 and dried under vacuum at 120°C for 24 hours. The temperature was then increased to 180°C for 2 hours to remove most of the remaining solvent. For Al-SPEEK and Zr- SPEEK, the neutralization agents were Aluminum acetylacetonate (Al(AcAc)3) and zirconium acetylacetone (Zr(AcAc)4). Both of these substances not soluble in water but they can dissolve in warm ethanol. The reaction was conducted by mixing a slurry of SPEEK in distilled water with either Al(AcAc)3 or Zr(AcAc)4 in ethanol and stirred for 24 hours at 100°C, then washed with water four times, three times with ethanol. The products were then filtered and dried under vacuum at 180°C for 24 hours. The metal salt products were titrated again following the same procedures as for H-SPEEK. No H+ was detected. Film samples of M-SPEEK were prepared by dissolving the M-SPEEK in refluxing N-methyl-2-pyrrolidone (NMP). The solution was cast onto a clean glass plate and then dried at 80C for 48 h to remove most of solvent. The film samples were dried at 150°C under vacuum for 24 h prior to characterization. PEEK film was also prepared for comparison by compression molding. The PEEK powder was molded into 0.3 mm films at 370°C . 3.2.4 Water uptake measurements Water uptake of H-SPEEK and M-SPEEK were measured by the difference in weight between the dry and swollen films. The films were dried in a vacuum oven at 180°C for 24 h. Films with uniform rectangle shape of all samples were weighted and immersed in deionized water at room temperature for 48 h. The films were swollen with water until no further weight gain was observed. The films were then removed from water and dried with tissue paper before weighing. The weight was recorded and the 48 amount of water gained by the films was calculated. Each M-SPEEK sample was tested with three specimens and the average results were reported. 3.3 Materials characterization 3.3.1 Titration The sulfonation level was determined by the titration of the sulfonic acid groups. SPEEK was first ion-exchanged with excess saturated aqueous sodium chloride solution overnight, which was found to be sufficient to convert SPEEK-H into SPEEK-Na. This ion-exchange procedure was similar as the neutralization method and was conducted under heterogeneous condition. The resultant HCl solution was titrated with a normalized (0.01N) sodium hydroxide (NaOH) solution using phenolphthalein as an indicator [123]. The sulfonation level has been expressed in different ways in common ionomer literature [124-125]: 1. Ion Exchange Capacity (IEC): defined as the amount of acidic ion exchangeable sites per weight, which indicates a concentration of sulfonate groups in the ionomer. 2. Equivalent weight (EW in g/mol): defined as the weight of polymer per mole of sulfonate group. IEC is the inverse of the equivalent weight IEC=1/EW. 3. Degree of sulfonation, DS, defined as the molar ratio of sulfonated units to the total repeat units and calculated from the ion exchange capacity measurements. The molar quantity of the sulfonic acid groups (-SO3H) in the H-SPEEK can be determine by measuring the volume of the normalized NaOH solution consumed in the titration. The IEC can be estimated by Equation (7). 49 (7) Where VNaOH and CNaOH, are volume and concentration of NaOH solution, respectively. MH-SPEEK is the weight of H-SPEEK power. Since the degree of sulfonation is the average number of sulfonate group per repeat unit and DS can be calculated from IEC as : IEC)(288 DS 1000 80 IEC)( (8) Where 288 is the molecular weight of PEEK repeat unit, 80 is the molecular weight of –SO3¯. 3.3.2 Structure Characterization Fourier transform infrared (FTIR) spectrometer in transmission mode was used to confirm the sulfonation of PEEK to H-SPEEK, and neutralization of H-SPEEK to Na- SPEEK, Ba-SPEEK, Zn-SPEEK, Al-SPEEK and Zr-SPEEK. FTIR spectra were obtained on cast films (thickness 2 µm) with a NicoletTm 380 FTIR using 32 scans with a resolution of 2 cm-1. Spectra covered a wavenumber range of 400-4000 cm-1. Small angle X-ray Scattering (SAXS) was used to characterize the nanophase separation of M-SPEEK ionomers arising from the ionic group aggregations. SAXS measurements were performed on solution cast thin films at room temperature by a Rigaku rotating anode (40KV accelerating voltage; 100 mA current). Thin film samples were annealed at their glass transition temperatures under vacuum for 48h. 3.3.3 Thermal Analysis The glass transition temperatures (Tg) and melting temperatures (Tm) were measured with a TA instruments differential scanning calorimeter (DSC) Q200 using a 50 nitrogen atmosphere, and heating and cooling rates of 10C/min. Tg was defined as the temperature at the half height of the heat capacity change. Tm was defined at the maximum rate of melting, i.e., the peak temperature of the melting endothem. The thermal stability of the materials was measured using a TA Instruments Q50 thermogravimetric analyzer (TGA) from room temperature to 800C with a heating rate of 10C/min. In order to determine whether the material gradually degrade at high temperature, the sample was held at a constant high temperature for 60 mins and then heat to 800C with a heating rate of 10C/min. 3.3.4 Mechanical Properties and shape memory cycle The viscoelastic properties of films were measured with a TA Instruments dynamic mechanical analyzer (DMA) Q800. Strain sweep was performed to determine the linear viscoelastic region. The film sample was then deformed in the tensile mode with an oscillatory strain (0.2%), a frequency of 1 Hz, and a temperature ramp from room temperature to 300C of 3C /min. The shape fixation and recovery efficiencies were measured from shape memory cycles carried out using a tensile film fixture and the controlled force mode of the DMA. For Na-SPEEK, Ba-SPEEK and Zn-SPEEK, the film sample was first heated from room temperature with a preload force of 0.005 N to 270C. which is higher than their glass transition temperatures. For Al-SPEEK and Zr-SPEEK, the film sample was heated to 300C due to their higher glass transition temperatures. After the sample equilibrated at 270C, the force was increased to 0.5 N and the sample was stretched isothermally. Once the strain equilibrated, the sample was cooled rapidly, under load, to 30C to fix the 51 temporary shape. The force was then lowered to 0.005 N, which was sufficient to prevent sagging of the sample when it was reheated above Tg. The shape recovery of the sample was achieved by reheating it to 270C with a constant force of 0.005 N and holding it isothermally at 270C for 20 min to allow the strain to equilibrate. 3.3.5 Rheology Dynamic shear measurements were performed with a TA Instruments advanced rheometric expansion system G2. A 8mm parallel plate fixture was used. Strain sweep experiments were conducted to determine the linear viscoelastic region and all frequency sweep tests were made within the linear viscoelastic limit. Frequency sweep tests at temperatures from 260 to 300 °C were performed covering a frequency range of 0.01- 100Hz. As the sample temperature was changed, the gap between two parallel plates was adjusted to account for thermal expansion of the sample. 3.4 Results and Discussion 3.4.1 Sulfonation Sulfonation is a powerful polymer modification method to improve materials properties such as better wettability, high water flux, and increased solubility in solvent for processing [126-131]. In particular, sulfonation has been used to modify aromatic thermoplastics to improve the hydrophilicity of the material in the development of alternative more economical non-perfluorinated polymer proton exchange membranes for high temperature use [132-135]. Many aromatic polymers have been sulfonated such as polystyrene, polysulfone, poly (ether ketone ketone) and poly(2,6-dimethyl- 1,4phenyleneoxide) [136-138]. Sulfonated materials have been used in various 52 applications, for example ion exchange resin, reverse osmosis membranes, proton exchange membranes, and biomedical devices [139]. Sulfonation can be achieved by either directly introducing the –SO3H group onto the polymer chain or polymerizing sulfonated monomers [140-142]. The sulfonation reaction rate was found to be first order with respect to the PEEK chain repeat unit [143]. Bailly and coworkers reported the sulfonation of PEEK in concentrated sulfuric acid and in the mixture of methanesulfonic acid with concentrated sulfuric acid. The sulfonation level was a function of the fourth power of the sulfuric acid concentrations [144]. Sulfonation of PEEK in concentrated sulfuric acid is an electrophilic substitution reaction. The sulfonate groups are introduced to the phenyl ring attached to two ether oxygen atoms, where the polymer chain activated for electrophilic substitution by the ether linkage. This was confirmed by 13C-NMR study [145]. Once the sulfonate group is attached to the ring, the electron-withdrawing effect deactivates the reaction. The other two phenyl rings connected through the carbonyl linkage are deactivated for electrophilic sulfonation due to the electron-withdrawing effect of the carbonyl group. Therefore, only one sulfonate group can be introduced to one repeat unit and the degree of sulfonation can not exceed 100 mol%, which is IEC<2.56meg/g. The degree of sulfonation can be controlled by sulfonation reaction time and temperature. Other variables may affect the reaction are the acid strength and polymer concentration, which are not discussed here. A Series of SPEEK with different sulfonation level was obtained by varying the reaction time and temperature, as shown in Table 3.2. PEEK film prepared by compression molding was listed for comparison. The presence of polar –SO3H groups 53 makes the SPEEK polymer become hydrophilic. The introduction of –SO3H groups also changes the polymer chain packing, leading to a reducing of crystallinity. In this research, the starting material, PEEK is a thermoplastic crystalline material with Tm at 338°C. 2.9 mol% SPEEK exhibited a reduced Tm at 333°C. 7.6 mol% SPEEK showed a further reduced Tm at 289°C and significantly decreased crystallinity (low ∆H value, Table 3.2). Above 11mol%, SPEEK is totally amorphous. Tg of SPEEK increases with increasing sulfonation level, see Table 3.2. The introduction of –SO3H groups into PEEK polymer chains also increased Tg by as much as 16 °C for only 2.9mol% sulfonation. This is due to increased intermolecular association through hydrogen bonding and increased polymer chain bulkiness. Table 3.2 Thermal Properties of SPEEK with different reaction conditions ΔCp ΔH ΔH Reaction IEC mol% Tg (°C) Tm(°C) Tc (°C) (J/g) (J/g) (J/g) condition PEEK 0 148 0.10 338 42.4 297 48.37 0.1 2.9 164 0.1 333 34.67 25°C 3hr 0.258 7.6 172 0.33 289 5.097 25°C 5hr 0.374 11 178 0.22 25°C 6hr 0.60 18 181 0.27 40°C 5hr 0.91 28 201 0.19 45°C 2hr 1.07 34 210 0.20 45°C 4hr As a consequence of increasing hydrophilic nature and reducing of crystallinity, the solubility of SPEEK is affected related to the sulfonation level. At relatively high sulfonation level such as above 40%, the polymer is soluble in dimethulformamide (DMF), N,N-dimethylacetamide (DMAc), dimethylsulfoxide (DMSO) at room temperature. SPEEKs with even higher sulfonation level (> 70%) can dissolve in 54 methanol [121]. The polymer becomes water soluble when the sulfonation level reaches 100% [121]. However, at lower sulfonation level such as less than 30%, the polymer is soluble in reflux N-Methyl-2-pyrrolidone (NMP). SPEEKs with very low sulfonation level (<11%) can only dissolve in very strong acid (concentrated H2SO4). SPEEK with sulfonation level less than 7.6mol% is crystalline material and is not soluble even in refluxing NMP. 11 mol% SPEEK exhibited limited solubility in refluxing NMP. In this research, SPEEK with sulfonation level of 18mol% was selected for the further neutralization and shape memory behavior studies, because of the processability. 3.4.2 Neutralization The 18mol% SPEEK was neutralized to different metal salts to achieve more robust physical crosslinking network provided by the ionic domains. Metal ions associate with sulfonate groups on the polymer chain back bone acting as multifunctional crosslinking point, which scaled with their coulomb energy. The properties of M-SPEEK are differentiated from each other due to their unique chemical structures. 3.4.2.1 Structure characterization of M-SPEEK FTIR was employed to detect the local structure changes as the sulfonate groups bonded to metal ions. Figure 3.3 shows FTIR spectra of M-SPEEK with wavenumber from 600-1400cm-1. The absorption around 1652cm-1 is assigned to the carbonyl band. Two absorptions at 1597 and 1491cm-1 are the characteristic peak of aromatic C-C band. The absorption at 1215cm-1 is attributed to the ether bridge in the SPEEK repeat unit. PEEK, H-SPEEK, and M-SPEEK show a band at ~860cm-1, which is the characteristic peak for out of plane stretching of the two hydrogen on the benzene ring (1,4 disubsitution). For PEEK, this band exhibited at 863 cm-1. For H-SPEEK and M-SPEEK, 55 this peak slightly shifted to lower wavenumber number e.g. 858cm-1. The band at ~840 cm-1 for PEEK, H-SPEEK, M-SPEEK is the characteristic peak of the out of plane C-H stretching on the benzene rings. The characteristic absorption bands at 1248, 1078, 1025, and 765 cm-1 are the evidence of successful sulfonation of PEEK [146], which are assigned as various sulfur oxygen vibrations: asymmetric O=S=O stretching, symmetric O=S=O stretching, S=O stretching and S-O stretching, respectively [147]. PEEK also shows absorption bands at 1250 and 769cm-1, which is due to the solvent, diphenyl sulfone, in the production of PEEK. Diphenyl sulfone has a boiling point of 379°C and is extremely hard to remove from the PEEK product. This residue solvent has a O=S=O group and exhibits IR absorption at the similar wavenumber as sulfonated groups in H-SPEEK. However, this solvent can dissolve in concentrated sulfuric acid and can be eliminated after the sulfonation reaction. Therefore, the peaks at 1248 and 765 cm-1 are the sulfonated groups in H-SPEEK. When the proton in H-SPEEK was exchanged with other metal ions, the electronegativity between O-H has been changed to O-M. The intermolecular interactions have been changed from hydrogen bonding for H-SPEEK to ionic interactions for M- SPEEK. These interactions influence the potential between the atoms and displace the stretching bands of O=S=O, S=O, and S-O [148]. The wavenumber of these absorption bands of M-SPEEK and H-SPEEK are summarized in Table 3.3. M-SPEEK displays completely different spectrum from H-SPEEK due to the rearrangement of electrons when the proton was exchanged to a metal ion. The –SO3¯ group has a pyramidal structure e.g. C3V symmetry and the metal ion disturbs this symmetry because of the 56 cation-anion interactions. For sulfonated polystyrene salts, for example Na-SPS, the -1 asymmetric stretching vibration of –SO3¯ ion exhibits doublet at around 1200cm , which is due to the removal of C3V symmetry by the electrostatic field of the cation [148]. In the case of M-SPEEK, the ethyl linkage between benzene rings has a strong absorption at -1 -1 1214cm . The doublet at 1200cm of –SO3¯ ion asymmetric stretching vibration was hard to resolve. Considering the structure difference of SPEEK and SPS, one could assume that the doublet of –SO3¯ ion asymmetric stretching vibration should appear at different wavenumber. The bands at around 1150 cm-1 exhibited splitting for Al-SPEEK and Ba-SPEEK, which is possibly due to the –SO3¯ ion asymmetric stretching vibration. This splitting became greater as the increasing of ion pair’s Coulomb energy e.g., increased field of the cation at the anion. The symmetric stretching of S=O, and S-O bands shifted slightly to lower wavenumber for Na-SPEEK, Zn-SPEEK and Ba-SPEEK, but shifted to much lower wavenumber for Al-SPEEK and Zr-SPEEK. This is due to the increased cation field of Al3+ and Zr4+. Table 3.3 S-O stretching bands of -SO3¯ groups in M-SPEEK Stretching vibration of the s s Electro q/a -1 S=O S-O Sample -SO3¯ ion (cm ) negativity (Å-1) (cm-1) (cm-1) Symmetric Asymmetric Asymmetric H-SPEEK 1078 - 1248 1025 765 Na-SPEEK 0.93 0.98 1078 1148 1244 1022 757 Ba-SPEEK 0.89 1.34 1071 1147 1243 1018 761 Zn-SPEEK 1.65 2.70 1078 1155 1242 1015 763 Al-SPEEK 1.61 5.61 1069 1153 1241 1000 755 Zr-SPEEK 1.33 4.65 1068 1150 1242 1004 759 s symmetric stretching 57 1248 1145 1078 1025 765 Zr-SPEEK Al-SPEEK Ba-SPEEK Zn-SPEEK Na-SPEEK Absorbance (a.u.) H-SPEEK PEEK 1400 1200 1000 800 -1 Wavenumber (cm ) Figure 3.3 FTIR spectrum of PEEK and M-SPEEKs in the spectral region 1400 – 600 -1 cm , which is characteristic for the sulfonate absorbances. 3.4.2.2 Water uptake High sulfonation level SPEEK has been widely used to develop proton exchange membranes (PEMs). It is known that water plays an important role in PEMs. The water molecules dissociate sulfonic acid group’s proton and facilitate it transport, therefore, the conductivity depends on the available acid groups and their dissociation ability in water. 58 However, excessive water uptake can lead to material fragility, failure in mechanical properties, and poor stability at elevated temperatures. In this research, the equilibrium water absorption for 18 mol% H-SPEEK hydrated at room temperature by immersing into deionized water is 13wt%. The absorbed water molecules form hydrogen bond with sulfonated groups Table 3.4 shows the water uptake values of H-SPEEK and M-SPEEK which was calculated using Equation 9 [149]. The water uptake of compression molded PEEK film was listed as a control experiment. M-SPEEK absorbs less water than H- SPEEK, which is due to less hydrogen bonding sites in M-SPEEK. The water molecules in M-SPEEK are probably located between the cation and the neighboring –SO3¯, therefore the water molecules interact with both metal cations and sulfonated groups. This interaction lead to certain properties of M-SPEEK change such as decreasing Tg, which was not measured in the present research. Na-SPEEK absorbed similar amount of water as H-SPEEK. However, this property was dramatically reduced as the proton of H- SPEEK was substituted by multiple valence cations. Zr-SPEEK almost dose not absorb water at room temperature. One can imagine under the extremely harsh environmental conditions such as high humidity, M-SPEEK will survive and maintain its major properties due to the high stability. W x wet 1 100% (9) Wdry 59 Table 3.4 Wake uptake property of M-SPEEK sample Water uptake(%) Standard deviation H-SPEEK 13.3 0.7 Na-SPEEK 12.9 0.6 Ba-SPEEK 5.63 0.7 Zn-SPEEK 6.73 0.2 Al-SPEEK 7.59 0.4 Zr-SPEEK 0.7 0.06 PEEK 0.4 0.07 3.4.3 Morphology characterization of M-SPEEK It is generally agreed that ion rich aggregates occur in the dry ionomers as a results of microphase separation [150]. This is due to the hydrophilic ionic groups are incompatible with hydrophobic polymer chains. Small angle X-Ray scattering has been proved as an efficient way to detect this microphase separation structure. The ionic aggregations demonstrate themselves as a broad peak in the structure factor on SAXS spectrum, which is believed to be due to a correlation length associated with an average inter-aggregate interference [150]. The microphase separated structure of sulfonic acid ionomer derivatives has been observed for sulfonated polystyrene ionomers by SAXS [151]. However, it is usually difficult to resolve ionic peaks for SPEEK and its adversatives, in which the sulfonate acid groups are attached to the aromatic rings in the backbone [152]. Figure 3.4 shows SAXS results for H-SPEEK and M-SPEEK samples which were run at room temperature. No pronounced scattering peak was observed for H-SPEEK, which is due to insufficient electron density contrast between the ionic domains and the PEEK matrix (Table 3.5) [123]. In the case of M-SPEEK, the sulfonate groups are attached to the aromatic ring in 60 the relatively rigid main chains, which is different from sulfonated polystyrene whose ionic groups are attached to the tertiary carbon in the relatively flexible main chains. It is more difficult for the ionic groups in the M-SPEEK to form multiplets due to the chain rigidity. It is possible that a certain amount of ionic groups may still exist as isolated ionic groups without forming multiplets. In addition, the ionic groups are randomly distributed in the materials because they are randomly attached to the main chain of SPEEK. It is difficult for the ionic groups from different polymer chains to come together and form multiplets due to steric hindrance. As a result, some M-SPEEK formed an imperfect phase separation such as Zr-SPEEK. As shown in Figure 3.4, there is no pronounced peak was observed for Zr-SPEEK. Na-SPEEK exhibits a weak and broad peak at q= 2.6 nm, which corresponds to correlation lengths in real space of d=2.4 nm, where d=2π/q according to Bragg’s law. This result is consistence with Leung and coworkers reported SAXS data for 30mol% sodium sulfonated PEEK sample [153]. The author observed a weak ionomer peak at a Bragg spacing of 2.6 nm, and they did not observe peak of a 10% sodium sulfonate sample. They concluded that the onset of clustering appears between 10 to 30% sulfonation level. Therefore, it is reasonable that we observed a weak ionic peak at d=2.4 nm. Al-SPEEK also shows a very weak and broad peak but at smaller scattering vector q=0.7 nm, which corresponds to a Bragg spacing of 8.9 nm. Ba-SPEEK exhibits the most pronounced peak at q=0.58 nm (d=11 nm). Table 3.5 shows the estimated electron density of PEEK, H-SPEEK, and M-SPEEK using mass density of the material. The mass density, ρ, was determined by measuring volume and mass of a rectangular prism sample at room temperature. Three specimens were measured and the average value was reported. 61 Barium has atomic number 56 and is the fifth element in Group 2. Ba2+ has largest ionic radius among other cation discussed here, see Table 3.1. The electron density in Ba- SPEEK ionic domains is high enough to enable ionic peak to show in SAXS measurement. Zn-SPEEK shows a very weak peak at the scattering vector q=0.52nm which is close to that of Ba-SPEEK. Zn-SPEEK and Ba-SPEEK both have divalent cations, but Zn-SPEEK is more rigid (higher Tg) than Ba-SPEEK. The phase separation in Zn-SPEEK is likely to be less perfect than Ba-SPEEK. Zn-SPEEK Ba-SPEEK Zr-SPEEK I (a.u.) Al-SPEEK Na-SPEEK H-SPEEK 0.0 0.5 1.0 1.5 2.0 2.5 3.0 q (nm-1) Figure 3.4 SAXS of H-SPEEK and M-SPEEK at room temperature 62 Table 3.5 Electron density of PEEK, H-SPEEK and M-SPEEK Electron density sample ρ(g/cm3) (#of electron/cm3) H-SPEEK 1.06 3.3E+23 Na-SPEEK 1.14 3.2E+23 Ba-SPEEK 1.41 4.4E+23 Zn-SPEEK 1.385 4.3E+23 Al-SPEEK 1.2 3.8E+23 Zr-SPEEK 1.28 3.6E+23 PEEK 1.32 4.1E+23 3.4.4 Thermal properties The parent PEEK was a semicrystalline thermoplastic polymer with a Tg of 148°C and a Tm of 338°C. The glass transitions temperature of the starting materials, H-SPEEK and the M-SPEEK ionomers are listed in Table 3.6 and Figure 3.5 shows the heating curves of second cycle from DSC. The introduction of –SO3H groups into PEEK increased Tg, lowered Tm and reduced he crystallinity. For 18 mol% sulfonation, SPEEK was completely amorphous and the Tg was increased about 36°C. Conversion of the sulfonic acid derivative to a metal salt significantly increased Tg than the free acid form with the same degree of sulfonation level as excepted. The Tg increase is a consequence of the restriction of segmental motion by intermolecular hydrogen bonding for the acid derivative and ionic or dipole-dipole interactions for the salts. For the metal salts, microphase separation of ion-rich domains, i.e., ionic clusters, also occurred, which increased Tg further due to the effect of multifunctional crosslinking by the ionic nanodomains, as has been generally observed for ionomers [154-155]. 63 The properties of ionomers, such as Tg, often scale with the Coulomb energy, Ec, of the ion-pair ~q/a in M-SPEEK. Tg of the M-SPEEK ionomers did increase with increasing Ec, see Table 3.6, which is a consequence of stronger interactions of the ionic dipoles as the electrostatic interactions increase. However, the glass transition is broadening as the conversion of H-SPEEK to metal salts (Figure 3.5) and ∆Cp values is decreasing (Table 3.6). This may be a result from the onset of structural fluctuation. As the salt groups act as crosslinking, the chain mobility is significantly reduced, which hindering the clustering process. As a result, the materials formed an imperfect phase separation between the ion-poor and ion rich domains. Some exists as single ion pair, and some are in the form of two ion pairs or higher order aggregates. Ba-SPEEK exhibits relatively lower Tg than other M-SPEEKs with a larger ion radius of a divalent cation. This can be understood by considering the steric effect of the ionic groups in Ba-SPEEK. The ionic clusters formed by Ba-SPEEK ionic groups may not be tightly and closely bound together due to the steric effect. The ion hopping may take place at relatively lower temperatures, which is part of cluster Tg. Tg of the M-SPEEK ionomers can be used as the switching temperature, Tc, for a SMP. This can be varied over a wide range – in this case from 181°C to 288°C – simply by changing the counterion. One can also change the sulfonation level of SPEEK and prepare corresponding metals salts to further change the glass transition temperature. 64 Table 3.6 Glass transition temperature of M-SPEEK -1 ΔCp Sample q/a(Å ) Tg (°C) (J/g) PEEK 148 0.15 H-SPEEK 181 0.27 Na-SPEEK 0.98 229 0.24 Ba-SPEEK 1.34 224 0.080 Zn-SPEEK 2.70 253 0.083 Al-SPEEK 5.61 272 0.11 Zr-SPEEK 4.65 288 0.084 Endo Al-SPEEK Ba-SPEEK Zn-SPEEK Heat Flow (W/g) Flow Heat H-SPEEK Na-SPEEK Zr-SPEEK 50 100 150 200 250 300 350 400 o Temperature ( C) Figure 3.5 DSC for H-SPEEK and M-SPEEK. Second heating curve are showing, heat rate was 10°C/min Figure 3.6 shows the thermalgravity analysis of M-SPEEK ionomers. PEEK and H-SPEEK were also plotted to compare with other M-SPEEK. PEEK is a highly heat resistant material and the onset of the significant weight loss starts from 520°C, which is 65 due to random chain scission of the ether and ketone bonds. The pyrolysis left about 50% carbonaceous char from the aromatic groups at 700°C, which is consistent with the results of Patel el al [156]. The introduction of –SO3H groups into PEEK polymer chain dramatically reduces its thermal stability due to desulfonation [124]. As shown in Figure 3.6, H-SPEEK exhibits two obvious degradation steps. The first weight loss began at ~260°C is the disassociation of sulfonic groups and the other one is the major chain scission. However, after neutralization, the first weight loss step became less pronounced and broadened (Zn-SPEEK, Ba-SPEEK), and eventually disappeared (Na-SPEEK, Al- SPEEK, and Zr-SPEEK). The desulfonation temperature increased to > 300°C. Figure 3.7 shows a TGA scan for Zn-SPEEK with an isothermal step. The sample was first heated to 270°C (at which shape memory tests were run) at a heating rate of 10°C/min, and then was held at this temperature for 60mins to evaluate the thermal stability at high temperature. Less than 1% weight loss was detected and the sample was then heated up to 800°C. This experiment was to demonstrate that M-SPEEK was stable during the shape memory tests. Figure 3.7 also indicates that Zn-SPEEK is thermally stable to 320°C and the major degradation occurs at 380°C (see Figure 3.6 Deriv. Weight curve). Form the TGA study, it can be concluded that M-SPEEK ionomers improved the thermal stability of H-SPEEK and the materials are thermally stable during the further shape memory behavior tests. 66 110 PEEK 100 Al-SPEEK 90 H-SPEEK 80 70 Ba-SPEEK Zr-SPEEK 60 Na-SPEEK Weight remaining (%) remaining Weight Zn-SPEEK 50 40 0 100 200 300 400 500 600 700 800 900 o Temperature ( C) Figure 3.6 TGA of PEEK, HPEEK, and M-SPEEK. Samples were heated under nitrogen environment with a heating rate of 10°C/min 120 1.0 100 0.8 110 C) 80 o 100 0.6 90 80 60 o 70 T=270 C 0.4 60 40 50 Weight remaining (%) remaining Weight 40 0.2 Driev. Weight (%/ Weight remaining (%) 20 30 40 50 60 70 80 90 20 Time (min) 0.0 0 0 100 200 300 400 500 600 700 800 900 Tempearature (oC) Figure 3.7 TGA for Zn-SPEEK. Sample was held at 270°C for 60mins and less than 1% weight loss was observed (inset picture). 67 3.4.5 Linear Viscoelasticity The linear viscoelastic (LVE) tensile properties of the neat PEEK, M-SPEEK ionomers are shown in Figure 3.8 and 3.9. The experiments were run only to 300°C to avoid desulfonation of the sample. Below Tg, the dynamic modulus, E’, of the M-SPEEK and PEEK were similar, but the addition of the metal sulfonate groups increased Tg. Above Tg, E’ for PEEK decreased from ~2 GPa to ~200 MPa, and exhibited a rubber-like region, though the modulus decreased slowly with increasing temperature in that region. The rubbery region is due in part to molecular entanglements, but more important with regard to the modulus are the physical crosslinks provided by the crystalline (solid) phase. The region where PEEK crystals melt is not shown in Figure 3.8, but E’ decreased rapidly above Tm. The viscoelastic behavior of neat PEEK suggests that it may exhibit shape memory behavior if the crystalline phase can effectively function as a “permanent” network and Tg is Tc. This will be discussed further in the shape memory section below. 10000 Zn-SPEEK Ba-SPEEK 1000 Zr-SPEEK PEEK 100 Al-SPEEK 10 Storage Modulus (MPa) Modulus Storage Na-SPEEK 1 0 50 100 150 200 250 300 350 Temperature ( C) Figure 3.8 Storage modulus of M-SPEEK as a function of temperature 68 1000 Zn-SPEEK PEEK Ba-SPEEK 100 Zr-SPEEK 10 Al-SPEEK Loss Modulus (MPa) 1 Na-SPEEK 0.1 0 50 100 150 200 250 300 350 400 Temperature ( C) Figure 3.9 Loss modulus of M-SPEEK as a function of temperature The neat, amorphous ionomer (M-SPEEK) samples remained glasslike (E’ ~ 1-2 GPa) to much higher temperatures than the PEEK due to the increase of Tg, For the ionomers, E’ decreased much more rapidly above Tg. For Na-SPEEK, Ba-SPEEK and Zn-SPEEK, the experiment could not be run above 300°C, because of desulfonation, so it was not possible to accurately measure the modulus of the rubbery region above Tg or a temperature when viscous flow became significant. However, the samples exhibited rubber-like behavior during the shape memory cycles as will be discussed later in this chapter. The E’ data in Figure 3.8 for Na-SPEEK suggest a leveling of the modulus at about 290°C. Those observations lead us to speculate that the rubbery modulus was ~1 MPa, which is a reasonable value for a crosslinked elastomer. The rubbery plateau can be readily seen for Al-SPEEK and Zr-SPEEK. For Al-SPEEK, the rubbery modulus was ~10MPa while Zr-SPEEK exhibits a rubbery modulus in the order of 200MPa. It is also 69 noticed that the dynamic moduli curve for Zr-SPEEK was similar to that of PEEK, but the curve was shifted to higher temperature. The rubbery region of Al-SPEEK and Zr- SPEEK can also partially attribute to the polymer chain entanglement, but more importantly related to the ion crosslinking. In this case, the crosslinks are due to the ionic interactions and the multifunctional nanodomains, and these results suggest that the neat ionomers may exhibit shape memory behavior using their supramolecular ionic structure to support a permanent network and Tg as Tc. M-SPEEK also showed increased tanfrom PEEKsee Figure 3.10. The increase of tan suggests the ability of the physical crosslinks formed by the ionic interactions of M-SPEEK can relax and dissipate energy. Tandecreases as the increasing of metal ion valence, and this is related to the crosslink density in M-SPEEK. 0.6 Na-SPEEK 0.5 Ba-SPEEK 0.4 Al-SPEEK 0.3 Tan Zr-SPEEK 0.2 Zn-SPEEK 0.1 PEEK 0.0 0 50 100 150 200 250 300 350 400 Temperatrue (oC) Figure 3.10 Tanδ of PEEK and M-SPEEK as a function of temperature 70 An advantage of developing a shape memory polymer from PEEK is the thermoplastic nature of the polymer. PEEK can be melt processed at temperatures in excess of 340C, but the ionomers will degrade at such high temperatures. The ionomers, however, are amorphous, so in principle they should be processable above Tg. Although the LVE data indicated that above Tg, these materials are still solid-like (E’ > E”), the high stresses that are used in conventional polymer processing operations such as extrusion, compression molding and injection molding might be sufficient to induce viscous flow of the melt. That is the case for other ionomers, e.g., moderately sulfonated polystyrene where the LVE behavior is characteristic of solid-like behavior, but the dynamic and loss moduli crossover at some critical strain rate or stress. 3.4.6 Dynamic shear The effect of ionic interactions on the viscoelastic behavior of M-SPEEKs under dynamic shear stress was studied. The samples were heated to the temperature at which shape memory performance were assessed, and the dynamic shear moduli, G' and G" were measured at frequencies from 0.01 to 100 rad/s. The measurements were made within the linear viscoelastic limit, which was determined from strain sweep experiments. H-SPEEK was not investigated in this experiment because at this high temperature H- SPEEK degrades. Figure 3.11 and 3.12 show the shear modulus of M-SPEEK as a function of frequency at constant temperatures. Na-SPEEK, Zn-SPEEK, and Ba-SPEEK were run at 270°C . Al-SPEEK and Zr-SPEEK were run at 300°C because these two materials have higher Tg. In the range of the experiment frequencies, the materials were solid like and no significant flow was observed at the temperature above their Tg. Storage modulus roughly 71 increases with the increasing of ion pairs’ Coulomb energy. Na-SPEEK showed the lowest storage modulus at the entire test frequencies, which is due to the relatively weaker physically crosslinked network, e.g. low q/a and no salt bridge forming. Other M- SPEEK ionomers, in which the metal ion associates with more than one sulfonate groups, have two possible crosslinking mechanisms: 1) ionic crosslinking which occurs in order to achieve electron neutrality and 2) dipole-dipole associations of M-sulfonate groups which produce ionic clusters. The ionic crosslinking is likely to be independent of temperature. Ba-SPEEK showed similar behavior as Zn-SPEEK with smaller storage modulus. At 300°C , which is above both Al-SPEEK and Zr-SPEEK’s Tg, Zr-SPEEK was still the most rigid material, which is due to its strong crosslinked structure. Al-SPEEK also exhibits high shear modulus, but the material tend to flow at very low frequencies, e.g. <0.1 rad/s. The ionic crosslinking of Al-SPEEK is weaker than that of Zr-SPEEK. q/a 107 Al-SPEEK Zr-SPEEK Ba-SPEEK Zn-SPEEK 106 G' (Pa) Na-SPEEK 105 0.001 0.01 0.1 1 10 100 1000 (rad/s) Figure 3.11 Storage Modulus of M-SPEEKs as a function of frequency 72 Table 3.7 Summary of the shear modulus ( G ′ ) at frequency=1 rad/s and molecular weight between crosslinks ( Mc ) at 270 and 300 °C. G' Mc Sample q/a(Å-1) ρ (g/cm3) (MPa) (g/mol) Na-SPEEK 0.98 1.14 1.16 4437a Ba-SPEEK 1.34 1.41 2.20 2894a Zn-SPEEK 2.70 1.39 3.13 1998a Al-SPEEK 5.61 1.20 6.22 919b Zr-SPEEK 4.65 1.28 9.93 614b a T=270°C b T=300°C The crosslinking density of M-SPEEK was estimated by calculating the molecular weight between crosslinks Mc using the classical equation of rubber elasticity [157]: (10) where Gn is the shear modulus in the plateau region. Here, the dynamic shear modulus G' at frequency=1rad/s was used to calculate the crosslinking density. ρ is the mass density of M-SPEEK, which was estimated by measuring volume and mass of a rectangular prism sample at room temperature. Three specimens were measured and the average values are listed in Table3.7. No adjustment was made for the temperature dependence of density. R is the gas constant and T is the absolute temperature. Table 3.7 summarized Mc value for M-SPEEK at high temperatures. The increase in modulus and decrease in Mc with the increasing of cation’s valence is due to the increasing of ionic crosslinking density. The decrease in tanδ value (Figure 3.10) also showed that the ionic crosslinking density increases with the increasing of cation’s valence. The material behaved more 73 elastic as the crosslinking density increases. Ba-SPEEK and Zn-SPEEK both have divalent cation, but Zn-SPEEK showed larger crosslinking density than that of Ba- SPEEK. This is due to the dipole-dipole interactions in Zn-SPEEK are stronger. Zr-SPEEK Ba-SPEEK 106 Zn-SPEEK G'' (Pa) Al-SPEEK Na-SPEEK 105 0.001 0.01 0.1 1 10 100 1000 (rad/s) Figure 3.12 Loss Modulus of M-SPEEKs as a function of frequency Frequency sweep tests at different temperatures from 260 to 300°C (Figure 3.13 and 3.14) were performed for Na-SPEEK under a frequency range of 0.01-100 rad/s. Na- SPEEK was chosen to study the temperature dependent linear viscoelastic behavior because Na-SPEEK has relatively lower Tg and showed tendency to reach its terminal flow region at high temperatures in Figure 3.13. Time temperature superposition (TTS) master curve of G' and G" were constructed at reference temperature of Tr=280°C using ARED G2 TRIOS software. Vertical corrections were made by multiply the modulus by a factor βT to compensate for the modulus temperature dependence. The shift factors were computed using the WLF equation as shown in Table 3.8. 74 Figure 3.15 shows the attempted superposition of the data at different temperature for Na-SPEEK. The TTS failed for this material in the low range of reduced frequencies. For sulfonated polystyrene metal salts ionomers, TTS worked well [158]. In that case, the material is only lightly sulfonated and the sulfonation level is less than 6 mol%. Na- SPEEK is 18mol% sulfonated and the material is nanophase separated. The relaxation times for the two phases are not that far apart that two mechanisms were involved in a single experiment. Table 3.8 Shift factors of Na-SPEEK at 280°C Temperature α (x variable) β base ∆β β (y variables) (°C) T T T T 260 2341 1.04 10 10 270 233 1.02 10 10 280 1.0 1.0 1.0 1.0 290 6.3e-4 0.98 0.63 0.61 300 1.37e-5 0.97 0.32 0.30 108 280 7 10 270 260 G' (MPa) 106 300 290 105 0.001 0.01 0.1 1 10 100 1000 (rad/s) Figure 3.13 Storage modulus of Na-SPEEK as a function of frequency at different temperatures 75 108 280 107 290 260 106 270 G'' (MPa) 105 300 104 0.001 0.01 0.1 1 10 100 1000 (rad/s) Figure 3.14 Loss modulus of Na-SPEEK as a function of frequency at different temperatures 8 8 7 7 G' (Pa) G' 6 6 (Pa) G'' Log LOG 5 5 4 4 -8 -7 -6 -5 -4 -3 -2 -1 0 1 2 3 4 5 6 Log aT Figure 3.15 G' (○) and G" (▼) master curves for Na-SPEEK. Reference temperature is 280°C 76 3.4.7 Shape memory behavior Shape memory behavior were evaluated using dynamic mechanical analysis (DMA) with appropriate programming procedures, see Figure 3.16. A stress-strain profile was obtained as a function of temperature and time. The shape fixation and recovery efficiencies were calculated using the following equations [1,2]. )N( R pm 100%x (11) pm 1)N( )N( )N(F u 100%x (12) m where εm, εp, εu and N denote the strain after the stretching step (before cooling or unloading the sample), the strain after recovery, the strain in the fixed temporary shape, and the cycle number, respectively. A value of F or R of 100% represents complete strain fixing or recovery. Figure 3.16 Schematic illustration of shape memory test programming The thermally activated shape memory behavior of neat PEEK is shown in Figure 3.17. PEEK crystals provided the permanent network and the glass transition served as switching temperature Tc. Although the neat PEEK exhibited some shape memory behavior, only 28% of the deformation was fixed by the temporary network (i.e., the glassy amorphous phase) and the shape recovery efficiency was only recover 35%. The 77 reason for the poor shape memory behavior was the extensive relaxation of the strain during the cooling step following the initial stretching of the sample at high temperature, see path 2 in Figure 3.16. During the cooling step, the sample was held at a constant tensile stress, 0.32 MPa, but the stress on the “network” chains increased as the sample was cooled, see the E’ data for PEEK in Figure 3.8. Apparently, the “permanent” network formed by the crystalline PEEK phase was insufficient to maintain the sample length as the internal stress increased, which led to the large strain relaxation. The change in strain at Tg in Figure 3.8 also clearly demonstrates that point. 3.0 2 2.5 3 2.0 1 1.5 4 1.0 Strain (%) Strain E 0.5 4 S 3 0.0 2 0 1 50 100 0 Temperature150 ( Stress (MPa) 200 C) 250 Figure 3.17 Shape memory cycle for PEEK. Point S denotes the start of the cycle. Path 1: The sample was heated to 200°C (Tc ~150°C) and stretched with a stress of 3.8 MPa. Path 2: The deformed sample was cooled under a constant stress of 3.8 MPa. Path 3: The stress was to set the temporary deformed shape. Path 4: The sample was reheated to 200°C to recover the permanent shape. Point E denotes the end of the cycle. In contrast to the neat PEEK, H-SPEEK and the M-SPEEK ionomers were amorphous and the “permanent” network was formed by hydrogen bonding of the sulfonic acid groups or ionic or dipole interactions for the salts. The shape memory 78 behavior of the neat H-SPEEK was still not very good as shown in Figure 3.18, in that the hydrogen bonding exhibited excessive creep under load. The material can only fix 54% temporary strain and recover 49% permanent shape. The ionic interactions worked much better than the hydrogen bonded network or the crystalline network in PEEK. 4 3 2 Strain (%) Strain 1.4 1 E 1.2 1.0 S 0.8 0 0.6 50 0.4 100 150 0.2 Temperature 200( o 250 0.0 Stress (MPa) C) 300 Figure 3.18 Shape memory cycle for H-SPEEK. The samples were stretched at 220°C with a stress of 1.1 MPa M-SPEEKs show much better shape memory behavior than both PEEK and H- SPEEK. M-SPEEKs still used glass transition as the switching temperature, however, the permanent work was provided by the ionic nanophase separated domains arising from the ionic aggregations, which act like multifunctional crosslinking points. Figure 3.19 shows one shape memory cycle of Zn-SPEEK. The sample was deformed at 270°C, which is 20°C higher than its Tg. At this temperature, the material was easily stretched to a certain strain (11%). The material was able to fix the temporary deformed shape after cooling 79 and unload process. Upon reheating, the sample started to recover as the temperature approaching Tg and recovered back to its original shape almost completely under stress free condition. The similar test protocol was performed for all the M-SPEEKs samples. The shape fixing and recovery ratio, which can be calculated from the strain-temperature profile, is summarized in Table 3.9. Multiple shape memory cycles were conducted for each sample to assess the reproducibility. As most of thermoplastic shape memory materials, the first cycle shows relatively lower values, which is due to the residue strain from the sample processing, e.g. compression molding and solution casting. Na-SPEEK exhibit relatively poor shape memory behavior among other M-SPEEK, but still better than PEEK. Na+ is a monovalent cation and can only associate with one sulfonic group. The permanent network supported only by the dipole-dipole interactions is less robust than those can form salt bridges, see Figure 3.20. Figure 3.23 summarized shape memory performance of M-SPEEKs related to the Coulomb energy of ion pairs. Na-SPEEK and Zn-SPEEK exhibit large creep of their ionic domains and can only fix less 90% strain. This was not case for Ba-SPEEK, which showed good shape fixing about 93%. Ba- SPEEK also exhibited excellent shape recovery, ~94%. This is a solution to improve the fixing efficiency of Na-SPEEK and Zn-SPEEK which will be discussed in the next Chapter. 80 Shape fixing 14 12 10 8 Shape 6 recovering Strain(%) End 4 1.2 1.0 2 0.8 Start 0 0.6 50 0.4 100 150 0.2 Temperature200 ( 250 0.0 300 Stress (MPa) C) Figure 3.19 One shape memory cycle of Zn-SPEEK, sample was stretched at 270°C with a stress of 1MPa. Figure 3.20 Intermolecular inter-actions that provide physical crosslinks in ionomers. Al-SPEEK and Zr-SPEEK can combine with three and four sulfonate groups coming from different polymer chains, therefore these materials are crosslinked by the metal ions. The linear viscoelastic behavior of these materials exhibited apparent rubbery 81 plateau above their Tg, which suggesting a dual shape memory behavior, see Figure 3.8. These “covalent like” crosslinking resulted in a much stronger permanent network and improved shape fixing ability. Figure 3.21 shows one shape memory cycle of Al-SPEEK with a switching temperature Tc=Tg= 275°C. The material was stretched at 300°C with an external stress of 0.5MPa. The sample was cooled to 30°C under the stress and then the stress was removed. The ionic crosslinked network of Al-SPEEK enabled the material to fix 94% strain. When reheated to 300°C, the sample recovered to its 90% of permanent shape. Zr-SPEEK exhibited similar excellent shape memory behavior with a slightly lower Tc, see Figure 3.22 and Table 3.9. Since the glass transition of Al-SPEEK and Zr- SPEEK are much higher than the M-SPEEK with mono- or di-valent cations, one has to deform the material above 300°C. Na-SPEEK and Zn-SPEEK exhibit large creep of their ionic domains and can only fix less 90% strain. Shape fixing 35 30 25 Shape recovery 20 End Strain (%) Strain 0.6 15 0.5 Start 0.4 10 0.3 0.2 50100 150 200 0.1 Temperature250 ( 0.0 Stress (MPa) 300 350 oC) Figure 3.21 One shape memory cycle of Al-SPEEK, sample was stretched at 300°C with a stress of 0.5MPa 82 Shape fixing 52 50 48 46 Shape recovery 44 Strain (%) Strain End 5 42 4 3 40 Start 2 50 100 1 150 200 Temperature250 ( 0 Stress (MPa) 300 350 oC) Figure 3.22 One shape memory cycle of Zr-SPEEK, sample was stretched at 300°C with a stress of 4.2 MPa 100 100 95 90 Al Zr 90 Ba 80 Zn 85 70 F (%) F R (%) 80 R 60 75 50 Na 70 40 0 1 2 3 4 5 6 q/a Figure 3.23 Shape fixing and recovery ratio for M-SPEEKs 83 Table 3.9 Shape memory properties of M-SPEEKs Deformation Material q/a(Å-1) T (oC) F(%) R(%) stress (MPa) c 79 32 Na-SPEEK 78.4 46 0.98 0.32 240 (4 continuous cycles) 78.4 43.9 78.6 44.3 88.3 63.3 Zn-SPEEK 89.3 92.5 2.7 1 250 (4 continuous cycles) 88.7 98.2 88.4 99.1 Ba-SPEEK 93.9 70.9 1.34 0.3 246 (2 continuous cycles) 92.9 94 Al-SPEEK 95 63.3 5.61 0.5 275 (2continuous cycles) 94 90 96.8 29.5 Zr-SPEEK 4.65 4.2 260 96 94.0 (3 continuous cycles) 96 93.9 3.5 Conclusions PEEK was functionalized by the introduction of –SO3H into one of the benzene ring in the polymer main chain to dramatically decease the crystallinity. Metal salts of sulfonated PEEK (M-SPEEK) were prepared by neutralizing the 18 mol% sulfonated PEEK to metal ions (Na+, Zn2+, Ba2+, Al3+, and Zr4+). M-SPEEKs were amorphous materials and exhibited significantly increase glass transition temperatures due to dipole- dipole interactions or ionic crosslinking of M-SPEEKs. Tg of M-SPEEK increases as a function of the ion pair’s Coulomb energy. M-SPEEKs exhibited moderate to excellent dual shape memory behavior depending on the coulomb energy of ion pairs. In M- SPEEK, the permanent network was provided by the ionic nanodomains due to ionic or dipolar interactions between metal sulfonated groups and the glass transition temperatures served as the switching temperature. Na-SPEEK showed moderate shape 84 memory behavior with a shape fixing of 80% and recovery of 46%. Zn-SPEEK showed much improved performance with a fixing of 90% and recovery efficiency as high as 99%. The ionic bonds formed by the Zn-sulfonates were more efficient at shape memory behavior than were dipole-dipole interactions formed from monovalent cations. Ba- SPEEK, Al-SPEEK and Zr-SPEEK showed excellent shape fixing efficiency (~93%) due to more robust permanent network, and almost complete shape recovery (~94%). The trivalent and tetravalent cation neutralized SPEEK exhibited improve thermal stability and also higher switching temperatures due to high Tg. Al-SPEEK and Zr-SPEEK exhibited “covalent like” crosslinked structure by a high rubbery plateau modulus at high temperature and low water uptake. 85 CHAPTER IV DUAL AND TRIPLE SHAPE MEMORY BAHAVIOR OF M-SPEEK/FATTY ACID SALT COMPOUNDS 4.1 Introduction This chapter presents the dual and triple shape memory behavior of blend compounds: Na-SPEEK/NaOl and Zn-SPEEK/NaOl. The effects of low molar mass compound sodium oleate (NaOl) on thermal and mechanical properties, morphology, and shape memory behavior of M-SPEEK were investigated. As discussed in Chapter III, metal salts of sulfonated PEEK exhibited moderate dual shape memory behavior, where the glass transition temperature of M-SPEEK served as the switching temperature and ionic nanodomains due to the ionic aggregation provided the permanent network. The temporary network, which formed by the glassy state of M-SPEEK was not efficient enough (80-90%), especially for Na-SPEEK and Zn-SPEEK. The problem of creep recovery of the temporary network for the neat M-SPEEK ionomers can be resolved by adding a high melting point fatty acid compound, NaOl. This approach to produce a shape memory polymer is identical to that described in refs. [159, 74] except that in this Chapter a high temperature ionomer and a high melting point fatty acid salt were used. As we introducing a temporary network into the M-SPEEK polymer matrix, the materials have two reversible thermal transitions: glass transition of M-SPEEK and 85 melting of NaOl crystals. Therefore, there is the possibility for tripe shape memory using the Tm of the NaOl and the Tg of the ionomer matrix to achieve two distinct switching temperatures, Tc1 and Tc2. 4.2 Experimental Details 4.2.1 Preparation of M-SPEEK/NaOl compounds Film samples of M-SPEEK and shape memory compounds of 70 parts M-SPEEK and 30 parts sodium oleate (NaOl, assay >99%, Aldrich Chemical Company), were prepared by dissolving the M-SPEEK and NaOl in refluxing N-methyl-2-pyrrolidone (NMP). The solution was cast onto a clean glass plate and dried at 80C for 48 h to remove most of solvent. The film samples were then dried at 150°C under vacuum for 24h prior to characterization. The sample notation used in this paper for the ionomer/FAS compound is M-SPEEK/NaOl(30), where M stands for the cation and 30 denotes the weight percentage of NaOl in the compound. 4.2.2 Thermal analysis Thermal properties such as glass transition, melting temperature, and thermal degradation temperature were determined using TA instruments DSC Q200 and TGA Q50. The tests details were described in Chapter III. 4.2.3 Wide angle X-ray diffraction (WAXD) WAXD of M-SPEEKs and M-SPEEK/NaOl(30) compounds were measured with a Bruker D8 Discover X-ray diffractometer with a general area detector diffraction system (GADDS) using copper radiation (λ=0.154 nm) at the room temperature with an exposure time of 5 mins. The distance between the diffraction patterns was calculated 86 from Bragg’s law, d=2π/q, where the scattering vector q=4πsin(θ)/λ and θ is one half of the scattering angle. 4.2.4 Mechanical tests The tensile stress-stain behavior was measured with an Instron model 1101 universal testing machine using a 100 N load cell. Film samples were cut into dog-bone specimens with gauge length dimensions of 7.3 × 3.3 × 0.3 mm. The tests were conducted at room temperature with a crosshead speed of 2 mm/min. The viscoelastic properties of films were measured with a TA Instruments dynamic mechanical analyzer (DMA) Q800. Strain sweep was performed to determine the linear viscoelastic region. The film sample was then deformed in the tensile mode with an oscillatory strain (0.2%), a frequency of 1 Hz, and a temperature ramp from room temperature to 300C of 3C /min. 4.2.5 Dual shape memory cycle test Shape memory effects were assessed using a tensile film fixture and the controlled force mode of the DMA. Before shape memory cycle test, a stress-strain test was performed to determine the extent of elongation and corresponding stress at temperature where deformation would take place during the shape memory cycle test. For the shape memory cycle test, the film sample was first heated with a preload force of 0.005N from room temperature to 270C, which is higher than the melting temperature of NaOl crystals. The sample was allowed to equilibrate at 270C for 5mins, the force was then increased. The sample was stretched isothermally to about 60% strain. Once the strain equilibrated, the sample was cooled rapidly under load to 30C to fix the temporary shape. The force was then lowered to 0.005 N, which was sufficient to prevent sagging of 87 the sample when it was reheated above Tm,NaOl. The shape recovery of the sample was achieved by reheating it to 270C with a constant force of 0.005 N and holding it isothermally at 270C for 20 min to allow the strain to equilibrate. The sample was then cooled to 30C to finish one shape memory cycle. Four continuous cycles were conducted to evaluate the reproducibility. 4.2.6 Triple shape memory cycle test Tripe shape memory effects were evaluated using the tensile film fixture and stress controlled mode of DMA similar as dual shape memory cycle tests. Since the glass transition temperature of M-SPEEK and melting temperature of NaOl crystals are separated, it is possible to observe triple memory behavior for both Na-SPEEK/NaOl and Zn-SPEEK/NaOl. The sample was first heated 270C under preload force 0.005N, which is higher that both thermal transition temperatures. After equilibrated for 5 mins at this temperature, the sample was stretched isothermally to about 20% strain. Once the strain equilibrated, the sample was cooled rapidly to Tg,MSPEEK< T< Tm, NaOl. The sample was allowed to equilibrate for 5 mins and then stretched to about 60% strain isothermally. After the strain equilibrated, the sample was cooled rapidly to 30C to fix the temporary shape. The step shape recovery was achieved by heating the sample first to Tg,MSPEEK< T< Tm, NaOl for 20 min to allow the first temporary shape recover and then heating to 270C for another 20mins. Three continuous cycles were conducted to evaluate the reproducibility. 88 4.3 Results and discussion 4.3.1 Thermal properties The thermal transitions of the starting materials, the M-SPEEK ionomers and the M-SPEEK/NaOl compounds are listed for comparison in Table 4.1. As discussed in Chapter III, sulfonation increased Tg, lowered Tm and reduced the crystallinity. In this research, the starting material PEEK possess crystallinity as high as 30%, which was calculated using the heat of fusion obtained from DSC comparing to the literature value (130J/g) [160]. The 18 mol% sulfonated SPEEK was completely amorphous and the Tg was 181°C, which was 36°C higher than PEEK. Na-SPEEK and Zn-SPEEK were also amorphous material, but showed further increased Tg, which is a consequence of the restriction of segmental motion by ionic or dipole-dipole interactions for the salts. In M- SPEEK, microphase separation of ion-rich domains, i.e., ionic clusters increased Tg further due to the effect of multifunctional crosslinking by the ionic nanodomains. As shown in Table 4.1, Zn-SPEEK increased Tg of H-SPEEK much higher than Na-SPEEK, which is a consequence of stronger interactions of the ionic dipoles as the electrostatic interactions increase. Sodium oleate (NaOl) is a low molar mass crystalline compound that has a melting temperature as high as 259°C. NaOl belongs to the family of fatty acid salts, which have polar characteristic and expect to be compatible with M-SPEEK ionomers. 89 Table 4.1 Thermal characteristics of materials T a Relative Materials q/a g T (°C) (°C) m crystallinity (%) NaOl -- -- 259 PEEK -- 148 338 30b H-SPEEK -- 181 -- 0 Na-SPEEK 0.91 229 -- 0 Zn-SPEEK 2.86 253 -- 0 Na-SPEEK/NaOl(30) -- 176 256c 36c Zn-SPEEK/NaOl(30) -- 208 250c 47c a Tg of the M-SPEEK b calculated from HDSC/H100% crystal c Tm and relative crystallinity of the NaOl in the compound (HNaOl in cmpd/HNaOl) The addition of NaOl to M-SPEEK reduced the Tg of the composite films by about 20%, which indicates some miscibility of the fatty acid salt and the ionomer. That was not unexpected, since other studies have reported some miscibility of compounds of fatty acids or fatty acid salts with ionomers [161]. The melting temperature of the NaOl crystals in the compounds was also lower in both compounds, which further supported some miscibility. The crystallinity of the NaOl in the two compounds was significantly suppressed from that of the pure NaOl, Table 4.1, which suggests that partial miscibility or strong interactions between the metal carboxylate groups (FAS) and the metal sulfonate groups (ionomer) restrained crystallization. The relative crystallinity of NaOl in the compounds was estimated from the ratio of ∆H/∆HNaOl, where ∆H was the measured value for NaOl in the compound and ∆HNaOl was the corresponding experimental value for the neat NaOl. 90 100 80 60 HSPEEK Na-SPEEK 40 Zn-SPEEK NaSPEEK/NaOl ZnSPEEK/NaOl Mass Remaining (%) 20 PEEK NaOl 0 0 100 200 300 400 500 600 700 800 Temperature ( C) Figure 4.1 TGA curves for PEEK, H-SPEEK, M-SPEEK and M-SPEEK/NaOl(30). Experiments were run using a nitrogen atmosphere Figure 4.1 shows the thermal stability of starting material, H-SPEEK, M-SPEEKs and M-SPEEK/NaOl(30) compounds. PEEK is thermally stable up to ~500°C, while sulfonation reduced it to 270°C, which is due to desulfonation. However, the metal salt ionomers improved the thermal stability to >300°C. The desulfonation step was broadened and shift to higher temperature for Na-SPEEK. For Zn-SPEEK the desulfonation step was not obvious. Neat NaOl was thermally stable to ~400°C and exhibited a single decomposition step, which is due to the scission of hydro carbon chain. M-SPEEK/NaOl(30) shows similar decomposition behavior as M-SPEEKs. Although, there was no desulfonation step observed, the material still started to degrade >300°C. Therefore, the stability of the sulfonate group is the limiting factor in these compounds. Thus, the limiting use temperature for the ionomers and the ionomer/NaOl compounds was ~300°C. 91 4.3.2 Wide angle X-ray diffraction Zn-SPEEK NaSPEEK Na-SPEEK/NaOl I (a.u) Zn-SPEEK/NaOl NaOl 0 5 10 15 20 25 30 35 40 2 Figure 4. 2 WAXD for M-SPEEK and M-SPEEK/NaOl Figure 4.2 shows wide angle X-ray diffraction for neat NaOl, M-SPEEK, and M- SPEEK/NaOl(30) compounds. The neat NaOl was a crystalline powder with a high degree of order. Six distinct peaks were resolved in the neat NaOl diffraction pattern with a periodical long spacing d=4.41nm, which is consistent with the 4.46nm reported by Ross and McBain [162]. The M-SPEEK ionomers were amorphous. M-SPEEK/NaOl (30) showed both NaOl crystalline peak and the amorphous peak of M-SPEEK. The NaOl crystalline peaks in the compounds are broad and less sharp, which may be evidence that partial NaOl was miscible with the ionomer and crystallinity was reduced. From Figure 4.2, one can conclude that Zn-SPEEK/NaOl (30) compound has more NaOl crystal content and the crystals form a separated phase in Zn-SPEEK polymer matrix, because its diffraction pattern is more analogous to neat NaOl and retain most of NaOl crystal’s 92 distinct peaks. This is consistent with the DSC data that the crystallinity of NaOl in Zn- SPEEK/NaOl (30) is higher than that in Na-SPEEK/NaOl (30), see Table 4.1. 4.3.3 Mechanical and Viscoelastic Properties Example engineering tensile stress versus strain curves for the neat PEEK, ionomers and the compounds are shown in Figure 4.3, and the tensile properties are summarized in Table 4.2. The Na-SPEEK and Zn-SPEEK films were relatively brittle at room temperature, but the M-SPEEK/NaOl(30) composite films were more ductile. The properties of PEEK film was also measured for comparison. Since PEEK is not soluble in any convenient solvent for casting film, the PEEK film was prepared by compression molding. The semi-crystalline PEEK was the most brittle of the materials used in this study. 100 80 PEEK ZnSPEEK 60 NaSPEEK NaSPEEK/NaOl (MPa) 40 ZnSPEEK/NaOl 20 HSPEEK 0 0 10 20 30 40 50 (%) Figure 4. 3 Engineering tensile stress versus strain curves at room temperature for neat PEEK, M-SPEEK and M-SPEEK/NaOl(30) compounds. 93 The sulfonation of PEEK lowered the tensile modulus and the yield stress, but increased the ultimate elongation. The lower modulus and yield stress are probably a consequence of the elimination of the crystallinity in the ionomer. The higher elongation is probably due to the physical network formed by association of the ionic species and microphase separation of the ionic species. These behave as physical, reversible crosslinks that provide a mechanism for dissipating strain energy and increase area under the stress – strain curve, i.e., the toughness – herein defined as strain energy per unit volume absorbed by the material. Table 4.2 Engineering tensile propertiesa of M-SPEEK and M-SPEEK/NaOl(30) Compounds E σ b ε c Sample u u (MPa) (MPa) (%) PEEK 2400 ± 45 87 ± 5.6 7.3 ± 0.47 H-SPEEK 1050±72 39±0.3 3.5±0.1 Na-SPEEK 1600 ± 56 43 ± 8.5 9.3 ± 1.2 Zn-SPEEK 2200 ± 27 61 ± 6.6 15 ± 2.8 Na-SPEEK/NaOl(30) 920 ± 87 44 ± 11 24 ± 7.9 Zn-SPEEK/NaOl(30) 1300 ± 54 40 ± 2.1 45 ± 4.7 a average and standard deviation of five specimens for each sample b stress at yield c strain at break Zn-SPEEK was stiffer than Na-SPEEK, which may be a consequence of the divalent cation, which provides a salt bridge between two sulfonate groups, as opposed to a dipole-dipole interactions of sulfonate groups in Na-SPEEK. The addition of NaOl to the ionomers significantly lowered the modulus and ultimate strength, to about 40– 55% of the modulus and 45 – 50% of the yield strength of neat PEEK film, but it also greatly improved the ductility of the film. Those results are consistent with the conclusion that 94 some NaOl was miscible with the ionomers and acts as a plasticizer with regard to the mechanical properties. The linear viscoelastic tensile properties of the neat PEEK, M-SPEEK ionomers and the M-SPEEK/NaOl compounds are shown in Figure 4.4. The experiments were run only to 300°C to avoid desulfonation of the sample. Below Tg, the dynamic modulus, E’, of the M-SPEEK and PEEK were similar, but the addition of the metal sulfonate groups increased Tg and tan . The increase of tan is consistent with the more ductile nature of ionomers and is due to the ability of the physical crosslinks formed by the ionic interactions to relax and dissipate energy. The PEEK sample also had a physical network, in this case the crystalline phase, but that does not contribute any significant energy dissipation below Tg. 104 1.0 (a) (b) Zn-SPEEK Na-SPEEK 0.8 103 ZnSPEEK/NaOl NaSPEEK/NaOl 0.6 102 E' PEEK tan 0.4 Tan E'(MPa) 101 0.2 100 0.0 0 50 100 150 200 250 300 350 0 50 100 150 200 250 300 350 T (C) T (C) Figure 4.4 Dynamic and loss tensile modulus versus temperature for (a) PEEK, ZnSPEEK, and Zn-SPEEK/NaOl(30); (b) NaSPEEK and Na-SPEEK/NaOl (30). The frequency was 1 Hz. 95 Above Tg, E’ for PEEK decreased from ~2 GPa to ~200 MPa, and exhibited a rubber-like region, though the modulus decreased slowly with increasing temperature in that region. Tan exhibited a maximum at the glass transition. The rubbery region is due in part to molecular entanglements, but more important with regard to the modulus are the physical crosslinks provided by the crystalline (solid) phase. The region where PEEK crystals melt is not shown in Figure 4.4, but E’ decreased rapidly above Tm and tan increased. Tan for the two neat ionomers (Na-SPEEK and Zn-SPEEK) in Figure 4.4 showed a small peak near 100°C. The origin of this relaxation is probably due to some residual water in the sample, but that was not investigated in this study. This relaxation was not seen in the PEEK or the composite samples. The PEEK is relatively hydrophobic and in the composites, the interactions between the Na-carboxylate from the fatty acid and the metal sulfonate groups are expected to occupy hydrogen bonding sites that would normally be taken up with water. The neat, amorphous ionomer (M-SPEEK) samples remained glasslike (E’ ~ 1-2 GPa) to much higher temperatures than the PEEK due to the increase of Tg, For the ionomers, E’ decreased much more rapidly above Tg. Unfortunately, the experiment could not be run above 300°C, because of desulfonation, so it was not possible to accurately measure the modulus of the rubbery region above Tg or a temperature when viscous flow became significant. However, the shape memory behavior of the ionomers, the samples exhibited rubber-like behavior during the shape memory cycles and the E’ data in Figure 4.4(b) for Na-SPEEK suggest a leveling of the modulus at about 290°C. Those observations lead us to speculate that the rubbery modulus was ~1 MPa, which is a 96 reasonable value for a crosslinked elastomer. In this case, the crosslinks are due to the ionic interactions and the multifunctional nanodomains. The viscoelastic behavior of compounds of the ionomers with NaOl, Figure 4.4, suggests that those materials may exhibit not only dual shape memory behavior, but also triple shape memory. The glassy modulus of the compounds was 50 – 75% of the modulus for the neat PEEK (E’ ~ 1.4 GPa for the Zn-SPEEK compound and 1.0 GPa for the Na-SPEEK compound), and the Tgs of the compounds were significantly lower than for the neat ionomers, which was likely to the limited solubility of NaOl in the ionomers. One then expects two specific mechanical transitions for the compounds, Tg of the ionomer and a Tm for the NaOl. These are clearly seen in the tan data for the Na- SPEEK composite in Figure 4.4, though there is overlap between the two loss dispersions. Both represent a possible Tc and, thus, dual and triple shape memory behaviors are conceivable. Two peaks were not resolved in the tan data for the Zn- SPEEK composite, but the single loss dispersion seen in Figure 4.4 is much broader than the Tg peak in either the neat ionomers for the Na-SPEEK composite. The similarity of the Tg of Zn-SPEEK and Tm of NaOl, 253°C and 259°C, respectively (see Table 4.1) suggests that the broad dispersion in Figure 4.4 is composed of the two transitions. As discussed in Chapter III, the advantage of developing a shape memory polymer from PEEK is the thermoplastic nature of the polymer. PEEK can be melt processed at temperatures beyond its melting points. However, the M-SPEEK ionomers are amorphous and possess high Tg. The linear viscoelsticity data shows that the materials are still solid like above their Tg, and high stresses are needed to process the material using conventional polymer processing operations. This situation is improved for M- 97 SPEEK/NaOl compounds, which can be processing above Tm of the NaOl. As shown in Figure 4.4, NaOl not only decreased Tg of M-SPEEK and also act as a plasticizer for the ionomer, reducing glassy modulus of M-SPEEK. 4.3.4 Shape memory behavior Although the neat PEEK exhibited some shape memory behavior, only 28% of the deformation was fixed by the temporary network (i.e., the glassy amorphous phase) and the shape recovery efficiency was only 35%. H-SPEEK showed slightly improved shape memory behavior by fixing 54% strain and recovering 49% deformation. In H- SPEEK, the permanent network was provided by hydrogen bonding of the sulfonic acid groups and glass transition was still the switching phase. The hydrogen bonding exhibited excessive creep under load, which led to poor shape memory performance of H-SPEEK. However, the ionic interactions in M-SPEEK form a much more robust permanent network. Both shape fixing and recovery were significantly improved. Figure 4.5 shows four consecutive shape memory cycles for Na-SPEEK, and the shape recovery and fixing efficiencies are summarized in Table 4.3. The sample film was heated to 270°C, stretched to about 6% strain and cooled under constant stress to 30°C. After allowing the strain to equilibrate at 30°C, temperature, the stress was removed. Shape recovery was achieved by reheating the film to 270°C, which was greater than Tc = Tg = 250°C. The shape fixing was significantly improved by Na-SPEEK, but the shape recovery was still poor, see Table 4.3. The difference in R between the first and subsequent cycles was due to relaxation during the first shape memory cycle of residual stresses from the processing history of the sample, which is commonly observed with other thermoplastic shape memory polymers [163]. 98 7 6 5 4 4 E 3 3 2 Strain(%) 2 1 400 1 300 S 0 200 0 50 100 100 150 200 Temperature( 250 0 Stress (kPa) C) 300 Figure 4.5 Four consecutive shape memory cycles for Na-SPEEK. The samples were stretched at 270°C with a stress of 0.32 MPa. (Tc = 250°C). The numbers denote the cycle number. Table 4.3 Shape fixing and recovery efficiencies of M-SPEEKs Na-SPEEK Zn-SPEEK Cycle F (%) R (%) F (%) R(%) 1 79 32 88 63 2 78 46 89 92 3 78 44 89 98 4 79 44 88 99 Significant improvements for shape fixing and shape recovery were achieved for M-SEEK ionomers by changing the metal cation to zinc, see Figure 4.6 and Table 4.3. After the first shape memory cycle, the subsequent shape memory cycles were reproducible and the recovery efficiency was nearly perfect. Another improvement is the elasticity of Zn-SPEEK at high temperatures. As shown in Figure 4.6, Zn-SPEEK can be 99 deformed at higher strain than Na-SPEEK. The temporary network, which is stabilized by the glassy state did, however, still exhibited some creep relaxation and the fixing efficiency was only ~90%. The difference between the behavior of the two M-SPEEK salts is probably due to the fundamental difference in the intermolecular association of the salt groups that comprise the ionic nanodomain crosslinks and the larger Coulomb energy of the zinc sulfonate ion-pair (~2.5 times that of sodium sulfonate). The Zn2+ cation can form a salt bridge, whereas the Na+ cation is associated with only one sulfonate group and the crosslink is due to dipole-dipole interactions of the ion-pairs. Tc 14 12 10 8 6 Strain(%) 4 2 1.2 1.0 0 0.8 0.6 -2 0.4 50 100 0.2 150 0.0 Stress (MPa) Temperature200 ( 250 C) 300 Figure 4.6 Four consecutive shape memory cycles for Zn-SPEEK. The samples were stretched at 270°C with a stress of 1 MPa. (Tc = 250°C). The numbers denote the cycle number. The problem of creep recovery of the temporary network for the neat M-SPEEK ionomers was resolved by adding a high melting point fatty acid compound, NaOl. This approach to prepare a shape memory polymer is identical to that described in refs. (1 and 100 2) except that in the present work a high temperature ionomer and a high melting point fatty acid salt were used. M-SPEEK was amorphous material and exhibited a physically crosslinked structure. M-SPEEK/NaOl(30) was crystalline material, in which some of NaOl still existed as crystals and some dissolved in M-SPEEK matrix, which can be concluded from DSC and WAXD data. Figure 4.7 shows the molecular structure of NaOl and the structure change of M-SPEEK after blending with NaOl. Ionic groups Figure 4.7 Illustration of molecular change after M-SPEEK blending with NaOl Figure 4.8 shows four consecutive shape memory cycles for Zn-SPEEK/NaOl(30), and the fixing and recovery efficiencies for Zn-SPEEK/NaOl(30) and Na- SPEEK/NaOl(30) are summarized in Table 4.4. The composite film was deformed above the Tm of the NaOl in the compound (250°C) and the Tg of the ionomer in the compound (208°C). The network formed by the ionic nanodomains served as the “permanent” network and the temporary network was formed by a combination of the glassy 101 amorphous phase of the ionomer and the crystalline NaOl. The crystallinity of the NaOl in the compound was much less than in the neat as-received NaOl, which for comparison purposes was assumed to be 100% crystalline. For the compound made with Zn-SPEEK, only 47% of the NaOl crystallized and for the Na-SPEEK only 36% crystallized. The reason for the low crystallization of the fatty acid salt is most likely due to partial miscibility of the NaOl with the ionomer, which is also responsible for the 40-50°C decrease of Tg of the ionomer in the compound, see Table 4.1. 1 2 100 3 4 80 60 40 Strain (%) Strain 2.0 20 1.6 1.2 0 0.8 50 100 0.4 150 Stress (MPa) 200 Temperature ( 250 0.0 300 oC) Figure 4.8 Four consecutive shape memory cycles for a Zn-SPEEK/NaOl composite film. The samples were stretched at 270°C with a stress of 1.5MPa. The numbers denote the cycle number. Except for the initial shape memory cycle, where the recovery efficiency of both compounds was relatively low, ~50-60%, the addition of the NaOl to the ionomers produced shape memory materials with excellent fixing and recovery efficiencies, see Table 4.4. Both approached 100% for F and R. The improvement of the shape memory 102 behavior by the addition of the fatty acid salt is attributed to the strong interactions between the metal sulfonate groups in the ionomer and the metal carboxylate groups in the NaOl, which has been demonstrated previously with another system to produce a robust temporary network. 80 1 3 60 2 4 40 Strain (%) Strain 0.6 20 0.5 0.4 0 0.3 0.2 50 100 150 0.1 Temperature 200( 250 0.0 Stress (MPa) 300 C) Figure 4.9 Four consecutive shape memory cycles for a Na-SPEEK/NaOl composite film. The samples were stretched at 270°C with a stress of 0.56MPa. The numbers denote the cycle number. Table 4.4 Shape memory properties for M-SPEEK/NaOl(30) compounds Na-SPEEK/NaOl(30) Zn-SPEEK/NaOl(30) Cycle Tc (°C) F (%) R(%) Tc (°C) F(%) R(%) 1 215 97 62 228 97 53 2 231 97 90 240 96 92 3 233 97 96 235 96 100 4 230 97 100 237 96 100 Figure 4.9 shows four continuous dual shape memory cycle for Na- SPEEK/NaOl(30). Although there were two separate temporary networks in the compounds, Figure 4.8 and 4.9 show that shape recovery began for both mateirla above 103 the Tg of the ionomer and below the Tm of the NaOl. For these systems, Tc was determined from the recovery step of the shape memory cycle, and it was defined as the temperature when the strain of the unloaded sample (temporary shape) started to change (in this case, decrease). Tc increased about 15°C after the first shape memory cycle for Zn-SPEEK/NaOl (30), but still was about 20°C lower than Tm of the NaOl in the compound. Na-SPEEK/NaOl(30) started to recover at slightly lower temperature than Zn-SPEEK/NaOl(30), which is probably due to the relatively lower Tg of Na-SPEEK in the compound. For the fixing step, both materials had to be heated above Tm of the NaOl to stretch the material. In addition, the compounds exhibited excellent elasticity at high temperature. NaOl act as plasticizer to the ionomer melt and enable the material achieve large deformation as temporary shape (>60%). This is another advantage of developing the shape memory polymer from compound, where the low molar mass crystalline component providing a secondary network and plasticizing the thermoplastic polymer matrix. 4.3.5 Triple shape memory behavior The previous section described the dual shape memory behavior of the ionomers and ionomer/NaOl compounds. Since there are two reversible physical networks, the NaOl crystals and the ionomer glass, in the compounds, there is the possibility for triple shape memory using the Tm of the NaOl and the Tg of the ionomer matrix to achieve two distinct switching temperatures, Tc1 and Tc2. Figure 4.10 shows a triple shape memory cycle program where Tg of M-SPEEK is Tc1 and Tm of NaOl is Tc2, and Figure 4.11 shows three consecutive triple shape memory cycles for Zn-SPEEK/NaOl(30). The Tg of Zn-SPEEK in the compound, 208°C, and Tm 104 of the NaOl crystals in the compound, 250°C, provides a 42°C temperature window for a second temporary shape. Figure 4.10 Schematic illustration of a triple shape memory cycle programming, where Tg is from M-SPEEK and Tm is from NaOl The Zn-SPEEK/NaOl(30) film was first heated to 280°C, which is above both transition temperatures, and the film was deformed to 57% strain. The sample was then cooled to 240°C while keeping the applied stress constant, and the sample was then held at 240°C to allow the NaOl to crystallize and form a temporary network (temporary shape B in Figure 4.10). The external stress was then removed to fix the temporary shape B, which corresponds to εB in Figure 4.11. There was about a 6% strain contraction after the stress was removed. A similar strain contraction was also observed in other triple shape memory polymer systems [164-165]. A second tensile deformation of 79% was then applied to the temporary shape B at 240°C. This was possible, since the amorphous ionomer phase was still above its Tg at that temperature. The temperature was then reduced to room temperature while holding the external stress constant. During this cooling step, the amorphous ionomer phase 105 vitrified as the temperature passed through its Tg, and a second temporary shape (temporary shape C in Figure 4.10), corresponding to C in Figure 4.11 was fixed by removing the external stress. Note that A in Figure 4.11 denotes the original (permanent) shape of the sample. 1st cycle 2nd cycle 3rd cycle 300 4 100 Temperature C 250 80 3 B C) 200 60 2 150 40 1 Strain Strain (%) 100 Stress (MPa) 20 Temperature ( Temperature rec 50 0 0 A A/rec 0 -20 0 100 200 300 400 500 600 Time (min) Figure 4.11 Consecutive tripe shape memory cycles for Zn-SPEEK/NaOl(30%) Sequential shape recovery was achieved by heating the sample that was in the temporary shape C to 240°C to recover the temporary shape B, and then heating to 280°C to recover the permanent shape A, see illustration in 4.10. Shape fixing (Rf) and recovery (Rr) can be calculated according to the following equations [164]: xy (13) R yx 100% ,loady x 106 ,recxy (14) F xy 100% xy where x and y denote two different shapes, εy,load is the maximum strain after applying load, εy and εx are fixed strains after unloading, and εx,rec is the strain after recovery. Table 4.5 Shape fixing and recovery efficiencies three consecutive triple shape memory cycles for Zn-SPEEK/NaOl(30). Cycle Zn-SPEEK/NaOl(30) Tc1 (°C) Tc2 (°C) F1 (%) F2 (%) R1(%) R2(%) 1 215 255 91 95 84 60 2 217 256 90 90 88 99 3 214 253 89 88 93 99 The shape memory characteristics for three consecutive triple shape memory cycles for Za-SPEEK/NaOl(30) and Na-SPEEK/NaOl(30) are summarized in Table 4.5 and 6, respectively. Figure 4.12 shows three continuous tripe shape memory behavior for Na-SPEEK/NaOl (30) compound following the same procedure as described in Figure 4.10. The linear viscoelastic data of Na-SPEEK/NaOl (30) shows two obvious mechanical transitions, which indicates the promising tripe shape memory behavior (Table 4.6). As with the dual shape memory behavior, the tripe shape memory cycles exhibit some permanent strain during the first cycle. For both M-SPEEK/NaOl systems, the shape recovery from temporary shape C to B was not as efficient as that from B to A. For the first recovery step, Tg,M-SPEEK < T 1st cycle 2nd cycle 3rd cycle 300 3.5 70 Temperature C 3.0 60 250 2.5 B 50 C) 200 2.0 40 150 1.5 30 1.0 Strain 20 (%) Strain 100 (MPa) Stress Temperature ( rec 0.5 10 50 0.0 0 rec 0 -10 0 100 200 300 400 500 600 Time (min) Figure 4.12 Consecutive tripe shape memory cycles for Na-SPEEK/NaOl(30%) Table 4.6 Shape fixing and recovery efficiencies three consecutive triple shape memory cycles for Na-SPEEK/NaOl(30). Na-SPEEK/NaOl(30)* Cycle Tc1 (°C) Tc2 (°C) F1 (%) F2 (%) R1(%) R2(%) 1 204 236 92 93 77 61 2 209 236 91 83 83 97 3 211 235 93 73 73 100 * 1 denotes temporary shape C and 2 denotes temporary shape B Tc1 represents the switching temperature for the recovery of the temporary shape C to temporary shape B (see Figure 4.10), where the glass transition of M-SPEEK serves as the temporary network. Tc2 is the switching temperature of temporary shape B recovering to permanent shape A, where the NaOl crystals form the temporary network. For Zn-SPEEK/NaOl(30), Tc1 (~215°C)was slightly higher than the Tg of Zn-SPEEK (208°C) and Tc2 was the melting temperature of the NaOl crystals (~255°C), which 108 provides a ~45°C temperature window for programming a second temporary shape. For Na-SPEEK/NaOl(30) system, the Tc1 (~208°C) was much higher than the Tg of Na- SPEEK (170°C), and Tc2 (236°C) was about 20°C lower than Tm of the NaOl. In this case, the temperature window for a second temporary shape was only 28°C. Varying the temperatures used for the triple shape memory experiments within the ranges between the transitions may affect the F and R values, but this was not investigated in this study. 4.4 Conclusions Dual and triple high temperature shape memory polymers can be prepared from sulfonated poly(ether ether ketones) ionomers and compounds of the ionomer and a high temperature fatty acid salt, sodium oleate (NaOl). Metal salts of sulfonated PEEK showed exhibited dual shape memory behavior, where the glass transition temperature of M- SPEEK served as the switching temperature and ionic nanodomains due to ionic or dipolar interactions between metal sulfonate groups provided the permanent network. The temporary network, which was formed by the glassy state of M-SPEEK, had fixing efficiencies of 80-90% and recovery efficiencies as high as 99%. More robust shape memory polymers were achieved using compounds of M- SPEEK and 30 wt% NaOl. In that case, two temporary networks were formed, one from the glassy state of the ionomer and the other from the NaOl crystals. Dual shape memory compounds had a switching temperature of 230-240°C and fixing and recovery efficiencies close to 100%. The M-SPEEK/NaOl compounds also exhibited Triple shape memory using the two separate switching temperatures of the two temporary networks, Tg of the ionomer and Tm of the NaOl. The shape memory effectiveness of the Zn- SPEEK was better than for the Na-SPEEK. The Zn-SPEEK/NaOl compounds had a 109 ~50°C between the two switching temperatures and shape fixing and recovery efficiencies for the two temporary shapes of ~90% and shape recovery of the temporary and permanent shapes of ~90% and 100%, respectively. 110 CHAPTER V IONOMER MODIFIED ASPHALT 5.1 Asphalt properties Asphalt is a dark brown to black cement-like material containing bitumens as predominating components, which occur in nature or are obtained in petroleum processing. Bitumens are defined as the natural or manufactured black colored solid, semisolid or viscous cement-like materials that are mainly composed of high molecular weight hydrocarbons [166]. Asphalt is a highly complex and its properties vary depending on the source of the crude oil from which the asphalt originates and on modifications that induced by treatments in the refinery [167]. 5.1.1 Chemical composition of asphalt Generally, asphalt contains saturated and unsaturated aliphatic and aromatic compounds with up to 150 carbon atoms, and the molecular weight of these constituent compounds range from several hundred to many thousands [168]. The compounds contain about 80wt% carbon, around 10wt% hydrogen, up to 6wt% sulfur, small amount of oxygen, nitrogen, and trace amounts of metals such as iron, nickel, and vanadium. These compounds are further classified as asphaltenes and maltenes according to their polarity [169-170]. Asphaltenes are aggregates of polar aromatic compounds and are not soluble in hexane or heptanes. Asphaltenes are regarded as being formed by the 111 condensation of resins. They are black colored and compounds with high molecular weight ranging between 1200 and 200,000g/mol. Asphaltenes affect strength and stiffness of asphalt. Maltenes are nonpolar aliphatic hydrocarbons and are soluble in hexane or heptanes. Maltenes are considered containing two fractions: heavy oils and resins. The heavy oils are the liquid part of the asphalt and consist of normal-, iso-, and cyclo-paraffins and condensed naphthenes with some alky aromatics. Maltenes have a key feature of dispersing polar asphaltenes and is responsible for viscosity and fluidity of the asphalt. The resins are semi-liquid and solid materials at room temperature. They are chemically very similar to asphaltenes but they are not as polar as asphaltenes and the molecular weight is much lower ranging from 300 to 2000g/mol. Normally, maltenes are the major components in asphalt, leaving the asphaltenes content to be only about 5~25wt% [171]. At the intermolecular level, polar molecule is able to attract another one as a result of forming dipoles. These associations lead to a three-dimensional intermolecular structure. This structure is held together by electrostatic and other short range forces [172]. These interactions are weaker than covalent bonds between the hydrocarbons and also determine the following behavioral characteristics of asphalt: (1) Elasticity through the effects of the polar molecule network. (2) Viscosity because the various parts of the polar molecule network can move relative to one another due to their dispersion in the fluid non-polar molecules. The polar asphaltenes also contribute to the adhesion between the polar surfaces of aggregates and asphalt. 112 A chemical model was built and refined by Strategic Highway Research Program (SHRP) to describe the asphalt structure. This model demonstrates asphalt microstructure as a continuous three-dimensional association of polar asphaltenes dispersed in a fluid of non-polar maltenes [173]. 5.1.2 Physical properties of asphalt The physical properties of asphalts vary due to their crude oil sources and the operations involved in the production. These properties are characterized using various standard tests, such as ductility, penetration, softening point, asphaltene content, temperature susceptibility, cold fragility, and viscosity [174]. The ductility of asphalt shows its ability to elongate or stretch without breaking under a standard testing condition [175]. The ductility test (ASTM D113) is conducted by pulling asphalt apart at a uniform rate at a specified temperature. The elongation before break is then measured in centimeters. Asphalt with a very low ductility is generally considered to have poor adhesive properties Penetration of asphalt is the oldest and most commonly used test to determine the consistency (hardness) of asphalt at a specified test condition [176]. Penetration test is performed with a penetrometer where a standard needle is applied to the sample. The needle penetrates vertically into the sample under fixed conditions of temperature, load, and time (ASTM D5). Usually, soft asphalt has deep penetration. The penetration test is used to determine grades of asphalt cement and to measure changes in hardness due to ageing hardening or changes in temperature. The softening point of asphalt is defined as the temperature at which a disk of the sample held within a horizontal ring is forced downward a distance of 2.54cm under the 113 weight of a 3.5g steel ball (ASTM D36) [177]. As the temperature increases, the asphalt sample gradually and imperceptibly changes from a brittle, slow flowing material to a softer, less viscous liquid. At this condition, the asphalt sample can no longer support the weight of the steel ball. The flash point of asphalt defines the temperature to which asphalt can be safely heated in the presence of an open flame [178]. This property is measured by heating the sample in an open cup at a specified rate. A small flame passing over the surface of the cup can cause the vapor from asphalt to temporarily ignite or flash at a certain temperature. This temperature is considered to be the flash point (ASTM D92). This test can also detect contaminating materials in asphalt by which the flash point suddenly drops. In addition, the flash point test is related to the hardening potential of asphalt. Asphalt with a high flash point will have a lower hardening potential. The viscosity test is an important physical property to evaluate the consistency of asphalt [179]. For asphalt, viscosity is usually measured under shear and is defined as the ratio between the applied shear stress and induced shear rate. Asphalt mostly behaves like a non-Newtonian fluid, especially at lower temperatures and the relationship between shear stress and shear rate in non-linear. The most common viscosity test is the absolute viscosity test by vacuum capillary viscometer (ASTM D2171), where the standard test temperature is 60°C. 5.2 Pavement Problems Asphalt and other bituminous products have been used for road construction and roofing materials for thousands of years due to their remarkable waterproofing and binding properties [180]. They can be obtained from nature’s bituminous deposits but 114 currently most asphalt is derived from crude oil. Asphalt for road surface use is produced by vacuum distillation, whereas for roofing materials, it is produced by air blowing. In real life, the performance of road pavement deteriorates due to the failure of asphalt binders. Asphalt behaves like a thermoplastic liquid, which is an elastic solid at low temperature or rapid loading and a viscous liquid at high temperature or slow loading. These behaviors lead to the performance failure such as thermal cracking, rutting, and fatigue cracking. Changes in temperatures and traffic loadings during the life of the asphalt pavement requires modification to improve the performance of asphalt binder to minimize the stress cracking that occurs at low temperatures and the plastic deformation at high temperatures. Thermal cracking of asphalt pavement often occurs in the cold climate regions such northern United States and Canada. This type of distress manifests itself as a series of transverse cracks that extend across the pavement surface as shown in Figure 5.1 (a). It is a consequence of low temperature shrinkage exceeding the ability of stress dissipation by the asphalt binder [181]. Pavement rutting is the accumulation of permanent deformation in a pavement structure leading to a distorted pavement surface as shown in Figure 5.1 (b). The progression of rutting is cracking and eventually complete disintegration [182]. It usually happens in regions with warm climates and heavy traffic loads. Rutting is induced by reduction in binder viscosity and permanent deformations [183]. Fatigue cracking in asphalt pavement is due to tensile stresses developed from the flexure of the layer and propagates to the pavement surface under repeated traffic loads [184]. The maintenance cost for pavement cracking is high, e.g $1-2/meter [185]. The accumulation of cracking is also a sign of significantly shortened residual pavement life 115 and eventually a much rougher ride. Therefore, it is necessary to modify asphalts in order to improve their mechanical properties and reduce the expense of maintenance. (a) (b) (c) Figure 5.1 (a)Low temperature cracking, (b)rutting and (c)fatigue cracking of pavements. (Reprinted with permission from [186]) 5.3 Polymer Modified Asphalt (PMA) Polymers have been widely used to modify asphalt binders to improve the mechanical properties due to their ability to highly enhance asphalt binder properties in both low and high temperature ranges [187-188]. The addition of small quantity polymers is able to significantly improve pavement performance and durability and therefore reduce the maintenance costs. Polymer modification enhances the asphalt binder’s resistance to rutting and thermal cracking, decreased fatigue, cracking, stripping, and less 116 temperature susceptibility. Usually, about 3 to 7wt% of polymer with respect to the asphalt phase are blended with asphalt. In addition, polymer modification also improves the pavement’s resistance to oxidation and aging. Polymers that can be used as the modifier have to meet the following requirements: (1) Polymers must be compatible with asphalt binders [189] (2) The high temperature viscosity of binder should not change too much after polymer addition so that the cost of road building process and apparatus would not increase. The only change for the production line should be the addition of asphalt-polymer mixing apparatus. [190] (3) The cost of the polymer must be balanced by both improved asphalt properties and the reduction in the costs of road construction and maintenance. [191-193] The addition of polymers in asphalt binders changed the original colloidal equilibrium of asphalt because the polymer interacted preferably with maltenes. It is known that materials tend to be miscible with other materials with a similar polarity. Most polymers used to modify asphalt have the similar solubility parameters with maltenes [194]. The interaction between polymers and maltenes lead to a local solvent extraction. Thus, the mixture will form a new three dimensional network, where a polymer-rich phase and an asphaltene-rich phase coexist in a metastable equilibrium. However, this mixture is thermodynamically unstable and tends to phase separate. If the polymer-rich phase segregates in the storage tank, the desired properties of asphalt binders are lost. Therefore, in real-life manufacture, high temperature storage is necessary at least during the transportation to the application sites. 117 The properties of asphalt binders are both time and temperature dependable, which is particularly true for rheological properties of asphalt binders. The desired condition is that the polymer maintains its basic morphological structure and physical characteristics due to its partial solubility and compatibility upon mixing. In this case, the polymer-rich phase is swelled by the oily aromatic part of asphalt and the asphaltenes rich phase is dispersed. The rheological properties of polymer modified asphalt will resemble those of polymers. Therefore, the small amount of polymer addition has significant effects on the overall asphalt binder properties. 5.3.1 Thermoplastic polymer modification Thermoplastic block copolymers are the most widely used polymer material for asphalt modification. Block copolymers are joined by two or more chemically distinct polymer blocks in a linear manner. Usually, the two blocks are not compatible and tend to form a phase segregated network, which is a similar situation as asphalt binders. Modification using block copolymers is achieved by simple mechanical dispersion of the polymers in molten asphalt under high shear. Styrene butadiene styrene (SBS) modifiers have become increasingly popular because of their apparent success in reducing the migration of cracks in pavement [195-199]. The polystyrene (PS) blocks are on both ends of the polymer with a glass transition temperature around 100°C, and the central segment is polybutadiene (PB) whose glass transition occurs around -80°C. These two blocks are not compatible and are phase separated to form a network with glassy PS domains dispersed in PB flexible matrix. Since the paving operation temperature is normally between these two glass transition temperatures, this tri-block copolymer exhibits a crosslinked elastomer network behavior under this condition. This block copolymer is 118 also melt-processable above PS glass transition temperature. The interaction between SBS and asphalt occurs by physical rather than chemical bonding, so the crosslinking system is reversible even after many heating and cooling cycles. Upon mixing, PB blocks absorb the maltenes and swell to a much larger volume than its initial state. The appropriate amount of polymer will be able to swell and form a continuous phase. After cooling down to room temperature, PS blocks segregated to form the rigid domains that are dispersed in an olefinic phase swollen by maltenes, as shown in Figure 5.2 [200]. Consequently, the physical network of SBS is restored and the elastic properties will transfer to the whole mixture. SBS modified asphalt binders are anticipated to have enhanced strength and elastic properties in a wide range of temperatures. The Florida Department of Transportation investigated the effect of SBS modification on cracking resistance and healing characteristics of the pavement. The results showed that SBS modification did improve pavement resistance to cracking and in a way of reducing rate micro-damage accumulation [201]. Figure 5.2 Schematic illustration of a network in SBS modified asphalt. (Reprinted with permission from [200]) 119 However, asphalt binders are composed of many components that have different solubility parameters, indicating that there are different degrees of miscibility with the blocks of SBS. Chen et al. [202] studied the morphology of SBS modified asphalt binder taken from the job site using transmission electron microscopy. The results showed that the polymer phase was not homogeneously distributed in the asphalt binder. The morphology varies depending on the source of asphalt and structure of polymer. There can be a continuous asphalt phase with dispersed SBS polymers or a continuous polymer phase with dispersed asphalt phase, and even two interlocked continuous phases. The formation of a network structure is very critical to the rheological properties of modified asphalt binder and therefore affects the pavement’s performance. Studies also show that there is a decrease in strength and resistance to penetration of SBS modified asphalt binder at high temperatures, although low temperature flexibility is improved [203]. This is due to the unsaturated nature of double bond in PB blocks. The double bond contributes to the low resistance to heat, oxidation and atmospheric agents, which leads to the limited duration and recycling chances of end-of-life road pavements. Research also shows that SBS tends to degrade even during both production and treatment of modified asphalt because its double bonds are subjected to free-radical reaction upon exposure to UV light or during thermal and mechanical processes [204- 205]. Efforts had been made to saturate the polymer. Therefore, styrene-(ethylene-r- butylene)-styrene (SEBS) triblock copolymer has been obtained by hydrogenation of SBS. The saturation of double bond made the polymer considerably less polar than SBS. EB blocks are less compatible with asphalt binder than polybutadienic ones, which increases the risk of phase separation and storage instability [206-207]. SEBS modified 120 asphalt solves the problem of asphalt being less resistant to degradation and oxidation, meanwhile the elastic properties are reduced as compared to those of modified SBS [208- 209]. Other approaches have been investigated to improve the storage stability of SBS modified asphalts such as the addition of a small amount of sulfur. Sulfur can vulcanize polymer molecules and chemically connect polymer and asphalt through sulfide and polysulfide bonds [210]. Compatibilizer agents or aromatic oils can be added to stabilize the mixture, which can still destroy the benefits of the SBS copolymer, thus resulting in a loss of polymer modified asphalt binder properties [211-212]. After all, the major disadvantage of SBS and SEBS modification is their high cost. The production of polymer highly increases the cost of paving. 5.3.2 Polyolefin modification Thermoplastic elastomers SBS and SEBS improve the elastic properties and reduce the permanent deformation of asphalt modified binder. However, the high cost and low resistance to degradation limits their application in pavement industry. Polyolefins were used to modify asphalt binders, which have advantages over thermoplastic elastomers [213]: (1) Polyolefins are a low cost material with a wider range of characteristics and technical grades (HDPE, LDPE, LLDPE), and are also collected for recycling. The global production of polyethylene (PE) in 1984 excessed 20 million tons along with huge quantities of reclaim PE at price not much different than that of asphalt binder itself. The recycling of PE can also provide extra source of supply. [214] 121 (2) Polyolefins molecular chains are saturated and have high resistance to degradation and oxidation. (3) Polyolefins provide asphalt binders high rigidity and resistance to the deformation under traffic loading. Polyethylene (PE), polypropylene (PP), and their copolymers have been used to modify asphalt binders. The mixture can be prepared at elevated temperatures under high shear conditions, at which the PE particles gradually absorb the aliphatic maltene and partially dissolve to form a highly viscous, elastic dispersion. PE is a crystalline material below its melting temperature and assumes the characteristics of a filler at the asphalt binder operation temperature. PE modification increases asphalt binder resistance to deformation at elevated temperature and impact fracture toughness at lower temperatures. It is confirmed that PE modified asphalt can last 2.7~2.9 times longer than an unmodified paving formulation during the past ten years in Austria, the United Kingdom, and Italy [215]. Fawcett and coworkers [216] studied the modification of asphalt binders using various polyolefins such as polypropylenes (PP) with different architectures and polyolefin copolymers. Semicrystalline atactic polypropylene (APP) and commercially available isotactic polypropylene (IPP) were blended with asphalt binders. The rheological properties results showed that the stiffness of asphalt was enhanced at low temperature by 100 fold and at high temperatures to an even greater extent. The high modulus resulted from the PP crystals. The copolymer showed similar high modulus, which suggests that the polymer-rich phase in asphalt binder is extensive. However, PE and PP are non polar materials and are almost completely immiscible with asphalt no matter the type considered. In addition, these materials have 122 high degrees of crystallization, which can further limit the interaction between asphalt and polymer [217]. Therefore, storage stability is still the main issue that delays their industrial application. Polyolefin modified asphalt binders are mostly used for the production of impermeable membrane for roofing or other waterproofing use, which can be stored at room temperature. Polyolefins have been modified to increase the compatibility with asphalt, such as the insertion of polar groups on to polymer chain. An example is ethylene-vinylacetate (EVA), which contains an ester group. This approach not only enhanced the polarity of the polymer chain, but also reduced its crystallinity due to the introduction of structural irregulatities [218-219]. Moreover, EVA has the similar ability to form a physical network as SBS thermoplastic elastomers. For instance, the crystals in EVA form the rigid parts, which are similar as polystyrene hard domains and the rest polymer chains are swelled by the asphalt components. The degree of crystallinity of EVA can not be too high so that the formation of large rigid domains restrict swelling by the asphalt components, but it can not be too low in order to have enough polarity to be compatible with asphalt components. Polyolefins blending with other materials were developed to increase the compatibility with asphalt binders. Ouyang and coworkers [220] reported LDPE compounded with silica to improve high temperature storage stability of asphalt binder. The silica used is a nonfiller with very fine particles. The modification is effective only when the ratio of LDPE/silica is around 100/60 (w/w) and the silica content in the whole asphalt binder is less than 3.2%. The incorporation of silica reduced the density difference between the LDPE/silica compound and the asphalt and equalized the polarity difference between asphalt and LDPE. Therefore, the high temperature stability is 123 improved while the mechanical properties are reserved as LDPE modified asphalt. Pérez- Lepe and coworkers [221] studied the asphalt modification using a blend of HDPE and ethylene-propylenediene monomer (EPDM). The addition of HDPE to asphalt binders can enhance their temperature susceptibility and resistance to rutting. EPDM acts as an emulsion stabilizer to improve the compatibility of HDPE and asphalt binder. EPDM can also serve as an additive to lower the rigidity of HDPE at lower temperatures. 5.3.3 Reactive polymers modification Reactive polymers containing both reactive and non-reactive polar groups were developed to increase the compatibility of olefinic polymers with asphalt binders. They are usually ethylene based copolymers containing epoxy rings and an ester group such as methyl, ethyl, or butyl acrylate. Ethylene is the main component so they are usually referred as reactive ethylene terpolymers (RET) [222]. The acrylic functional groups in RET improve the polarity of asphalt binder, meanwhile the epoxy rings react with the asphalt’s carboxylic groups. RET are commercially available on the market. For example, Dupont produces Elvaloy® , which is ethylene glycidyl acrylate (EGA) terpolymer. Elvaloy have been used for road pavement since 1991 [223]. However, RET are only used as polymer compatibilizers by a small amount of addition because of their shortcomings. The inter-chain reactions between polymer chains are highly possible, leading to chemically crosslinking of the material rather than physical interactions. The chemically crosslinks hinder the swelling of polymer with maltene and can not form the asphaltene dispersed in the swollen polymer-rich phase three dimensional network structure. Therefore, there is an upper limit to the amount of RET that can be used. Although RET are highly compatible with asphalt due to the chemical bonds formed with 124 asphaltene molecules, the use of RET is limited. Witzeak and coworkers [224] studied the laboratory performace of the commercial product Elvaloy modified asphalt binder. The results showed that the mixtures increased asphalt binder resistance to moisture damages and permanent deformations. However, Babcook et al [225] at the Dupont institute studied the modified binder properties at high temperatures. The results indicated that the binder failure occurred at a temperature above 6°C due to the loss of adhesive. Another type of reactive polymer, such as a maleated polypropylene, was investigated by Yeh et al. [226]. The functional group in this polymer positively improved the compatibility between the asphalt and polymer. Nevertheless, the application of reactive polymer in asphalt modification is limited because they are intrinsically expensive materials. Nevertheless, the addition of a polymer causes a significant increase in the production costs in terms of operative complications and storage. The polymer itself is expensive if copolymers, e.g. SBS are used. The low compatibility between asphalt and polymer can lead to severe phase separation when the binder is stored at relative high temperature in the absence of stirring. This causes the polymer rich phase to migrate to the higher part of the storage tank while the asphalt rich phase segregates to the lower part due to their viscosity difference. Consequently, an inhomogeneous binder forms and is useless for paving. Therefore, a continuing effort still needs to be made to develop efficient polymer modified asphalt. 5.4 Research Objective The objective of this research is to modify paving-grade asphalt with ionomers, particularly partially neutralized ethylene-acrylic acid copolymers and terpolymer 125 ionomers to improve the performance of the pavement, especially for low temperature application. All the described failures (cracking, rutting and fatigue) are directly related to the viscoelastic behavior of the binder. By studying the rheological properties of modified asphalt binders, their response to traffic loads, environmental conditions can be evaluated. The internal structure of asphalt binders plays a significant role in viscoelastic performance due to the complex morphology of asphalt itself. The ionomer modified asphalts were characterized by optical microscope and the thermal transitions were studied by differential scanning calorimetry. The strategic highway research program (SHRP) has a unique set of tests to evaluate grade of pavement involving rheological methods, such as dynamic shear and bending beam shear. The performance of ionomer modified asphalt was graded using these methods. 5.5 Superpave Test Superpave is an acronym for Superior Performing Asphalt Pavements. Superpave is a comprehensive asphalt mix design and asphalt binder specification system developed by the Strategic Highway Research Program (SHRP). $50 million of the SHRP research funds was invested to develop performance based asphalt binders and asphalt-aggregate mixes [227]. Asphalt binders used to be graded by either penetration or viscosity tests as shown in Figure 5.3. Both methods reflect the temperature effect on asphalt behavior, which can not fully characterize asphalt binder. However, SHRP researchers developed the new binder tests and specifications to address the asphalt contribution to three distresses (rutting, fatigue cracking, and low temperature cracking) found in asphalt concrete pavement. The Superpave binder tests measure physical properties, which directly describe how it will perform as a constituent in asphalt concrete pavement. The 126 binder test also provides information about the climate conditions and aging considerations. This is the most advanced asphalt binder grading system: Performance Grade (PG) [228] . Figure 5.3 Penetration grading and viscosity grading. (Reprinted with permission from [229]) The new Superpave tests characterize asphalt at a wide range of temperatures and at the period of time when the asphalt failure is most likely to occur, as shown in Figure 5.4. The typical set of Superpave asphalt binder tests include: (1) Rotational Viscometer (RV) test (2) Dynamic Shear Rheometer (DSR) test (3) Bending Beam Rheometer (BBR) test (4) Direct Tension (DTT) test (5) Rolling Thin Film Oven (RTFO) aging (6) Pressure Aging Vessel (PAV) aging 127 Figure 5.4 Schematic illustration of Superpave test types and temperature ranges. (Reprinted with permission from [229]). 5.5.1 Rotational Viscometer test Rotational viscosity is used to evaluate the workability of asphalt binders at high temperature. Asphalt viscosity at high temperature is measured to ensure that it sufficiently remains in a fluid form when pumping and mixing. According to the Superpave binder specification [230], the rotational viscosity should not exceed 3 Pa·S measured at 135°C. The asphalt binder is tested in an unconditioned or neat state, which is the state of material in the tank at the asphalt plant. Rotational viscosity is determined 128 by measuring the torque required to maintain the constant rotational speed of a spindle submerged in asphalt binder (Figure 5.5), as per the AASHTO TP48 standard test method [231]. Figure 5.5 Schematic illustration of rotational viscosity test. (Reprinted with permission from [228]) 5.5.2 Dynamic Shear Rheometer test Dynamic shear rheometer is used to characterize the viscoelastic behavior of asphalt binders at intermediate to high service temperatures. The test results provide an indication of the ability of the binder to resist permanent deformation at high service temperatures in the early stages of pavement life [232]. The DSR results also indicate the resistance of the binder to fatigue cracking at immediate service temperatures in the later stages of service life. DSR evaluates the behavior of an asphalt specimen when subjected to oscillatory (sinusoidal) stresses. A thin asphalt specimen is sandwiched between two parallel metal plates held in a constant temperature environment as per AASHTO TP5 standard test [233]. One plate is fixed while the other oscillates at an angular frequency of 10 radians per second for 10 cycles, as shown in Figure 5.6. The specimen’s response to the sinusoidal stress was measured and the complex shear modulus and phase angle was 129 calculated. The complex modulus, G*, represents the total deformation resistance under shear stress and the phase angle, δ, indicates the delayed strain response of the binder to the applied shear stress during steady state conditions. Figure 5.6 Schematic illustration of DSR measurement. (Reprinted with permission from [229]) Rutting parameter Rutting results from the deformation within the surface layer. A higher complex modulus and a smaller phase angle are desired for rutting resistance [234]. Higher complex modulus indicates that the asphalt binder is stiffer and provides increased resistance to deformation. A small phase angle means the binder has a greater elastic component and allows more of the total deformation to be recovered. Rutting is regarded as a stress-controlled, cyclic loading phenomenon. In each cycle of loading, the work done during deforming the asphalt binder at high temperature is partially recovered by the elastic rebound of the surface layer and partially dissipated through the permanent deformation as well as an associated generation of heat. The dissipated work is a function of stress and strain and can be expressed as [235]: (15) 130 Where σ0 is the applied stress, ε is the strain amplitude, δ is phase angle, and G* is shear modulus. The parameter G*/sinδ was chosen for SHRP specifications with respect to rutting. An increase in G* and a decrease in sinδ will both lower the amount of work dissipated per loading cycle. The Federal Highway Administration Asphalt Binder Expert Task Group (ETG) established the specifications for the DSR test parameters. The minimum limits of G*/sinδ for unaged asphalt is 1.0KPa, which is based on the tests on AC-10 viscosity graded asphalt. This asphalt provides reasonable service in moderate climates and yielded G*/sinδ value of approximately 1.0 KPa. The minimum limit of G*/sinδ for RTFO aged asphalt is 2.2KPa. [236-237] 5.5.3 Bending Beam Rheometer test The bending beam rheometer was developed to measure the stiffness of asphalt binder at low service temperatures and to evaluate the potential for thermal cracking [238]. According to the AASHTO TP1 standard test method, the asphalt beam sample is subjected to a creep load for 240 seconds at a constant low temperature. The load and deflection of the sample beam are measured and a plot is continuously generated. After 240 seconds, the load is automatically removed and the computer calculates the parameters as shown in Figure 5.7. The creep load is used to simulate the stresses that gradually build up in the pavement as temperature decreases. Two parameters can be obtained from this test: the creep stiffness and the creep rate (m-value). The creep stiffness is characterized by measuring the creep response of the beam sample at critical temperatures. 131 The m-value is defined as the slope of the log creep stiffness versus log loading time on a master curve [239]. The creep stiffness indicates how the asphalt sample resists the constant creep loading and the m-value is a measure of the rate at which the creep stiffness changes with loading time. The m-value also indicates asphalt’s ability to relax stresses. A low m-value indicates a slow rate of relaxation, where stresses build as the binder loses its ability to relieve the thermal-induced stresses by permanent flow. Figure 5.7 Schematic illustration of BBR test. Reprinted with permission from [229] The specification criterion for maximum creep stiffness at 60 seconds is 300MPa, which restricts the level of stresses developing in the pavement [239]. The specification criterion for minimum m-value at 60 seconds is 0.300, which maintains the rate of relaxation so that stresses can be relieved through permanent flow. 5.5.4 Rolling thin film oven aging (RTFO) The rolling thin film oven test (RTFO) simulates aging hardening of asphalt binder during the production and construction. RTFO is used to determine the effect of heat and air on a moving film of asphalt and to evaluate the resistance to aging [240]. The moving film was made by placing the asphalt binder sample in a jar and then placing in a circular metal carriage that rotates within the oven. According to the AASHTP T240 132 method, the RTFO test was conducted at 163°C with air blowing at 4000ml/min for 85minutes to achieve the accelerated age hardening, as shown in Figure 5.8. The rolling feature allows continuous exposure of asphalt binder to heat and air flow, maintains modifiers dispersed in the sample, and prevents the formation of a surface skin on the sample. The purpose of RTFO is to prepare aged asphalt binder for further testing and evaluation using the Superpave specifications. RTFO test measures the mass loss, which was vaporized by the testing procedure. A high mass loss value indicates that the material contains excessive volatiles and could age excessively. Mass loss can be calculated as: Mass loss,%=[(Original mass-Aged mass)/Original Mass]×100% (16) Figure 5.8 Rolling thin film oven (left) and RTFO samples (right). (Reprinted with permission from [229]) 5.5.5 Pressure Aging Vessel aging (PAV) The pressure aging Vessel is used to simulate the age hardening of asphalt during the first 5-10 years of pavement service life [241]. Similar as RTFO, PAV is also a conditioning procedure and its primary purpose is to prepare the aged binder materials for 133 further testing such as DSR and BBR to evaluate the binder’s performance following aging. According to AASHTO PP1 standard method, PAV expose the RTFO-aged asphalt residue to high air pressure (2.1MPa) and temperature (90°C, 100°C or 110°C depending on expected climatic conditions) for 20 hours to achieve accelerated hardening of the asphalt. Aging the asphalt under pressure can limit loss of volatiles and accelerate the oxidation process without heating the asphalt to extremely high temperatures. 5.6 Asphalt binder grade The asphalt performance grading specifications are based on the rheological characterization as discussed above. These tests can evaluate the pavement performance at three critical stages of service life: 1. The production and transportation of original asphalt binder before mixing 2. The short-term aging following the production and construction 3. The long-term aging during the pavement service life. Asphalt binder graded by this system should meet all specified criteria by this common set of tests. Asphalt grade is represented by two numbers, which indicate the temperature rating in degrees Celsius. For example, a PG 70-22 donates the average seven-day maximum pavement temperature of 70°C and the minimum pavement design temperature likely to be experienced is -22°C. These temperatures are pavement temperatures and not air temperatures and are usually 20°C higher than air temperature because the asphalt is dark color that absorbs heat and retains it. The average seven days maximum pavement temperature usually ranged from 46°C to 82°C and minimum is ranging from -46°C to -10°C. The reliability of this PG grading system is 98% and the detailed information can be found in SHRP reports [242]. Tests are run on the original 134 binder, where RTFO residue and PAV residue fully characterize the asphalt throughout its service life. RV is used to evaluate the pumpability at the asphalt plant. DSR test is designed to evaluate the binder properties at intermediate to high temperature related to the rutting or fatigue cracking. BBR and DTT are used to determine the binder properties at low temperatures related to thermal cracking. RTFO simulates the short-term aging during construction and PAV simulates long-term aging. 5.7 Innovation of this research Conventional polymer modifiers such as thermoplastic elastomers, polyolefins and reactive polymers are reported to improve asphalt performance to some extent. However, they all have several shortcomings that limit their application in asphalt industry. SBS is a high cost modifier and is easy to degrade because of the unsaturated blocks. Olefin polymers are not compatible with asphalt binder which leads to macro phase separation that embrittle the asphalt. Reactive polymers are expensive too and can only be used as compatilizer with a small amount. Ionomers are ion-containing polymers with a maximum ionic group content of about 15mol% [243]. The bulk properties of the polymers are governed by ionic interactions in discrete regions of the materials. It is another class of thermoplastic elastomers and has the property of nanophase-seperation. The polyolefin based ionomer would be a natural extension of the use of polyethylene modified asphalt. Besides the advantage of lower costs, it can provide good adhesion with asphalt, which is a crucial problem in modified asphalt binders. As an ion-containing polymer, the polar and non polar phase are incompatible and the electrostatic interactions will lead to nanophase seperation of about 1~5 nm aggregates (ionic clusters). At room temperature, the hard 135 ionic cluster domains that act as physical crosslink agents are dispersed in the hydrocarbon polymer matrix phase, providing the network structure that is need for polymer asphalt modification. Different from SBS thermoplastic elastomer, the domain ionic cluster formed is about 1 nm, which is smaller than SBS. In addition, the ionomer is amphiphilic and the polarity is improved so that the polar groups are expected to strongly interact with polar components in asphalt to improve adhesion. Meanwhile, the non-polar hydrocarbon component is likely to be swollen by maltenic molecules and then to form a three-dimensional network. This aspect of ionomer contrasts to SBS and polyolefin where the entire polymer is non polar. Nevertheless, there are few reports about ionomer-modified asphalts, and nothing commercial turns out. Cliplijauskaset al [244-245] synthesized asphalt ionomers, which was a chemical modification of a Shell Venezuelan asphalt by sulfonation or maleic anhydride addition. They concluded that the asphalt ionomers retained high cohesive strength under wet conditions and improve the low temperature flexibility. Yoshikane [246-247] described paving asphalt compositions that incorporated ionomers for improving strength. Gorbaty [248] patented sulfonated styrene-butadiene and SBS, sulfonated EPDM, and acrylic acid terpolymers modified asphalt with improved viscoelastic properties and elastic stablility. Dow Chemical Co. patented [249] using a carboxylate olefinic ionomer to modify paving asphalt. The ionomer is used to enhance adhesion of the binder. These two patents demonstrate the advantage of ionomer modification, but the information about the structure and properties of this modified asphalt binder were not discussed in detail and the role of the choice of cation on the adhesion properties was not well addressed. 136 Therefore, it is worthwhile to revisit the subject of asphalt modification using ionomers. Even though the previous researches provide limited detail, it is still encouraging that ionomer may be a good candidate material to improve the performance of asphalt pavement. There are several types of ionomers that could be used: random copolymers of ethylene, partially neutralized acrylic acid, lightly sulfonated polystyrene, sulfonated ethylene-propylene-dieneterpolymer(SEPDM) rubber, and sulfonated thermoplastic elastomers. Among them, the ethylene ionomers are commercially available with a range of acrylic acid concentration from 3-20% and it is easy to obtain sulfonated ionomers by modification of commercial polymers. Ionomers with various cations, including Na+, Li+, Zn2+ and Mg2+ are commercially available. The choice of backbone and ions allow one to control mechanical and viscoelastic properties and the choice of ions make it possible to adjust adhesion between ionomer and asphalt. The goal of this research is to modify paving grade asphalt binders by addition of ionomers, particularly partially neutralized copolymers of ethylene and metharcylic acid to improve the viscoelastic properties and measure the SuperPave properties of ionomer modified asphalt. 137 CHAPTER VI STRUCTURE AND PROPERTIES OF IONOMER MODIFIED ASPHALTS 6.1 Introduction This chapter presents the studies of structure and physical properties of ionomer modified asphalt. Unmodified asphalt suffering from several major distresses that shorten the pavement life and increases maintenance cost. Thermal cracking occurs when asphalt falls below a critical temperature during a thermal cycle. Fatigue cracking is the consequence of the mechanical stresses associated with repeated traffic loading. Polymers are widely used to modify asphalt by blending of asphalt and about 3-7 wt% polymers to improve pavement performance. Commercial polymer modified asphalts currently use polyolefins or styrenic block copolymers as the modifier. However, polyolefins are highly incompatible with asphalt and SB block copolymers are expensive. In this research, we proposed a new idea to use polyethylene based ionomers as modifier since ionomers have several promising advantages. First: ionomer used in this research is an ion –containing polymer with hydrophobic backbone and less than 6mol% ionic groups distributed through the polymer chain. Therefore, ionomers are amphiphilic with improved hydrophility by ionic groups. This feature improved compatibility comparing to polyolefin, because the ionic groups can strongly interact with the polar components in asphalt while the hydrophobic polymer matrix is likely to be swollen by the non polar component of asphalt. The adhesion between the ionomer and asphalt can be enhanced. 138 Second; as an ion-containing polymer, the polar and non polar phase are incompatible and the electrostatic interactions will lead to nanophase seperation of about 1~5 nm aggregates (ionic clusters) (Figure 6.1). At room temperature, the hard ionic clusters domains that act as physical crosslink agent are dispersed in the hydrocarbon polymer matrix phase. Different from styrene butadiene thermoplastic elastomer, the domain ionic cluster formed is about 1 nm, which is smaller than SBS. Figure 6.1 Schematic of ionomer nano-structure The blue nanophase constitutes ion-rich domains dispersed in a continueshydrophobic matrix. The key aspect of using polymer as a modifier is the ability of polymer to form a network structure, so as to achieve sufficient elasticity of asphalt/polymer mixture. The widely used polymer modifier is styrenic block copolymer thermoplastic elastomers (SBCs), e.g., Kraton® elastomers (typically SBS) [250]. SBCs are linear polymers that exhibit rubber-like properties up to ~80 - 100ºC by virtue of a physically-crosslinked network that results from nanophase-separation of ~10 nm diameter glassy polystyrene 139 nanodomains. The advantage of using SBS is the good interfacial adhesion between SBS and asphalt because of partial miscibility of asphalt in the polybutadiene phase [251]. The glassy polystyrene phase improves the toughness of the asphalt. It has been proved that the addition of 3-6% SBS to asphalt binder decrease the tensile modulus of the asphalt at low temperature and increase it at high temperatures [252]. Ionomer is also a nanophase- separated polymer, the physically crosslink of ionic clusters can provide the network structure that needed in polymer modified asphalt similar as SBS. Last; ionomers are low cost material and are commercially available in a wide range of compositions. In this research, a performance grade 64-28 asphalt and partially neutralized copolymers of ethylene and methacrylic acid ionomer were mixed at four concentration levels to yield ionomer modified asphalt blends. Thermal properties and morphology were investigated to study the miscibility of ionomer and asphalt. After establishing the linear viscoelastic range of response through strain sweep, frequency sweep tests at a temperature range of 30-80C were conducted to study the dynamic mechanic properties of the modified blends. The viscosity results were analyzed in terms of common models for viscoelastic fluid. Using the time-temperature superposition principle, the isothermal response curves were reduced to dynamic master curves of modulus and complex viscosity. The effects of ionomer concentration and mixing time on the visocoelastic behavior of ionomer modified asphalts were studied. The performance grade of ionomer modified asphalt was determined based on the results of a series of SuperPave tests, e.g. dynamic shear (DSR) and bending beam rheometer (BBR) tests. 140 6.2 Experiment details 6.2.1 Preparation of asphalt and ionomer blends Neat asphalt from vacuum distillation with Performance Grade 64-28, was used as base. The polymer modifier was DuPont’s Surlyn 9910, EMAA15Zn50, which was a zinc salt of PEMAA containing 15 wt% (~5.7 mol%) methacrylic acid that was 50% neutralized. The density of this material, as reported by DuPont, is 0.97 g/cm3 and the melt index (MI) is 0.7g/10min (190°C, 2.16 kgf). All the ionomer modified asphalt blends were prepared in a metal container equipped with a high shear mixer. The neat asphalt was added to the container and heated to 180°C, followed by the addition of a known quantity of ionomer. The mixture was continuously stirred for 45 and 60 minutes at the same temperature. Four sets of blends were prepared, which contains 3%, 5%, 7%, 9% of ionomer, respectively. 6.2.2 Materials Characterization Thermal analysis was performed by a TA instruments differential scanning calorimeter (DSC Q-100). The samples were sealed in aluminum sample pans using an empty aluminum sample with cap as a reference. The data were collected from -70°C to 150°C with a heating rate of 10°C /min. All the samples were evaluated by the thermal gravimetric analysis (TGA) on a TA instrument, model 2950 TGA HR V5.4A. The samples were heated from room temperature to 800°C with a rate of 10°C /min under nitrogen atmosphere. The morphology of mixes (phase structure and phase size) was studied by an optical microscope, model Nikon diaphot-tmd 24i inverted microscope. Images were taken with a 10× and 20× magnification. 141 Rheological measurements were conducted with a stain controlled Advanced Rheometic Expansion System (ARES) rheometer. Parallel plate geometry with a diameter of 25mm and gap of 1mm were used. All the mixtures were subjected to a dynamic frequency sweep test from 0.01 to 100 rad/s. All the tests were performed in the temperature range between 30 to 80°C. SuperPave test were conducted with 5% ionomer IMA and base asphalt using dynamic shear (DSR) and bending beam rheometer (BBR). All the samples were tested in accordance with AASHTO M 320. The basic BBR test uses a thin asphalt binder sample, which is sandwiched between two curricular plates. The lower plate is fixed while the upper plate oscillates back and forth across the sample at 10 rad/s to create shear. The DSR test measures the complex shear modulus, G*, and phase angle, δ, of the binder. Rolling Thin Film Oven test (RTFO) simulates the short term aging, oxidation and volatilization, which occur during the mixing and placement of the asphalt pavement. After completing the simulation, samples were collected for the rheological test using DSR. Pressure Aging Vessel (PAV) test simulates long-term aging that occurs during the in-service life of the pavement. The test is conducted at 19C and involves exposing PAV test takes RTFO-aged samples to heat and pressure over 20 hours. Then the samples were collected for the rheological test using DSR. Bending Beam Rheometer (BBR) was used to test asphalt binders at low temperature, -12C, where thermal cracking normally occurs. The BBR test measures the low temperature stiffness and the rate of change of that stiffness (m-value) of the samples. 142 6.3 Results and discussion 6.3.1 Thermal behavior of ionomer modified asphalt (IMA) DSC data for neat asphalt, ionomer, and IMAs contain 3-9 wt% ionomers are shown Figure 6.2. Two transitions were observed for the neat asphalt, at -29°C and 1.3°C, which correspond to glass transition temperatures (Tgs) for the maltene and asphaltene fractions, respectively [253]. Glass transitions for IMAs shifted to lower temperatures upon the addition of ionomers, see Table 6.1, which indicates mixing of the ionomer into both the maltene and asphaltene phase of the asphalt. The neat ionomer is semi- crystalline polymer had a melting point of 90°C and a heat of fusion of 43.3 J/g. When it was mixed with the asphalt, the melting point dropped to 80°C and the crystallinity of the polymer in the compound decreased. The largest decrease in the ethylene crystallinity occurred for the samples with the two lower ionomer concentrations. Those samples retained about 65-70% of the original ionomer crystallinity. As the amount of ionomer added increased, the crystallinity also increased, which is consistent with a limited solubility of the ionomer in the asphalt. The amorphous ionomer fraction also contains all of the methacrylic acid and methacrylate salt groups. The melting point of the ethylene crystals in the ionomer within the IMA, was about 10°C less than that in the neat ionomer, which is consistent with a diluent effect (i.e., the two compounds mix at some level in the amorphous phase). 143 1 3 0 2 -1 1 Ionomer -2 0 Y Axis 5 Axis Y 4 Axis Y Heat flow (w/g) 5% IMA -3 -1 7% IMA -4 asphalt 3% IMA -2 Exo up 9% IMA -5 -3 -50 0 50 100 150 200 250 Temperature oC Figure 6.2 DSC of base asphalt and ionomer modified asphalt Table 6.1 Thermal properties of IMAs * Xyl sample Tg1 (°C) Tg2(°C) Tm (°C) H (J/g) (%)* asphalt -29.07 1.26 3% -32.37 1.42 80 0.8473 65 5% -31.59 0.22 80 1.552 69 7% -33.65 0.18 81 2.794 92 9% -34.19 -2.15 81 3.100 80 ionomer -120** 90 43 * heat of fusion per mass ionomer ** value for polyethylene The thermal stability of the asphalt, ionomer, and IMAs is shown by the thermogravimetric analysis (TGA) curves in Figure 6.3. The asphalt exhibits a two-step decomposition stage. The first step occurs between 194 – 360°C corresponding to a mass 144 loss of 26 wt%, which is attributed to maltene. The second step, from 360 to 506°C represents the decomposition of the asphaltene. The neat ionomer is stable to higher temperature, and it exhibits a single degradation step starting at ~360°C. The IMAs showed three degradation steps that corresponded to the two degradation processes in the asphalt and the one for the ionomer. The onset temperature for the maltene degradation, however, was about 30°C higher than that of the neat asphalt, which may be due to the interaction of the ionomer with the asphalt or simply to the increase in viscosity that suppresses the diffusion of the degradation products out of the sample 1.0 asphalt 3% 0.8 5% 7% saturates 9% Ionomer 0.6 0.4 Mass remaining fraction Aromatics,resins and Asphaltenes 0.2 0.0 0 200 400 600 800 1000 o Temperature ( C) Figure 6.3 TGA of base asphalt and ionomer modified asphalt 6.3.2 Morphology of ionomer modified asphalt Optical micrographs of an IMA containing 5% ionomer and a similar polymer- modified asphalt containing 5% low density polyethylene (LDPE) are shown in Figure 145 6.4. The same asphalt was used in each sample. The black regions are the asphalt, and bright regions are the ionomer. Figures 6.4(a,b) shows that the LDPE is highly incompatible with the asphalt, forming macroscopic dispersed phases of ~1 mm. This phase separation embrittle the asphalt and affect the performance. However, the ionomer mixed much better with the asphalt forming a much smaller dispersed phase, with the majority of the domains < 100 μm. There are still some large ionomer domains in the IMAs, as shown by the higher magnification photo in Figure 6.4(c), but in that case, they are an order of magnitude smaller and it is also clear that that phase encapsulates some of the asphalt (the black droplets within the ionomer phase). Unfortunately, the nature of the asphalt within the ionomer phase cannot be determined from these data. One could make the argument that the ionomer might preferentially imbibe either the asphaltene or maltene component. The asphaltenes may interact preferentially with the polar ionic species or the maltenes may swell the hydrocarbon phase of the ionomer. Figure 6.4 Optical polarized micrographs of asphalt modified with (a) 5 wt% LDPE, (b,c) 5 wt% ionomer 146 6.3.3 Viscoelastic behavior of IMAs Dynamic shear tests provide information on the behavior of the asphalt and blends in loading modes similar to that of traffic loading. Figure 6.5 shows the dynamic viscosity as a function of frequency at 40°C and 60°C for IMAs mixed at 180ºC for 45 min and 60 min, and Table 6.2 lists the viscosity at a fixed frequency of 1 rad/s as a function of mixing time at 180°C. The mixing time had a significant effect on the viscosity. For a 45 min mixing time there was not a large difference between the viscosities for the IMAs with different compositions, while significant differences were observed when a 60 min mixing time was used. That indicates that a significant mixing time was needed to disperse and mix the ionomer into the asphalt, which may be a consequence of the complicated phase behavior of the system, i.e., interaction of the ionomer with both the maltene and asphaltene components and a swollen ionomer network formed during mixing. Table 6.2 Effect of mixing time at 180°C on η' [Pa⋅s] at 40°C and f = 1 rad/s 30 45 60 asphalt 4500 5000 10,000 3% 3600 4600 4500 5% 4500 6000 50,000 7% 5500 6000 23,000 9% 4400 4700 – For the IMAs mixed for 45 min, the viscosity of the asphalt decreased by an order of magnitude when the temperature was increased from 40°C to 60°C. For low frequencies, the viscosity of the IMA’s were lower than the unmodified asphalt for ionomer concentrations < 3 wt%, and higher than the viscosity of the asphalt for the higher ionomer concentrations. The latter observation was a consequence of the greater 147 shear-thinning behavior of the unmodified asphalt. The reason for the lower viscosity of the IMAs at low frequency is not clear, but may be due in part to poor mixing of the samples for the 45 min. mixing time, though the viscosities of the IMAs at 60ºC were generally higher than that of the base asphalt. For both temperatures, however, the differences in the viscosities of the asphalt and the IMA’s was less than a factor of 2 for the different compositions and less than a factor of 4 due to frequency, 10000 40oC asphalt 3% (a) 5% 7% 9% 1000 ' (Pa-s) 60oC asphalt 3% 5% 7% 9% 100 0.01 0.1 1 10 100 1000 (rad/s) 148 1e+6 40oC (b) asphalt 1e+5 3% 5% 7% 1e+4 ' (Pa-S) 1e+3 60oC 1e+2 asphalt 3% 5% 7% 1e+1 0.01 0.1 1 10 100 1000 (rad/s) Figure 6.5 Dynamic viscosity of base asphalt and inomer modified asphalts at 40C and 60C :a 45minutes mixing ; b, 60minutes mixing For the samples mixed for 60 min, the viscosity changed by a couple of orders of magnitude as the ionomer concentration changed, though the changes were not monotonic with increasing ionomer concentration. At 40°C, the viscosity of the IMA containing 5 wt% ionomer increased by 25 times that of the unmodified asphalt at low frequency and about 3-4 times at high frequency. The unmodified asphalt was not very frequency sensitive, while the viscosity of the IMA with 5% ionomer changed by 1.5 orders of magnitude as the frequency increased by 4 orders of magnitude. That sample did not exhibit a Newtonian region, and the slope of the viscosity vs. frequency plot suggests that that sample exhibited considerable solid-like behavior, which would be expected if the swollen ionomer formed a three-dimensional network. 149 6 5 4 Pa 3 log G' 2 400C 600C asphalt asphalt 1 3% 3% 5% 5% 7% 7% 0 9% 9% -1 -2 -1 0 1 2 3 4 5 log rad/s (a) 6 400C 5 asphalt 3% 4 5% 7% 3 log G' Pa 2 600C 1 asphalt 3% 0 5% 7% -1 -2 -1 0 1 2 lograd/s Figure 6.6 Elastic modulus of neat and inomer modified asphalt at 40C and 60C:a 45minutes mixing ; b, 60minutes mixing 150 Elastic modulus is illustrated in Figure 6.6 as a function of frequency. High value of elastic modulus indicates low resistance to low-temperature cracking, which indicating the material become tough at low temperature and easy to break. At high temperature, the high elastic modulus gives high flexibility. Similar as the results of dynamic viscosity, ionomer modification increased value of elastic modulus over the entire test frequency range. G' of 5% ionomer modification from long time mixing is about 4 times that of base asphalt (Figure 6.6 b). Short time mixing results in poor performance (Figure 6.6 a), in which G' does not show much difference from the base asphalt under the test frequency range. This is probably due to the insufficient interaction between ionomer and asphalt. Zero shear rate viscosity (η0), can be used to determine rutting resistance of modified asphalt. When a shear stress was applied to the sample in a dynamic shear rheometer for a duration long enough that the deformation reaches a constant value, which corresponds to steady state flow. At this stage, the viscosity of the binder was the steady state viscosity or the zero shear viscosity. The shear thinning behavior of our system can be described by Carreau model [254] (Eq10), and η0 can be predicted (Table 6.3). (17) Where, η0 is the viscosity where shear rate is almost zero. is shear rate, which can be converted from frequency. n is a parameter related to the slope of the shear-thinning region and λ is a characteristic relaxation time of the material. Excellent fit of Carreau model was obtained for the 5% ionomer modified asphalt sample (Figure 6.7). The parameters n, λ and fitting precision R2 are summarized in Table 6.3. 151 6.0 5.5 experiment data Carreau Model 5.0 4.5 ' (Pa-s) 4.0 3.5 3.0 0 1 2 3 -1 (s ) Figure 6.7 viscosity data fitted to Carreau model (5% ionomer 60minutes mixing) Table 6.3 Characteristic parameters for viscosity data fitted to Carreau model Sample n λ R2 asphalt 0.77 0.65 0.992 3% 0.85 0.11 0.984 5% 0.67 0.39 0.996 7% 0.84 0.76 0.992 η0 were obtained for both short and long time mixing samples, which is shown in Table 6.4. For the study of asphalt binders, η0 is considered as an indicator of two rutting related binder characteristics [255]. One is the stiffness of binder and the other is the binder’s resistance of permanent deformation under long term loading. Long time mixing led to better performance, which has a higher value of η0. 5% ionomer modified asphalt shows the best rutting resistance of the samples, which is corresponding to the dynamic rheological results discussed above. 152 Table 6.4 Zero shear rate Viscosity ionomer modified asphalt 45 min 60 min asphalt 3.84 4.26 3% 3.75 3.79 5% 3.87 5.2 7% 3.87 3.82 9% 3.74 6.3.4 Time Temperature Superposition Viscoelastic mastercurves of the asphalt data were constructed using the principle of time-temperature superposition (TTS). TTS is based on the equivalence of the effects of time and temperature on the spectrum of relaxation times for a viscoelastic fluid and allows one to determine the viscoelastic properties of a material over a much wider range of time or frequency than is assessable from real-time experiments. The validity of the time-temperature superposition principle has been previously shown for the linear viscoelastic behavior of other unmodified and modified asphalts [256-257]. Frequency sweep of neat asphalt and IMAs at five different temperatures between 30°C and 80°C were conducted. Figure 6.8-6.12 presents the storage and loss modulus curves of neat asphalt and IMAs at each temperature. Mastercurves of the dynamic and loss moduli referenced to 40°C are shown in Figure 6.13 and 6.14 for the IMAs and the neat asphalt. TTS worked fairly well for neat asphalt, 3% and 5% with some deviation observed. However, TTS worked not quite well for higher concentrations, 7% and 9%, especially for G". It is not surprising that TTS failed for 7% and 9% IMAs, since these systems are phase separated and are thermo-rheologically complex. There may be a phase change between 5% and 7% of ionomer in asphalt. The relative success of TTS for lower concentrations may be a consequence of a fairly large difference in the relative terminal 153 relaxation times of the asphalt and low concentration ionomer phase, such that both are not measured in a single isothermal experiment. That would not be surprising, as the ionic interactions in the ionomer have exceedingly long relaxation times. 107 106 (a) 105 104 103 102 G' (Pa) 30 40 101 50 100 60 70 10-1 80 10-2 0.1 1 10 100 1000 (rad/s) 107 (b) 106 105 4 10 30 G'' (Pa) 40 103 50 60 70 2 10 80 101 0.1 1 10 100 1000 (rad/s) Figure 6.8 Shear Modulus of neat asphalt versus frequency at different temperatures (30- 80°C ) (a) Storage modulus and (b) loss modulus. Numbers denote temperatures (°C) 154 107 106 (a) 105 104 103 G' (Pa) 30 2 10 40 50 101 60 100 70 80 10-1 0.1 1 10 100 1000 (rad/s) 107 (b) 106 105 104 30 G'' (Pa) 40 3 10 50 60 102 70 80 101 0.1 1 10 100 1000 (rad/s) Figure 6.9 Shear Modulus of 3 wt% IMA versus frequency at different temperatures (30- 80°C ) (a) Storage modulus and (b) loss modulus. Numbers denote temperatures (°C) 155 107 106 (a) 105 104 103 102 G' (Pa) 70 101 30 100 40 50 10-1 60 10-2 80 10-3 0.001 0.01 0.1 1 10 100 1000 (rad/s) 107 (b) 106 105 104 103 70 G'' (Pa) 102 30 40 101 50 60 0 10 80 10-1 0.001 0.01 0.1 1 10 100 1000 (rad/s) Figure 6.10 Shear Modulus of 5wt% IMA versus frequency at different temperatures (30- 80°C). (a) Storage modulus and (b) loss modulus. Numbers denote temperatures (°C) 156 107 (a) 106 105 104 103 G' (Pa) 30 2 10 40 50 1 10 60 80 100 70 10-1 0.1 1 10 100 1000 (rad/s) 107 (b) 106 105 4 10 30 G'' (Pa) 40 103 50 60 80 2 10 70 101 0.1 1 10 100 1000 rad/s) Figure 6.11 Shear Modulus of 7wt% IMA versus frequency at different temperatures (30- 80°C). (a) Storage modulus and (b) loss modulus. Numbers denote temperatures (°C) 157 106 (a) 105 104 103 102 G' (Pa) 30 101 40 50 0 10 60 70 10-1 80 10-2 0.01 0.1 1 10 100 1000 (rad/s) 107 (b) 106 105 104 3 G'' (Pa) 10 30 40 102 50 60 101 70 80 100 0.01 0.1 1 10 100 1000 (rad/s) Figure 6.12 Shear Modulus of 9wt% IMA versus frequency at different temperatures (30- 80°C). (a) Storage modulus and (b) loss modulus. Numbers denote temperatures (°C) 158 8 6 4 2 Log G' (Pa) asphalt 3% 5% 0 7% 9% -2 -6 -4 -2 0 2 4 6 Log (rad/s) T Figure 6.13 Elastic modulus master curve for base asphalt and ionomer modified asphalt 8 7 6 5 4 LogG'' (Pa) asphalt 3 3% 5% 2 7% 9% 1 -6 -4 -2 0 2 4 6 Log (rad/s) T Figure 6.14 Loss modulus master curve for base asphalt and ionomer modified asphalt 159 The validation of TTS also indicates that IMAs exhibit thermo-rheologically simple linear viscoelastic behavior within the temperature range considered, 30 – 80°C. The reference temperature was chosen as 40C and the shift factors, summarized in Table 6.5, were fit to the Willams-Landel-Ferry (WLF) equation (Figure 6.15) [258], 1 TTC 0 )( T Ta )(log (18) 2 TTC 0 Where, C1 and C2 are empirical constants that can be obtained from experiment data. T0 is the reference temperature. The slope of the low frequency region for storage and loss modulus curves is also listed in Table 6.5. For terminal response, i.e., viscous flow, one expects a slope of 2. The data in Table 6.5 indicate that neither the asphalt, nor the IMAs reached terminal response, i.e., the terminal relaxation of the fluid was longer than the experimental timeframe. However, the low frequency slope increased with increasing ionomer concentration, which indicates that the presence of the ionomer shifted the relaxation times of the compound to shorter times. The constants in equation 1 were c1 ~ 15 and c2 ~ 165 for all the IMAs, which is similar to the values reported by Polacco et al. for SBS- modified asphalt [259]. The elastic modulus, G', increased as the increasing ionomer concentration was increased to 5%, but then decreased as more ionomer was added. The elasticity of the 3% ionomer IMA was about twice that of the base asphalt and at 5% ionomer the modulus was about an order of magnitude greater than that of the asphalt. The shape of G' master curve was roughly the same for all the samples, except for the 7% sample, which showed a flatter frequency dependence in the mid-frequency range. The reason for the increase and then decrease in modulus with increasing ionomer concentration is not know for 160 certain, but may be related to the phase behavior of the ionomer/asphalt mixture. Note from Table 6.1 that the amount of ethylene crystallinity of the ionomer in the compound increases above 5% composition, which was tentatively attributed to the limit of miscibility of the ionomer with the asphalt. 2 asphalt 1 3% 5% 7% 0 9% T WLF fit -1 Log a -2 -3 -4 -20 0 20 40 T-T (K) 0 Figure 6.15 Shift factors (aT) vs. (T – To) used for the G’ and G” data shown in Fig. 12 for the asphalt and IMAs. The reference temperature, To, was 40°C. The curves are the WLF equation fits (see Table 6.5 for the WLF constants used). 161 Table 6.5 WLF constants and slope of viscoelastic master curves Gd )'log( Gd )''log( cionomer C1 C2 ad )log( ad (wt %) T 0 T )log( 0 0 6.4 55.2 1.3 0.99 3% 17.9 199 1.4 0.82 5% 9.0 88.8 1.5 0.82 7% 12.2 94.5 1.3 0.97 9% 10.7 134 1.7 0.95 107 106 105 104 * (Pa-S) 103 asphalt 3% 102 5% 7% 9% 101 10-4 10-3 10-2 10-1 100 101 102 103 104 105 (rad/s) T Figure 6.16 Complex viscosity master curves at 40°C for base asphalt and the IMAs Figure 6.16 shows the master curve of the complex viscosity η* for the neat asphalt and IMAs as a function of frequency. Both the neat asphalt and IMAs exhibit shear thinning behavior, but shear thinning for the IMAs occurs at a lower frequency, 162 which indicates a shorter relaxation time. The onset of shear thinning varies with ionomer concentration, but not in a systematic way. Similar to the effect of ionomer addition on the modulus, the complex viscosity of the IMAs, which were all higher than that of the asphalt at all frequencies, increased with concentration up to 5% ionomer and then decreased for 7% ionomer and increased again for 9% ionomer. That behavior, which was reproducible, is consistent with the idea of a phase change occurring between 5% and 7% ionomer. The increases in the asphalt viscosity with the addition of ionomer were more than two orders of magnitude for the 5% ionomers concentration. 6.3.5 SuperPave test The asphalt industry uses dynamic shear data to evaluate the performance of asphalt binders at medium to high temperatures. Specifically, this characterization is used in the SuperPave asphalt binder performance grade (PG) evaluation. Various SuperPave tests [260] were performed on the unmodified asphalt and an IMA with 5% ionomer by CAP Lab at the University of Connecticut. Viscoelastic behavior was measured with a dynamic shear rheometer (DSR) at 64°C and a frequency of 10 rad/s, which is representative of the time for one cycle of loading corresponding to a vehicle speed of 55 mph. The actual temperatures of the area where the asphalt binder will be placed determine the test temperature. The basic test uses a thin asphalt binder sample, which is sandwiched between two curricular plates. The lower plate is fixed and the upper plate oscillates to apply shear stress. This test is often conducted on unaged, RTFO aged and PAV aged samples. According to Strategic Highway Research Program (SHRP), modified asphalt should have a minimum G*/sin of 1 KPa at its upper service temperature. The experimental results for the neat asphalt and an IMA containing 5% 163 ionomer are given in Table 6.6. Both unaged asphalt and ionomer modified asphalt samples show the G*/sin value higher than 1 KPa at 64C. This indicates that these samples have an upper service temperature higher than 64C. The ionomer modification increased G*/sin value by 44%, suggesting the upper service temperature for ionomer modified asphalt is even higher. The rotational viscosity for the unaged samples was obtained from high temperature test (135C). The purpose of rotational viscosity requirement is to insure the pumpability of asphalt binder during storage and transport. Polymer modified asphalt is difficult to test with standard capillary rheometer. The ionomer modified asphalt has a higher rotational viscosity (0.62 Pa·S) than neat asphalt (0.485 Pa·S) while still under the maximum value (3.0 Pa·S). This indicates the ionomer modified asphalt still has the pumpability. The Rolling thin film oven (RTFO) test simulated the short term aging, oxidation and volatilization, which would occur during the mixing and placement of the asphalt pavement. This test provides a quantitative measurement of the volatiles lost during the aging process. It is known that aging increases modulus and viscosity with little or no influence on activation energy. These rheological property changes of aged sample depend on the combined effect of asphalt oxidation and polymer degradation. Both modified and unmodified samples were exposed to elevated temperatures to simulate manufacturing and placement aging. The results (Table 6.6) show that ionomer modified asphalt still has higher G*/sin value, 4.4 KPa which is twice above the minimum requirement. The higher rutting factor G*/sin indicates less permanent deformation of ionomer modified asphalt at the test condition. The modified asphalt also showed much 164 lower mass loss (0.015%) after RTFO aging. This suggests that ionomer modification improved stability of asphalt at elevated temperature. The Pressure Aging Vessel (PAV) test simulated long-term aging that occurs during the in-service life of pavement, which was conducted at 19C. Usually the PAV test takes RTFO aged samples and exposes them to heat and pressure over a long period (20 hours). After long time aging, the rutting factor, G*/sin of ionomer modified sample was still higher than neat asphalt (Table 6.6), which is a consequence of improved resistance of asphalt to oxidation and degradation. The low rutting factor of neat asphalt indicates the hardening of asphalt during aging process due to oxidation. The bending beam rheometer (BBR) was used to test asphalt binders at low temperature where the thermal cracking usually occurs (-12C). The BBR test provided the creep stuffiness and the rate of change of that stiffness (m-value) of the samples. The creep stiffness is associated to thermal stress in the sample due to shrinking as temperature drop while m-value is related to the ability of the sample to release that stress. The role of maximum limit on creep stiffness and minimum limit on m-value is to make sure that the thermal stress would not be too high and the sample has minimum ability to relieve stress. The test results in Table 6.6 indicate that ionomer modified asphalt built lower thermal stress at low test temperature and had better ability to release that thermal stress without cracking. In General, the super pave test results indicate that ionomer modified asphalt is less sensitive to the aging process than unmodified one. And the modification improves the performance of neat asphalt when subjected to simulated real life service conditions. 165 Table 6.6 Super Pave Test results for ionomer modified and neat asphalt* Original** DSR @64C Spec. req. Ionomer modified asphalt G*/sin 1.905 kPa min. 1.0kPa Unmodified asphalt 1.320 kPa min. 1.0kPa Rotational Viscosity Original @ 135C Ionomer modified asphalt Viscosity 0.62 Pa-S max 3.0Pa-S Unmodified asphalt 0.485 Pa-S max 3.0Pa-S RTFO DSR @64C Ionomer modified asphalt G*/sin 4.41 kPa min. 2.2 kPa Unmodified asphalt 3.457 kPa min. 2.2 kPa Mass Loss/Gain Ionomer modified asphalt -0.015 % max 1.0% Unmodified asphalt -0.603 % max 1.0% PAV DSR @19C max 5000 Ionomer modified asphalt G*/ sin 3617 kPa kPa max 5000 Unmodified asphalt 3401 kPa kPa BBR @-12C max 300 Ionomer modified asphalt Stiffness 121 MPa Mpa m-Value 0.337 min. 0.300 max 300 Unmodified asphalt Stiffness 267 MPa Mpa m-Value 0.317 min. 0.300 *Sample were tested in accordance with AASHTO M 320 **“Original” means the test used the sample without aging Table 6.7 shows DSR and BBR test results at different temperatures for the ionomer-modified asphalt that are used to determine the performance grade of an asphalt. These tests were used to evaluate IMA’s ability to resist permanent deformation and fatigue cracking. For high temperature classification, the binder was tested at 64°C and 70°C. The temperature where G*/sin δ is 1 kPa was taken as the average seven-day 166 maximum allowable pavement temperature (Tmax) for the unaged sample and 2.2 kPa for the RTFO-aged sample (AASHTO specifications for performance grade asphalt binders). An interpolation between the two data points was used to calculate Tmax, which was 69.5°C and 69.2°C for the unaged and aged samples, respectively. The DSR test was also run at 16 – 25°C on PAV-aged IMA b to evaluate its susceptibility to fatigue cracking. The criterion for fatigue cracking is a maximum G*/sinδ of 5000 kPa. (AASHTO M320). The sample was aged at 100°C to simulate a pavement age of 5 to 10 years. The results in Table 6.7 show that the dynamic shear increased with decreasing temperature and exceeded 5000 kPa at ~ 16°C According to the performance grade asphalt binder specifications in AASHTO MP 1, the IMA would be classified as a PG 64. The BBR test is used to determine the minimum pavement design temperature. This test measures the low temperature stiffness and relaxation properties of asphalt binders. PAV-aged IMA was tested at -18°C and -12°C. The test results were obtained 60 seconds after loading the sample. Those test conditions correspond to an experiment run at 10°C lower (i.e., -28°C and -22°C) with a creep time of 2 hours. As a consequence, the experimental data for creep stiffness at 60 seconds and the slope of the master stiffness curve at 60 seconds (m-value) can be used to interpolate between -28°C and -22°C to determine the low temperature limit of the performance grade classification. The specification limits the minimum pavement design temperature are creep stiffness = 300 MPa and m = 0.3. The data in Table 6.7 fall below those limits, so they were extrapolated to determine the lower temperature limit. The stiffness extrapolation indicated a lower temperature of -30.2°C and the m-value extrapolation produced a value of -26.5°C. The 167 higher value was used to designate the lower temperature classification for the IMA. As a result, the actual performance grade of the IMA was 69.2-26.5, which according to AASHTO MP 1 would be sold as PG64-22. Table 6.7 Performance Grade Determination of IMA* DSR Original RTFO aged DSR PAV (IMA) G*/sinδ (kPa) G*/sinδ (kPa) G*/sinδ (kPa) 64°C 1.905 4.409 70°C 0.9398 1.980 25°C 2044 22°C 2609 19°C 3617 16°C 5107 TG (°C) 69.5 69.2 BBR Stiffness (MPa) m-value -18°C 252 0.288 -12°C 121 0.337 TG (°C) -26.5 6.4 Conclusions The ionomer modified asphalt exhibited much better dispersion and smaller phase separation than did polyethylene modified asphalt. As expected the addition of ionomer to asphalt increases its stiffness (high η') and resistance to heavy traffic loading (high G'). The extent of the increase is related to the content of ionomer. 5% ionomer addition has the most profound effect on the dynamic mechanical properties of asphalt. The role of ionomer here is to form a network with asphalt that can provide elasticity for the system. Higher concentration leads to poor compatibility between the crystalline ionomer and asphalt. According to SuperPave test, the resistance to the long-term and short-term aging was enhanced. The results indicate that IMAs are able to withstand permanent 168 deformation and fatigue cracking to some extent. The performance grade 64-28 was transformed to 69.5-26.2 after modification. . 169 CHAPTER VII SUMMARY AND FUTURE WORK This dissertation discussed two research projects: high temperature shape memory polymers and ionomer modified asphalts. The high temperature shape memory project explored the development of shape memory polymers based on thermoplastic material PEEK to achieve high switching temperatures. Sulfonated PEEK can be easily prepared by reacting PEEK with concentrated sulfuric acid and the sulfonation level can be controlled by the reaction time and temperature. PEEK is a thermoplastic semicrystalline material that has good thermal stability and mechanical properties. PEEK can memorize shape by thermally actuation but only to a small extent due to the creep relaxation of crystals. Metal salts of sulfonated PEEK (M-SPEEK) exhibited improved shape memory performance. M-SPEEK were prepared by neutralizing sulfonated PEEK to different metal ions e.g. Na+, Zn2+, Ba2+, Al3+ and Zr4+. The glass transition of SPEEK was significantly increased by dipole-dipole interactions or ionic crosslinking of M-SPEEKs. Tg of M-SPEEK also increases as a function of the ion pair’s Coulomb energy. M- SPEEK exhibited nanophase separated structure due to the strong intermolecular interactions. Linear viscoelastic studies suggested dual shape memory behaviors of M- SPEEK, where the ionic nanodomains due to ionic or dipolar interactions between metal sulfonated groups provide permanent network and the glass 170 transition temperature serves as the switching temperature. The temporary network, which was formed by the glassy state of M-SPEEK, had fixing efficiencies of 80-90% for Na-SPEEK and Zn-SPEEK, and recovery efficiencies as high as 99%. The ionic bonds formed by the Zn-sulfonates were more efficient at shape memory behavior than were dipole-dipole interactions formed from monovalent cations. Ba-SPEEK, Al-SPEEK and Zr-SPEEK showed improved shape fixing efficiency (~93%) and excellent shape recovery (~94%). The advantage of developing trivalent and tetravalent cation neutralized SPEEK is to achieve improved thermal stability and more robust permanent network. By simply changing the cation of SPEEK, one can have a series of shape memory polymers with tunable high switching temperatures. This switching temperature can be further adjusted by changing the sulfonation level of SPEEK. The relatively low shape fixing efficiency of Na-SPEEK and Zn-SPEEK can be resolved by blending with a low molar mass crystalline compound NaOl. In this case, the temporary networks were provided by strong dipolar interactions between the ionomer and a dispersed phase of crystalline NaOl. Chapter IV presents the approaches to develop a shape memory polymer system based on M-SPEEK/NaOl compounds and studies of their properties. The compounds were prepared by solution blending 70wt% M-SPEEKs and 30wt% NaOl. NaOl strongly interacted with M-SPEEK ionomers and formed a dispersed crystalline phase in the ionomer matrix. Part of NaOl was miscible with M-SPEEK, which can be proved by the melting depression of NaOl in the compound and the decrease in Tg of M-SPEEK. The compounds showed excellent dual shape memory behavior with a switching temperature of 230-240°C and fixing and recovery efficiencies close to 100%. In the compounds, two 171 temporary networks were formed, one from the glassy state of the ionomer and the other from the NaOl crystals. The M-SPEEK/NaOl compounds also exhibited Triple shape memory using the two separate switching temperatures of the two temporary networks, Tg of the ionomer and Tm of the NaOl. The Zn-SPEEK/NaOl compounds had a ~50°C temperature window between the two switching temperatures and shape fixing efficiencies for the two temporary shapes of ~90% and shape recovery of the temporary and permanent shapes of ~90% and 100%, respectively. Na-SPEEK/NaOl exhibits similar shape fixing and recovery efficiency but a smaller programming temperature window (~30°C). Changing the composition of the blends, e.g. 20 wt% or 40wt% of NaOl, may affect the shape memory performance. This was not investigated in this research. Previous study in our group using compounds of S-EPDM/ZnOl ranging in composition from 0-50wt% ZnOl developed shape memory polymer systems. In that case, the increasing content of ZnOl led to a increasing of compounds’ modulus and decreasing of shape recovery efficiency. Similar approaches can be conducted with M-SPEEK/NaOl systems. The role of NaOl can be expected to be either plasticizer or reinforce filler depending on the NaOl concentration. The melting temperature of NaOl is 259°C which is about 40°C lower than Tg of Al-SPEEK and Zr-SPEEK. Compounds of Al-SPEEK, Zr-SPEEK and NaOl can be expected to exhibit excellent dual and triple shape memory behavior at a wide range of temperatures. This approach will further explore the application of M-SPEEK based high temperature shape memory polymers. 172 Chapter V discussed asphalt modification by polyethylene base ionomer to increase the asphalt resistance to low temperature cracking. Ionomer modified asphalts with different ionomer content were prepared by addition of ionomer to paving asphalt and their properties and morphology were characterized. Both thermal and morphological studies indicate that ionomer and asphalt give rise to mixtures where certain interactions are established between the two components. From rheological point of view, the improved viscoelastic properties were achieved by ionomer modification. The effect of ionomer concentration on rheological properties was investigated, and 5% seems to be a concentration at which ionomer modified asphalts exhibit best performance. The resistance to the long-term and short-term aging was enhanced by ionomer modification according to the SuperPave tests. The ionomers interact strongly with one or both of the two components of asphalt, i.e., the maltenes and the asphaltenes. The nature of the interaction was not identified, but is expected to be dipolar interactions between the ionic dipoles of the ionomer and the carbonyl or other polar groups in the asphaltene phase and/or compatibility of the hydrocarbon component of the ionomer with the maltene phase. The latter would be promoted by repulsive interactions within the ionomer between the polar and nonpolar species. These interactions produce relatively good compatibility of the ionomer with the asphalt binder, which is manifest as a much finer dispersion of the polymer in the asphalt than produced using polyethylene. The addition of the ionomer to asphalt generally improved the elasticity of the asphalt, but formulation of specific conclusions was complicated by the effect of mixing time on the dispersion and, hence, the properties of the ionomer-modified asphalt. 173 Although the SuperPave properties were improved by a 5% (w/w/) addition of the ionomer to asphalt, in this case, the net improvement was a transformation of the performance grade from PG64-28 to PG69.5-26.2. That improvement may not be economical. However, the initial hypothesis that ionomers may represent an alternative to the polymers currently used for modifying asphalt was demonstrated, at least in part, and future work in this area probably would be worthwhile. Because of the short study time, characterization of a wider variety of ionomers was not achieved. As a consequence, the achievement of only an incremental improvement in the performance grade should not be viewed as a terminal conclusion. There are numerous material variable that were not thoroughly probed in this study, e.g., ionomer molecular weight, degree of ionic functionality, choice of the cation and the anion and the chemical composition of the non- polar part of the ionomer. All of these are expected to impact the viscoelastic properties of the IMA. The key conclusion of this study is that the ionomers were more compatible with asphalt than polyethylene, which is a commonly used polymer modifier for asphalt. In addition to a more thorough investigation of different ionomer functionality and backbone chemistry, a major problem encountered in this work that needs to be addressed in future work is control of the mixing and dispersion of the ionomer into the asphalt. Many of the measurements in this study were not reproducible with different mixes, due largely to differences in the dispersion, even when the mixing conditions (time, temperature, rotor speed) were seemingly the same. Hence, the effect of composition showed diverse and complicated results, especially in the viscoelastic behavior. While those data and the lack of clear trends may be correct (due to a very complicated 174 morphology and, perhaps, partitioning of the ionomer into the maltene and asphaltene phases, it is imperative that morphology be better controlled and confirmed before drawing any conclusions from those data. 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