POLYMERIZATION INDUCED PHASE SEPARATION (PIPS) IN EPOXY / POLY(ε-CAPROLACTONE) SYSTEMS XIAOFAN LUO. For the degree of Master of Science

Transcription

1 POLYMERIZATION INDUCED PHASE SEPARATION (PIPS) IN EPOXY / POLY(ε-CAPROLACTONE) SYSTEMS by XIAOFAN LUO Submitted in partial fulfillment of the requirements For the degree of Master of Science Thesis Advisor: Dr. Patrick T. Mather Department of Macromolecular Science and Engineering CASE WESTERN RESERVE UNIVERSITY January 2008

2 CASE WESTERN RESERVE UNIVERSITY SCHOOL OF GRADUATE STUDIES We hereby approve the thesis/dissertation of candidate for the degree *. (signed) (chair of the committee) (date) *We also certify that written approval has been obtained for any proprietary material contained therein.

3 Dedication To Jialing Shen For all the love and happiness you gave me and the great moments we shared i

4 Table of Contents Dedication...iii List of Tables...vi List of Figures...vii List of Schemes...x Acknowledgements...xi Abstract...xii Chapter 1. Introduction Theory of Polymerization Induced Phase Separation (PIPS) Thermodynamics and Kinetics of Polymerization Induced Phase Separation Application of PIPS PIPS in Epoxy/PCL Systems: Motivation of the Study References...9 Chapter 2. Study of PIPS Process in Epoxy/PCL Systems Background of PIPS in Epoxy/PCL Systems Materials Selection Sample Mixing Technique Cure Technique and Cure Kinetics Mold Design Cure Procedure...18 ii

5 2.4.3 Cure Kinetics Evolution of Morphology during Cure Optical Microscopy (OM) and Turbidity Experiment Technique Determination of Final Morphologies Phase Separation Onset Time (Cloud Point) Morphology Development during Cure Morphology Control References...31 Chapter 3. Thermomechanical Study of Cured Epoxy/PCL Systems Scope Thermal Stability DSC Study Dynamic Mechanical Analysis (DMA) Study Basics of Dynamic Mechanical Analysis (DMA) Experimental Results and Discussion References...60 Chapter 4. Conclusions and Future Work Plans Conclusions Future Work Plans...68 iii

6 List of Tables Table 3.1: Summary of the DSC experimental results...63 iv

7 List of Figures Figure 1.1: Schematic representation of polymerization induced phase separation...12 Figure 1.2: Evolution of the miscibility gap with conversion, P, for (a) UCST and (b) LCST...12 Figure 2.1: Intermolecular hydrogen-bonding vs. intramolecular hydrogen-bonding in epoxy/pcl systems...33 Figure 2.2: Chemical structures of DGEBA, DDS and the crosslinked network...33 Figure 2.3: Chemical structure of poly(ε-caprolactone)...34 Figure 2.4: DSC results of uncured stoichiometric mixtures...35 Figure 2.5: The plot of glass transition temperature vs. PCL weight fraction for uncured mixtures...36 Figure 2.6: The aluminum mold and cast silicone mold for producing up to six cured epoxy/pcl sample bars (with 2 bars shown in the picture)...37 Figure 2.7: Conversion-time curve plotted from the kinetics model...38 Figure 2.8: The turbidity experiment set-up...39 Figure 2.9: Transmittance vs. time curves for DGEBA/DDS/PCL samples...40 Figure 2.10: Phase separation onset time (a) and onset conversion (b) plotted vs. PCL weight fraction...41 Figure 2.11: OM images showing the morphology evolution during cure for (a) DGEBA/DDS/PCL(10) and (b) DGEBA/DDS/PCL(20)...42 Figure 2.12: SEM images for (a) DGEBA/DDS/PCL(10) and (b) v

8 DGEBA/DDS/PCL(20) cured at 180 o C for 3 hrs...43 Figure 2.13: Transmittance vs. time curve for DGEBA/DDS/PCL(20)...44 Figure 2.14: Polarized optical microscope (POM) image of DGEBA/DDS/PCL(20) cured at 180 o C for 3 hrs...45 Figure 2.15: Hot-stage OM images for DGEBA/DDS/PCL(15)...45 Figure 2.16: OM (a) and POM (b) images of DGEBA/DDS/PCL(15) cured at 180 o C for 3 hrs...46 Figure 2.17: SEM images of DGEBA/DDS/PCL(15) cured at 180 o C for 3 hrs (refer to the text for descriptions of region 1, 2 and 3)...46 Figure 2.18: Hot-stage OM images of (a) DGEBA/DDS/PCL(30) and (b) DGEBA/DDS/PCL(40)...47 Figure 2.19: SEM images for (a) DGEBA/DDS/PCL(30) and (b) DGEBA/DDS/PCL(40) cured at 180 o C for 3 hrs...48 Figure 2.20: POM images of (a) DGEBA/DDS/PCL(30) (b) DGEBA/DDS(80)/PCL(30) and (c) DGEBA/DDS(70)/PCL(30) cured at 180 o C for 3 hrs...49 Figure 2.21: POM images of (a) DGEBA/DDS(80)/PCL(30) cured at 180 o C and (b) DGEBA/DDS(80)/PCL(30) cured at 150 o C...49 Figure 3.1: Thermogravimetric curves for cured stoichiometric epoxy/pcl samples and pure PCL...61 Figure 3.2: The chemical structure of 4,4 -methylene-bis(2-chloroaniline)...61 Figure 3.3: DSC results of cured epoxy/pcl and pure PCL samples (a) cooling runs vi

9 at 10 o C/min (b) 2 nd heating runs at 10 o C/min...62 Figure 3.4: DMA results of cured epoxy/pcl samples (a) tensile storage modulus vs. temperature (b) tan delta vs. temperature...64 Figure 3.5: Envisioned network structure of DGEBA/DDS/PCL(40)...65 Figure 3.6: DGEBA/DDS/PCL(40) and pure PCL sample bars under stepwise heat treatment...65 vii

10 List of Schemes Scheme 2.1: Melt blending technique for the mixing of DGEBA, DDS and PCL...34 viii

11 Acknowledgements I would like to acknowledge Air Force Research Lab (AFRL) for funding this research project via the Small Business Innovation Research (SBIR) program, and NEI Corporation for the collaboration on the project. Personally I want to first thank my advisor, Prof. Patrick T. Mather, for all his guidance and support on my research. I will always appreciate his generosity in sharing his scientific knowledge and experience. I would also like to thank Prof. Alexander M. Jamieson and Prof. Christoph Weder for their precious comments and suggestions on my research and future work. Also I want to thank all the members in Mather Research Group at Case Western Reserve University. Working with you is so enjoyable and I have learned so much from you. Special thanks go to Dr. Kyungmin Lee, who helped me tremendously develop my experimental skills; Timothy Marsh and Pamela Knight for the help in my English writing; and Pritesh Patel, who collaborated with me in the late stage of the project. Finally I would like to thank my parents, Ping Fan and Min Luo, and my girlfriend Jialing Shen. It is your love and strong support that keeps me moving forward. ix

12 Polymerization Induced Phase Separation (PIPS) in Epoxy / Poly(ε-caprolactone) Systems Abstract by Xiaofan Luo Polymerization induced phase separation (PIPS) is a widely existing process where an initially miscible, single-phase mixture undergoes phase decomposition during the polymerization of one component, and finally transforms to a phase separated blend. In this thesis, a study of PIPS in DGEBA based epoxy/poly(ε-caprolactone) systems is presented in two parts. The first part involves monitoring of the phase separation process and final morphologies, including an investigation on the impact of processing conditions on ultimate morphologies. The second part contains a series of thermomechanical studies of cured epoxy/poly(ε-caprolactone) blends. It has been discovered that at relatively high temperatures (60 o C~200 o C), depending on the morphology, the material can behave as a high-strength glassy polymer, a low-t g chemically crosslinked semi-interpenetrating network (semi-ipn), or a physically crosslinked pseudo-elastomer. Such epoxy/poly(ε-caprolactone) blends have a great potential and versatility for a large range of applications. x

13 Chapter 1. Introduction 1.1 Theory of Polymerization Induced Phase Separation (PIPS) Polymerization induced phase separation, or PIPS, is a widely used method for producing a large variety of composite and functional materials and has been previously reviewed in literature by different authors. 1,2 In PIPS, the system generally starts with a miscible homogeneous blend of polymer A and monomer (or pre-polymer) B. The polymerization of B is triggered by either heat or ultraviolet exposure in the presence of initiator or crosslinking agent, depending on the chemistry of the system. Polymerization decreases the miscibility of the two components because of the increased molecular weight of B. At a certain point, phase separation occurs via either nucleation and growth mechanism, if the system is far off the critical composition, or spinodal decomposition mechanism, if the system is close to the critical composition. (Figure 1.1) The morphology further evolves with polymerization until the system is frozen by either chemical gelation or vitrification. A large range of morphologies can be obtained by PIPS in a controlled manner, with the two major types being particle/matrix (or sea/island) morphology, and co-continuous morphology, as predicted by the two phase separation mechanisms. However, in reality, the situation is complicated by the competition of reaction rate and phase separation kinetics, as well as the interplay of multiple sub-factors such as viscosity, temperature, non-covalent interactions, and viscoelasticity. Multiple-stage phase separation events can occur in a single system creating complex, sometimes 1

14 hierarchically organized structures. Another way to interpret the polymerization induced phase separation is to use the phase diagram 1, illustrated in Figure 1.2. Here Ф, T and P are the system composition (volume fraction), temperature and reaction conversion, respectively. Considering a system with an upper critical solution temperature (UCST, Figure 1.2a), since most polymer-polymer mixtures exhibit UCST behavior, the system with an initial composition Ф o is initially miscible at the polymerization temperature, T p. As the polymerization proceeds, the miscible mixture becomes less stable as the phase boundary curve shifts upwards. When it reaches the system point (Ф o, T p ), phase separation begins. It is defined as the cloud point (the point when phase separation begins), defining a cloud point conversion, P cp. Similar situations happen for LCST systems where the phase boundary curve shifts downwards during the polymerization, shown in Figure 1.2b. More practically, the time-temperature-transformation diagram (TTT), conversion-temperature-transformation diagram (CTT), and conversion-phase diagram are used to depict the multiple events taking place during the polymerization process. Several references provide details on the use of such diagrams Thermodynamics and Kinetics of Polymerization Induced Phase Separation The basic thermodynamic analysis of polymerization induced phase separation involves the Flory-Huggins (F-H) theory. 6 Considering a system with two components, A and B, the F-H expression for the entropy of mixing is given by 2

15 Δ S = R( N lnφ + N ln φ ) (1) M A B B B where R is the gas constant and N i and φ i are the mole number and volume fraction of the component i, respectively. By defining V i as the molar volume of the component i, and total system volume V T = N A V A + N B V B, the entropy of mixing per unit volume can be expressed as ΔSM φa φb Δ S = = R( lnφa + ln φb) (2) V V V T A B If component A is undergoing polymerization, V A increases with conversion P, since molecular weight increases. The absolute value of entropy of mixing decreases, making mixing less favorable (less contribution from entropy to free energy of mixing). The enthalpy of mixing of the system, according to F-H theory, is expressed as Δ H = Z ε ( V / V ) φφ (3) M c T r A B V r, the reference volume, is defined as the molar volume of the unit lattice in F-H theory. This could be considered as the molar volume of the monomer, or basic repetitive unit of the polymer chains. ε is the contact pair exchange energy (ε = 0.5ε AA + 0.5ε BB ε AB ) and Z c is the coordination number. Thus the enthalpy of mixing per unit volume is Δ H = ( Z ε / V ) φφ (4) M c r A B The above expression indicates that the enthalpy of mixing does not change as a component undergoes polymerization. Therefore it is generally believed that the decrease of entropic contribution is the reason for the phase separation during polymerization. However the prerequisite here is that the interactions between the 3

16 basic units do not change during the polymerization. In reality, this may not hold true due to the change of chemical structure during polymerization, resulting in a change of ε, or χ (F-H parameter, χ = Z c ε/rt). An increase of χ will favor demixing and a decrease of χ will facilitate mixing. The dynamics are more complicated than the thermodynamics for PIPS. In a conventional phase separation process without any polymerization events, it is clear that the dynamics are controlled by concentration fluctuations in the early stage and by diffusion and interfacial tension in the late stage. 7 The domain size follows a simple scaling law, R(t) ~ t α, where α is the growth exponent and depends on space dimension, hydrodynamic effects and composition However, in PIPS the dynamics are complicated by the competition between phase separation dynamics and polymerization rate, or is determined by the ratio 1 : K = phase separation rate / polymerization rate. When the system evolves in the metastable region (the region between binodal and spinodal curves), the two extreme cases are: If K =, equilibrium is instantaneously reached and the system evolves along the binodal curve. Phase separation then proceeds via the nucleation and growth (NG) mechanism. If K = 0, no phase separation will be observed until the spinodal curve is reached. In this case, the phase separation takes place via the spinodal decomposition mechanism (SD), which is a continuous and spontaneous process by a diffusional flux against concentration gradients. The morphology generated displays some degree of 4

17 connectivity (co-continuous). The location of the initial system composition (Φ 0 ) relative to the critical composition (Φ critical ) also plays an important role in determining the phase separation mechanism. SD is highly favored if the system composition is in close proximity to the critical composition, and NG has a higher probability to happen at off-critical compositions. According to a Monte Carlo simulation of the phase separation process in a thermoplastic-modified epoxy 12, SD occurs near the critical point irrespective of the magnitude of the cure reaction rate. There is also sufficient experimental evidence supporting this conclusion. 1 The domain size growth accompanying PIPS can be monitored by different experimental techniques such as small-angle laser light scattering (SALLS) 13 and optical microscopy (OM). For SD mechanism, the Cahn-Hilliard (C-H) theory 14,15 or C-H theory modified with polymerization kinetics 7,13 is useful in explaining many experimental results. The effects of hydrodynamics 16,17 and viscoelasticity 18,19 on the dynamics of PIPS have also been studied by various researchers. 1.3 Applications of PIPS The traditional and most extensive application of polymerization induced phase separation is the toughening of brittle thermosets 20-27, for example, epoxy resins. These highly crosslinked thermoset materials usually have high strength, but their intrinsic brittleness, or poor fracture toughness, limits their use greatly. The fracture toughness, indicated by G IC (fracture energy) or K IC (critical stress intensity factor), 5

18 can be enhanced in several ways. The first general approach is to change the formulation, for example reducing the crosslinking density by using less curing agent, or copolymerizing with a mono-functional component or a long chain flexible spacer. However, this approach introduces a compromise in glass transition temperature. The second approach is the inclusion of a modifier to create a dispersed second phase. According to past investigations 1,2,28, a dispersed phase can successfully increase G IC by several different mechanisms such as particle cavitation, matrix shear yielding, and rubber bridging Polymerization induced phase separation provides an efficient and versatile tool to create two-phase morphologies for the toughening of brittle thermosets such as epoxy resins, and has several advantages over the conventional procedure, the physical blending and dispersion of second phase particles in the original thermosetting monomers. It has a much lower initial viscosity compared with the conventional procedure and there is no problem associated with the stability of the particle dispersion. However, the most important advantage of PIPS is the ability to create a much larger variety of morphologies, many of which cannot be produced by the conventional procedure. Examples include co-continuous, sponge-like 27, or even highly confined crystal-like array 32,33 morphologies. The morphology has a large impact on the mechanical properties of the final material, and therefore the control of morphology is of vital importance to obtain desirable material mechanical performance. To gain maximum enhancement of fracture toughness while maintaining adequate material strength, or G IC, it is generally 6

19 accepted that a co-continuous morphology is the most desirable 1,2,20-22,24,25 because cracks possess a larger tendency to propagate through the more ductile thermoplastic domains. Another opinion is that it is essential to obtain a sponge-like or sandwich-like morphology for toughening purpose 27, but those types of morphologies only exist in systems where the components display large asymmetry of viscoelastic properties. 19,27 Another area where PIPS is extensively used is the polymer dispersed liquid crystals (PDLC) Here, the system normally starts with a homogeneous mixture of liquid crystal (rod-like) molecules and polymerizable monomers, such as epoxy monomers. As a result of the subsequent phase separation upon polymerization (PIPS), a morphology with LC-rich domains uniformly dispersed in a polymer matrix can be created. If the size of the LC-rich domains are controlled to be comparable or larger than the wavelength of visible light, then the resulting material is highly opaque due to extensive light scattering. The LC director can then be oriented upon the application of an external electrical field, making the materials transparent in the direction of the field. This unique property makes PDLC the desired candidate for a large range of electro-optical applications such as optical switches, reflective displays, and variable transmittance windows. 1.4 PIPS in Epoxy/PCL Systems: Motivation of the Study This study is a constituent part of an ongoing project on shape memory assisted self-healing (SMeASH) materials within the research group of Prof. Patrick T. 7

20 Mather. Inspired by both the unique mechanical response of shape memory polymers (SMPs) and S.R. White, et al. s recent advances in the field of self-healing materials 44,45, we are pursuing the design of a novel shape memory assisted self-healing materials that can achieve simultaneous, effective crack closure and restoration of mechanical properties upon proper thermal treatment. Epoxy resins are a class of covalently crosslinked glassy networks and are intrinsically T g -based shape memory polymer. It will be used as the high-strength shape memory matrix material. Poly(ε-caprolactone) (PCL) is selected as the healing agent and also serves as a toughening agent. 46,47 A phase separated morphology is essential to realize the anticipated material functions and can be created by PIPS. Therefore this thesis will mainly focus on the systematic study of PIPS in epoxy/pcl systems. This thesis is structured logically as follows: Chapter 2 will present a study on the PIPS process including the sample preparation method, monitoring of the morphological evolution during PIPS, and some tentative results on the control of morphology by several processing variables. Chapter 3 will involve the study of the thermomechanical properties of epoxy/pcl composite materials after being fully cured, or polymerized, and the coupling between microscopic morphologies and macroscopic material properties. Finally, in Chapter 4, the concluding remarks will be made, and future work plans proposed. 8

21 1.5 References (1) Williams, R. J. J.; Rozenberg, B. A.; Pascaullt, J.-P. Adv. Polym. Sci. 1997, 128, (2) Inoue, T. Prog. Polym. Sci. 1995, 20, (3) Aronhime, M. T.; Gillham, J. K. Adv. Polym. Sci. 1986, 78, (4) Kim, B. S.; Chiba, T.; Inoue, T. Polymer 1993, 34, (5) Biolley, N.; Pascal, T.; Sillion, B. Polymer 1994, 35, (6) Gedde, U. W. Polymer Physics; Chapman & Hall: London; New York, (7) Chen, J.; Huang, H.; Li, M.; Chang, F. J. Appl. Polym. Sci. 1999, 71, 75. (8) Chiantore, O.; Trossarelli, L.; Lazzari, M. Polymer 1988, 39, (9) Grassie, N.; McNeil, I.; McLaren, I. Eur. Polym. J. 1970, 6, 679. (10) Luo, K. Eur. Polym. J. 2006, 42, (11) McNeil, I.; Grassie, N.; Samson, J.; Jamieson, A.; Straiton, T. J. Macromol. Sci., Chem. 1978, A12, 503. (12) Jo, W.; Ko, M. Macromolecules 1994, 27, (13) Kyu, T.; Lee, J. Phys. Rev. Lett. 1996, 76, (14) Chan, P. K.; Rey, A. D. Macromolecules 1996, 29, (15) Chan, P. K.; Rey, A. D. Macromolecules 1997, 30, (16) Huo, Y.; Jiang, X.; Zhang, H.; Yang, Y. J. Chem. Phys 2003, 118, (17) Tang, X.; Zhang, L.; Wang, T.; Yu, Y.; Gan, W.; Li, S. Macromol. Rapid Commun. 2004, 25, (18) Tanaka, H. Phys. Rev. Lett. 1996, 76, 787. (19) Tanaka, H. J. Phys.: Condens. Matter 2000, 12, R

22 (20) Francis, B.; Lakshmana Rao, V.; Jose, S.; Catherine, B. K.; Ramaswamy, R.; Jose, J.; Thomas, S. J. Mater. Sci. 2006, 41, (21) Francis, B.; Thomas, S.; Jose, J.; Ramaswamy, R.; Lakshmana Rao, V. Polymer 2005, 46, (22) Girard-Reydet, E.; Vicard, V.; Pascault, J. P.; Sautereau, H. J. Appl. Polym. Sci. 1997, 65, (23) Johnsen, B. B.; Kinloch, A. J.; Taylor, A. C. Polymer 2005, 46, (24) Oyanguren, P. A.; Aizpurua, B.; Galante, M. J.; Riccardi, C. C.; Cortazar, O. D.; Mondragon, I. J. Polym. Sci., Part B: Polym. Phys. 1999, 37, (25) Ragosta, G.; Musto, P.; Scarinzi, G.; Mascia, L. Polymer 2003, 44, (26) Siddhamalli, S. K.; Kyu, T. J. Appl. Polym. Sci. 2000, 77, (27) Wang, M.; Yu, Y.; Wu, X.; Li, S. Polymer 2004, 45, (28) Unnikrishnan, K. P.; Thachil, E. T. Designed Monomers and Polymers 2006, 9, (29) Wang, X.; Guild, F. J.; Kinloch, A. J. Proc. Annu. Meet. Adhes. Soc. 2004, 27th, (30) Kishi, H.; Uesawa, K.; Matsuda, S.; Murakami, A. J. Adhes. Sci. Technol. 2005, 19, (31) Kim, N. H.; Kim, H. S. J. Appl. Polym. Sci. 2006, 100, (32) Wang, X.; Okada, M.; Han, C. C. Macromolecules 2007, 40, (33) Matsushita, Y.; Furukawa, H.; Okada, M. Phys. Rev. E: Stat., Nonlinear, Soft Matter Phys. 2004, 70, / /3. (34) Serbutoviez, C.; Kloosterboer, J. G.; Boots, H. M. J.; Touwslager, F. J. Macromolecules 1996, 29, (35) Kyu, T.; Chiu, H. W. Polymer 2001, 42, (36) Kloosterboer, J. G.; Serbutoviez, C.; Touwslager, F. J. Polymer 1996, 37, 10

23 (37) Jin, J.-M.; Parbhakar, K.; Dao, L. H. Macromolecules 1995, 28, (38) Hoppe, C. E.; Galante, M. J.; Oyanguren, P. A.; Williams, R. J. J. Macromolecules 2002, 35, (39) Crawford, N. J.; Dadmun, M. D.; Bunning, T. J.; Natarajan, L. V. Polymer 2006, 47, (40) Boots, H. M. J.; Kloosterboer, J. G.; Serbutoviez, C.; Touwslager, F. J. Macromolecules 1996, 29, (41) Liu, C.; Qin, H.; Mather, P. T. J. Mater. Chem. 2007, 17, (42) Rousseau, I. A.; Mather, P. T. J. Am. Chem. Soc. 2003, 125, (43) Liu, C.; Chun, S. B.; Mather, P. T.; Zheng, L.; Haley, E. H.; Coughlin, E. B. Macromolecules 2002, 35, (44) Cho, S. H.; Andersson, H. M.; White, S. R.; Sottos, N. R.; Braun, P. V. Advanced Materials 2006, 18, (45) White, S. R.; Sottos, N. R.; Geubelle, P. H.; Moore, J. S.; Kessler, M. R.; Sriram, S. R.; Brown, E. N.; Viswanathan, S. Nature 2001, 409, (46) Siddhamalli, S. K. Polym. Compos. 2000, 21, (47) Barone, L.; Carciotto, S.; Cicala, G.; Recca, A. Polym. Eng. Sci. 2006, 46,

24 Figure 1.1: Schematic representation of polymerization induced phase separation (a) (b) Figure 1.2: Evolution of the miscibility gap with conversion, P, for (a) UCST and (b) LCST 12

25 Chapter 2. Study of PIPS Process in Epoxy/PCL Systems 2.1 Background of PIPS in Epoxy/PCL Systems The PIPS in diglycidyl ether of bisphenol A (DGEBA) based epoxy resin/poly(ε-caprolactone) (PCL) blend systems have long been of interest to many researchers It has been discovered that whether or not PIPS occurs at all primarily depends on the curing agent used. Only cured with 4,4 -diaminodiphenylsulfone (DDS) could phase separation occur 3,4,6,10 ; otherwise PCL stays miscible with epoxy even after polymerization. 1,2,11 In light of this compositional sensitivity, it is generally agreed that hydrogen-bonding plays an important role in the process. 7,9 In particular, PCL can form intermolecular hydrogen bonding with the pendant hydroxyl groups (resulted from the ring opening reaction with the amine curing agent) of the epoxy, leading to a negative ΔH m, or a decrease of χ relative to the initial uncured system. Therefore the system remains fully miscible and forms an inter-penetrating network (IPN). However, when cured with DDS, the sulfone group of DDS serves as a much stronger electron donor and can preferentially form intramolecular hydrogen bonds with the epoxy hydroxyl groups, which suppresses the intermolecular hydrogen bonds between PCL and epoxy, schematically shown in Figure 2.1. This results in phase separation during the polymerization because of the increased molecular weight during cure and decreased entropic contribution to free energy of mixing, as discussed in the previous chapter. 13

26 In the field of PIPS, DDS cured epoxy resins have been of great interest since they can phase separate with PCL when being polymerized. Chen et al. 10 constructed the phase diagram for a DDS-cured epoxy blended with PCL having a molecular weight of 80,000 Da, and qualitatively explained the macro- and micro-phase separation mechanisms using the phase diagram. Poel et al. 6 studied the morphologies generated by a DDS-cured epoxy blended with various PCL loadings of two different PCL (M w = 5,000 Da and 50,000 Da) and provided experimental evidence for the multiple-stage phase separation events. However, given the fact that a large amount of knowledge has been obtained on the PIPS of epoxy/pcl systems, even more questions are yet to be answered. For example, how can the morphology be controlled by different variables such as polymerization temperature, stoichiometry of the curing agent, and the use of more than one curing agents? What are the thermo-mechanical properties of the cured phase separated blends and the correlations between morphology and mechanical properties? The study presented in this thesis will try to address, to some extent, the questions above. 2.2 Materials Selection The epoxy pre-polymer used in this study was DGEBA (diglycidyl ether of bisphenol A, Aldrich) with an epoxide equivalent weight (EEW, defined as weight in grams of material containing 1 epoxide group) of 172~176. The curing agent selected was DDS (4,4 -diaminodiphenylsulfone, Aldrich). It reacts with DGEBA through the 14

27 ring opening of the epoxide functional groups. Since the functionality of DDS and DGEBA is 4 (4 amino protons) and 2 respectively, the two species create an infinite network when reacted in stoichiometric equivalence of the two functional groups. The chemical structures of DGEBA, DDS, and the crosslinked network are shown in Figure 2.2. The poly(ε-caprolactone), or PCL used has a denoted M n of 42,500 Da and M w of 65,000 Da (Aldrich). The chemical structure is shown in Figure 2.3. Regarding the range of different compositions studied, the samples were named as DGEBA/DDS(A)/PCL(B). The variable A is used to indicate the percentage of the stoichiometric amount of curing agent DDS used. Under stoichiometric conditions, the moles of amino protons of DDS and epoxide groups of DGEBA are equal. If A=100 (stoichiometric condition) then it is neglected. The variable B designates the weight fraction of PCL in respect to DGEBA, calculated as: B mpcl = 100% m + m PCL DGEBA For example, DGEBA/DDS(80)/PCL(30) indicates a sample which contains 80 wt% of stoichiometric amount of DDS, and 30 wt% PCL relative to DGEBA. 2.3 Sample Mixing Technique According to previous studies, DGEBA, DDS, and PCL are miscible over a large 15

28 temperature range. However, the viscosity of the mixture increases dramatically with increasing PCL loading. A relatively high mixing temperature and strong shear are required to ensure thorough mixing. At the same time, there should be little or no polymerization of the epoxy component during mixing in order to maintain the same starting point of study for all the samples. Based on these considerations, a mixing technique was designed utilizing mechanical stirring, summarized in Scheme 2.1 and described in the following. DGEBA and PCL weighed in a glass vial were first mixed at 120 o C by constant mechanical stirring with an Arrow 6000 motor and a custom-made glass stirring rod for 1-2 hours until all the PCL pellets dissolved in DGEBA and the mixture became transparent. The temperature was raised to 140 o C and DDS powder was quickly added. The mixture was mechanically stirred for another 5 minutes, until it became visually homogeneous, and then put in a vacuum oven at 120 o C for 1 hour to remove air bubbles. Finally, a bubble-free, transparent, light yellow homogeneous ternary blend was obtained (Scheme 2.1). To evaluate the quality of the mixing, DSC experiments were carried out for mixed, uncured samples containing stoichiometric amounts of DDS. For the DSC experiments, the samples were first heated up to 100 o C at 10 o C/min, and cooled down to -80 o C at the same rate. Finally a heat run at 10 o C/min to 200 o C was conducted. The 2 nd heat runs for all the stoichiometric samples are shown in Figure

29 In Figure 2.4, each sample only shows a single glass transition temperature and no melting peak of PCL, indicating thorough mixing of the components. In addition, there is little reaction exothermicity at the mixing temperature of 140 o C. Since the mixing lasted only for 5 minutes, it was assumed that cure reaction during that stage was negligible. Figure 2.5 shows the plot of glass transition temperature vs. PCL weight fraction, which displays a reverse proportional relationship. This is due to the plasticizing effect of PCL. The glass transition temperature was determined as the mid-point of the transition in Figure Cure Technique and Cure Kinetics Mold Design A silicone mold for producing cured sample bars with confined dimensions (25 mm * 5 mm * 2.5 mm) was designed. To make the silicone mold, a machined aluminum mold containing raised sections corresponding to the test bar dimensions was used (Figure 2.6 left). A two-part silicone resin system (Polytek Development Corporation, Platsil 73-45, with a cured Shore A hardness of 45) was mixed, degassed at room temperature, and quickly poured into the aluminum mold. It was allowed to stay at room temperature for a sufficiently long time (>24 hrs) to ensure complete cure. The gelled silicone mold was finally peeled off the aluminum mold (Figure

30 right) Cure Procedure To obtain optimum cured resin performance and desired material properties, an appropriate cure procedure is crucial. For most epoxy/pcl systems (unless explicitly stated), the mixture and the mold were first pre-heated at 120 o C, and the mixture was carefully and quickly poured into the silicone mold. An in-mold degas stage was carried out at 120 o C for 2 hours to completely remove any volatile species and trapped air. The temperature was then increased to 180 o C and the samples were allowed to cure for 3 hours, to ensure complete cure, as established by previous studies. 6, Cure Kinetics A previously published kinetics model for DGEBA/DDS epoxy system was adopted. 10 It is expressed as a conversion-time equation as: dp / dt k(1 p) p = (1) with k = Z exp( E / RT) (2) a where p, t, and k are the conversion, reaction time, and reaction rate constant, respectively. Z is the collision factor and log Z = 8.10 min -1. E a is the activation 18

31 energy and in DGEBA/DDS case E a = 79.8 kj/mol. T is the reaction temperature in kelvin. The conversion-time curve plotted from the above equations for 180 o C is shown in Figure 2.7. From the trend shown by the curve it can be inferred that 3 hour curing at 180 o C leads to almost completion for pure DGEBA/DDS. The major limitation of the kinetics model is that it neglects the effects of the PCL loading and the phase separation. The PCL was considered to have a diluting effect and may essentially slow down the reaction rate. Also, in the separated phases, the polymerizations may proceed at different rates, because of the difference in the actual compositions of the two phases. However, this model provided a quick estimation of the reaction conversion at a certain time and was a good starting point. Further experiments, such as chemo-rheology and isothermal DSC, need to be carried out to confirm the true reaction kinetics for the blend systems. Nevertheless this model provides easy approximation of reaction kinetics and is a good starting point for current studies. 2.5 Evolution of Morphology during Cure Optical Microscopy (OM) and Turbidity Experiment Technique Optical microscopy is generally the simplest and most direct way to observe the morphology change during PIPS. In this study, an Olympus BX51 optical microscope, equipped with crossed polarizer and analyzer and an Instec hot-stage were used for 19

32 this purpose. The uncured sample, a bubble-free mixture of DGEBA, DDS and PCL, was sandwiched between a glass slide and cover slip. The temperature of the hot-stage was set to 180 o C. The sample slide was quickly inserted into the hot-stage when the temperature was stable. In order to accurately measure the phase separation onset time, or the cloud point, a turbidity experiment was carried out to assist the optical microscopy. The basic principles as well as the experimental set-up were schematically shown in Figure 2.8. The light from a visible region light source, in our case a halogen lamp, passed through the sample, and finally was detected by a UV/Vis spectrometer (Ocean Optics S2000 miniature fiber optic spectrometer, Figure 2.8a). The transmittance was calculated by It () Tt ( ) = 100% It ( ) 0 where I(t) and I(t 0 ) are the integrated intensity of the nm light wavelength range of time t and t 0, respectively. Transmittance was then plotted with time. When phase separation occurred, because of the fact that the phase domain size was comparable or larger than the wavelength of the visible light, extensive light scattering happened and the transmittance dropped dramatically (Figure 2.8b). Therefore the phase separation onset time, or cloud point, was determined by the point of the initial drop of the transmittance in the curve. 20

33 2.5.2 Determination of Final Morphologies To determine the final morphologies after curing, polarized optical microscopy (POM) and scanning electron microscopy (SEM) were used. In the former technique, a cross-polarizer perpendicular to the polarized light source was utilized so that it could distinguish the PCL (or PCL-rich) and epoxy (or epoxy-rich) phase regions by showing the birefringence of the PCL crystals. SEM provided the view of morphologies at much higher magnifications than OM. For SEM analysis, all cured samples were first immersed in liquid nitrogen, then quickly fractured to avoid any plastic deformation from distorting the morphology. The fractured surfaces were immersed in chloroform for 30 minutes to selectively etch out the PCL. The samples were then dried, coated with platinum and examined with SEM (Philips, Model XL30 ESEM). As the PCL was pre-dissolved, only the epoxy component was visible on the SEM images Phase Separation Onset Time (Cloud Point) The results of the turbidity experiments were shown in Figure 2.9 for all the stoichiometric DGEBA/DDS/PCL samples. For each sample the cloud point was determined by the point of the initial drop of transmittance in the curve. In Figure 2.10 the phase separation onset time and onset conversion calculated using equations (1) and (2) are plotted versus PCL weight fraction. The general trend was that the 21

34 higher the PCL weight fraction, the more time / higher conversion it required for phase separation to begin. Also it s quite interesting that in all cases the phase separation happened at a relatively high conversion of higher than 50%. The dashed line in Figure 2.10b represents 50% conversion Morphology development during cure OM images were taken periodically during the 3 hour cure reaction in order to visualize the evolution of morphology during cure. The results for DGEBA/DDS/PCL(10) and DGEBA/DDS/PCL(20) are displayed in Figure DGEBA/DDS/PCL(10) (Figure 2.11a) started with a transparent, homogeneous solution (0 min). At around 14 minutes (in good accordance with the result of the turbidity experiment, the same for all samples described below) the mixture became cloudy and more turbid as time went on. The images were slightly blurred possibly due to strong back scattering of the dispersed small second-phase droplets. The final morphology could be confirmed by SEM, shown in Figure 2.12 (A). As previously mentioned, PCL was pre-dissolved so only the epoxy component is visible in the SEM image. From the SEM image it is clear that PCL formed spherical particles dispersed in a continuous epoxy matrix. The particles varied in sizes, but typically were several microns in diameter. It is worth mentioning here that PIPS followed by solvent treatment of DGEBA/DDS/PCL(10) might be an easy means of producing porous epoxy thermosets. A similar approach utilizing PIPS was reported in the 22

35 literature, but a thermal treatment (high temperature oxidative degradation of the thermoplastic) rather than a much easier solvent treatment was used. 12 The phase separation process for DGEBA/DDS/PCL(20) was also interesting (Figure 2.11b). The phase separation started at about 18 min and continued via the nucleation and growth mechanism. At a certain point (21 min), all of the small primary droplets quickly dissolved, followed by the appearance of droplets fewer in number, but larger in size (21.5 min). After 22.5 min all of the small primary droplets were not visible by OM. The secondary droplets continued growing slowly until 33 min, after which there was little change in the morphology. The transmittance vs. time curve also recorded this phase dissolution event (Figure 2.13). In Figure 2.13, the first sudden drop of transmittance corresponded to the nucleation and growth of the small primary droplets. At around 20 min the transmittance started to increase due to the dissolution of the dispersed droplets. After passing a maximum peak around 23 min the transmittance decreased again and was approaching a constant value, due to the growth of the secondary larger droplets. All of the times were in good accordance with OM observations. Phenomenologically the best explanation for the experimental observation of this phase dissolution event was the so-called Ostwald ripening process, or evaporation-condensation process, which is a phase coarsening mechanism widely existing in different systems such as superalloys 13, water-oil emulsions 14, and polymer-polymer systems. 15 The process occurs by the growth of large particles at the 23

36 expense of the dissolution of smaller particles. 13 Considering the current case of DGEBA/DDS/PCL(20), the phase separation was initiated by nucleation and growth and this process was kinetically controlled. There was a large surface area in the system because of the large number of small size particles. Thermodynamically this state was not favored because it presented a large amount of excess surface energy. The system therefore continued to evolve to lower the surface energy. Theoretically there exists a so-called critical particle radius, R *, being in thermodynamic equilibrium with the mean matrix composition 13 ; particles with R > R * would grow and those with R < R * would shrink and finally dissolve. The driving force of this process is the total decrease of surface free energy. The final morphology of DGEBA/DDS/PCL(20) was studied by POM and SEM, with the results shown in Figure 2.14 and Figure 2.12b. In the POM image (Figure 2.14), the light white phase regions were PCL-rich because it showed the birefringence of the PCL crystals, while the dark black regions were mainly composed of amorphous epoxy. Therefore it is clear, even a little surprising, that only a 20% inclusion of PCL could form the matrix phase. The result was further confirmed by SEM (Figure 2.12b). Therefore, as the PCL fraction changed from 10% to 20% a phase inversion occurred. In other words, the critical concentration lies between 10% and 20% PCL weight fractions. Thus a sample with 15% PCL content, DGEBA/DDS/PCL(15), was studied. The OM images taken during the PIPS were displayed in Figure

37 From what Figure 2.15 has shown, the phase separation in DGEBA/DDS/PCL(15) occurred via the spinodal decomposition mechanism, which largely differed from DGEBA/DDS/PCL(10) and DGEBA/DDS/PCL(20) where nucleation and growth was the dominant phase separation mechanism. This also indicates the proximity of 15% PCL content to the critical concentration. 6,16 The OM images (unpolarized and polarized) of the final cured samples are shown in Figure A co-continuous type morphology, showing a phase structure with a certain degree of connectivity, was formed. Another interesting point is that there were lots of small blobs within the primary PCL-rich phase domains (see the dotted area in Figure 2.15), which might indicate secondary phase separation (or micro-phase separation termed in some literature) taking place within the primary phase domains. To verify this, SEM images were taken to visualize the morphology profile at a larger magnification, shown in Figure The SEM image of DGEBA/DDS/PCL(15) revealed a very complex morphology. The epoxy (again only epoxy component is visible here) existed in three different forms with different origins: (1) continuous matrix in the primary epoxy-rich phase region, resulted from primary spinodal phase separation as visualized in hot-stage OM images, (2) spherical particles in the primary PCL-rich phase region, and (3) spherical particles in the holes within the epoxy-rich primary phase region. Epoxy regions (1), (2) and (3) are all denoted in Figure In order to elucidate the origin of this complex, somewhat hierarchically 25

38 organized morphology, the physical explanation of the secondary phase separation will be introduced first. When phase separation happens in a system, two phases with different compositions are generated. In the beginning, the system is not very viscous because of low polymerization conversion. The mass transfer between the two phases is relatively easy so the two phases quickly reach the thermodynamic equilibrium compositions, and the system evolution follows the phase boundary curve. However, as the polymerization continues, the viscosity of the system keeps increasing. Thus the mass transfer becomes more and more difficult with time, and the two phase domains gradually become semi- and non-permeable to mass transfer. In this situation the two different phases are considered to be two miscible systems (with different starting compositions) isolated from each other and are behaving the same way as in bulk. As the polymerization progresses the two initially miscible phases becomes less and less miscible, and secondary phase separation may happen within the two primary phase domains. Chen et al. 10 explained the secondary phase separation (which they termed micro-phase separation) in their epoxy/pcl system based on the phase diagram, which provides another way to clearly illustrate the matter. Coming back to the discussion for DGEBA/DDS/PCL(15), it is not difficult to explain the origins of different existing forms of the epoxy component, at least qualitatively. Form (1) resulted from the primary phase separation, as visualized by hot-stage OM images (Figure 2.15). Form (2) existed in the primary PCL-rich phase (where PCL matrix was etched away, see the light region in Figure 2.16 B), and it was formed from the secondary phase separation within the primary PCL-rich phase 26

39 domains. Form (3) is more interesting in that it exists in the holes on the primary epoxy-rich phase region. It could be inferred that the holes were originally PCL particles resulting from the secondary phase separation within the primary epoxy-rich phase domains, so the form (3) epoxy particles are from the tertiary phase separation that happened within the secondary PCL-rich phase domains. Similarly, PCL could exist in different forms as well, with at least two visible in the SEM images: the continuous matrix PCL in the primary PCL-rich phase and PCL particles in the primary epoxy-rich phase resulting from secondary phase separation. For stoichiometric samples with a PCL fraction higher than 20% matrix/particle morphologies were obtained where epoxy formed the particle phase dispersed in a continuous PCL matrix. The hot-stage OM and SEM images for DGEBA/DDS/PCL(30) and DGEBA/DDS/PCL(40) are shown in Figure 2.16 and In both cases the phase separation happened via nucleation and growth. As the PCL content increased the size of the epoxy particles decreased and became more uniform, although these aspects were not quantified. 2.6 Morphology Control It is important to gain control over the morphology to be able to tune the properties of the final phase separated systems. Some preliminary results about various ways to control the morphology will be presented in this section, but more systematic study on this subject will need to be conducted in the future. 27

40 One way to control the morphology is to change the processing variables, such as cure temperature and the amount of curing agent. Both factors could change the reaction rate, relative volume fraction and interactions, and thus change the phase separation process and the morphology generated. Figure 2.20 shows the effect of curing agent amount on the final morphology. All of the samples have the same PCL content (30%) but different (100%, 80% and 70% of the stoichiometric amount of DDS) curing agent amounts. The stoichiometric case, DGEBA/DDS/PCL(30) (Figure 2.20a) as already described before, formed a particle/matrix morphology with uniformly sized epoxy particles distributing in a PCL-rich matrix (see Figure 2.19a for the SEM image). When DDS was used in 80% of the stoichiometric amount (Figure 2.20b), the epoxy-rich particles became much larger in size and fewer in number. The particles were roughly spherical in shape, but much less regular and uniform than DGEBA/DDS/PCL(30). The morphology was somewhat similar to DGEBA/DDS/PCL(20), as shown in Figure When the DDS amount further changed to 70% of the stoichiometric one (Figure 2.20c), phase inversion occurred and the dispersed particle phase became PCL-rich. This approach provides a general way to control the morphology and to incorporate relatively large amounts of PCL while still maintaining epoxy as the continuous matrix, which is generally desirable. The disadvantage of this approach is that it would lead to a significant decrease of epoxy T g, since T g is largely determined by the crosslink density. Figure 2.21 represents an example of controlling the morphology by means of varying the cure temperature. The morphology changed from particle/matrix (Figure 28

41 2.21a) to co-continuous as the cure temperature decreased from 180 o C to 150 o C (Figure 2.21b). This change of morphology may be due to the retarded phase separation rate because of the increased viscosity at a lower temperature. Therefore the system tended to continue phase separation until it crossed the metastable region and reached the spinodal curve. Another variable to control the morphology is the molecular weight of DGEBA, the epoxy pre-polymer. Interestingly, when DGEBA with a much higher molecular weight (~ 1,075 Da, compared with 340 Da for all the other samples) was used, the sample after curing at 180 o C appeared transparent, indicating that there was no phase separation, at least at the length scale of visible light. In addition the DSC curve displays only a single T g around 60 o C, much lower than the T g of fully cured epoxy resins, showing the plasticizing effect of dissolved PCL. This miscibility, which to our knowledge has never been reported in literature, might be due to the excess amount of pendant hydroxyl groups, or hydrogen bonding sites on the backbone of DGEBA, which could not be sufficiently occupied by DDS as in the case of lower molecular weight DGEBA. In between the extreme cases, DGEBA with intermediate molecular weights between 340 Da and 1,075 Da are believed to have the ability to tune the morphology. However at this point we lack adequate experimental evidence and will address this issue in the future. Finally, the morphology can possibly be controlled by some interfacial additives, such as diblock or triblock copolymers. Such additives are often amphiphilic in nature, 29

42 and would act essentially as surfactants to decrease the surface energy thus affecting the ultimate morphology. All of these will be further discussed in the future work plans section in Chapter 4. 30

43 2.7 References (1) Zheng, S.; Zheng, H.; Guo, Q. J. Polym. Sci., Part B: Polym. Phys. 2003, 41, (2) Zheng, S.; Guo, Q.; Chan, C.-M. J. Polym. Sci., Part B: Polym. Phys. 2003, 41, (3) Siddhamalli, S. K. Polym. Compos. 2000, 21, (4) Remiro, P. M.; Cortazar, M. M.; Calahorra, M. E.; Calafel, M. M. Macromol. Chem. Phys. 2001, 202, (5) Remiro, P. M.; Cortazar, M.; Calahorra, E.; Calafel, M. M. Polym. Degrad. Stab. 2002, 78, (6) Poel, G. V.; Goossens, S.; Goderis, B.; Groeninckx, G. Polymer 2005, 46, (7) Ni, Y.; Zheng, S. Polymer 2005, 46, (8) Guo, Q.; Harrats, C.; Groeninckx, G.; Reynaers, H.; Koch, M. H. J. Polymer 2001, 42, (9) Clark, J. N.; Daly, J. H.; Garton, A. J. App. Polym. Sci. 1984, 29, (10) Chen, J.-L.; Chang, F.-C. Macromolecules 1999, 32, (11) Barone, L.; Carciotto, S.; Cicala, G.; Recca, A. Polym. Eng. Sci. 2006, 46, (12) Loera, A. G.; Cara, F.; Dumon, M.; Pascault, J. P. Macromolecules 2002, 35, (13) Baldan, A. J. Mat. Sci. 2002, 37, (14) Koroleva, M. Y.; Yurtov, E. V. Colloid Journal 2003, 65, (15) Chuang, W.-T.; Shih, K.-S.; Hong, P.-D. J. Polym. Res. 2005, 12, (16) Williams, R. J. J.; Rozenberg, B. A.; Pascaullt, J.-P. Adv. Polym. Sci. 1997, 128, 31

44

45 Figure 2.1: Intermolecular hydrogen-bonding vs. intramolecular hydrogen-bonding in epoxy/pcl systems O CH 3 CH 3 O O H 2C HCH 2CO C OCH 2CHCH 2O C OCH 2CH CH 2 H 2N S NH 2 CH 3 OH n CH 3 O DGEBA DDS O N O S N H 2C OH CHCH 2O OH OH OCH 2CH CH 2 N S O N O H 2C CHCH 2O OCH 2CH CH 2 OH N S N Crosslinked Network Figure 2.2: Chemical structures of DGEBA, DDS and the crosslinked network 33

46 O HO (CH 2 ) 5 O C (CH 2 ) 5 O H n Figure 2.3: Chemical structure of poly(ε-caprolactone) Scheme 2.1: Melt blending technique for the mixing of DGEBA, DDS and PCL 34

47 DGEBA/DDS/PCL(40) Heat Flow (Endo Up) DGEBA/DDS/PCL(30) DGEBA/DDS/PCL(20) DGEBA/DDS/PCL(10) 0.5 W/g Temperature ( o C) Figure 2.4: DSC results of uncured stoichiometric mixtures 35

48 10 Glass Transition Temperature ( o C) PCL Weight Fraction (%) Figure 2.5: The plot of glass transition temperature vs. PCL weight fraction for uncured mixtures 36

49 Figure 2.6: The aluminum mold and cast silicone mold for producing up to six cured epoxy/pcl sample bars (with 2 bars shown in the picture) 37

50 1.0.8 Conversion (%).6.4 dp/dt = k(1-p) 1.66 p k = Z*exp(-Eα/RT) T = 180 o C Time (min) Figure 2.7: Conversion-time curve plotted from the kinetics model 38

51 (a) (b) Figure 2.8: The turbidity experiment set-up 39

52 Transmittance (%) 100 DGEBA/DDS/PCL(40) DGEBA/DDS/PCL(30) DGEBA/DDS/PCL(20) 80 DGEBA/DDS/PCL(10) Time (min) Figure 2.9: Transmittance vs. time curves for DGEBA/DDS/PCL samples 40

53 24 Phase Separation Onset Time (min) PCL Weight Fraction (%) (a) 100 Reaction Conversion (%) PCL Weight Fraction (%) (b) Figure 2.10: Phase separation onset time (a) and onset conversion (b) plotted vs. PCL weight fractions 41

54 0 min 14 min 17 min 21 min (a) 0 min 18 min 19 min 20 min 21 min 21.5 min 22.5 min 33 min (b) Figure 2.11: OM images showing the morphology evolution during cure for (a) DGEBA/DDS/PCL(10) and (b) DGEBA/DDS/PCL(20) 42

55 (a) (b) Figure 2.12: SEM images for (a) DGEBA/DDS/PCL(10) and (b) DGEBA/DDS/PCL(20) cured at 180 oc for 3 hrs 43

56 DGEBA/DDS/PCL(20) Transmittance (%) Time (min) Figure 2.13: Transmittance vs. time curve for DGEBA/DDS/PCL(20) 44

57 Figure 2.14: Polarized optical microscope (POM) image of DGEBA/DDS/PCL(20) cured at 180 o C for 3 hrs 14 min 15.5 min 18.5 min 37 min Figure 2.15: Hot-stage OM images for DGEBA/DDS/PCL(15) 45

58 (a) (b) Figure 2.16: OM (a) and POM (b) images of DGEBA/DDS/PCL(15) cured at 180 o C for 3 hrs Figure 2.17: SEM image of DGEBA/DDS/PCL(15) cured at 180 o C for 3 hrs (refer to the text for descriptions of region 1, 2 and 3.) 46

59 0 min 20 min 23 min 33 min (a) 0 min 18 min 19 min 20 min (b) Figure 2.18: Hot-stage OM images of (a) DGEBA/DDS/PCL(30) and (b) DGEBA/DDS/PCL(40) 47

60 (a) (b) Figure 2.19: SEM images for (a) DGEBA/DDS/PCL(30) and (b) DGEBA/DDS/PCL(40) cured at 180 oc for 3 hrs 48

61 (a) (b) (c) Figure 2.20 POM images of (a) DGEBA/DDS/PCL(30) (b) DGEBA/DDS(80)/PCL(30) and (c) DGEBA/DDS(70)/PCL(30) cured at 180 o C for 3 hrs (a) (b) Figure 2.21 POM images of (a) DGEBA/DDS(80)/PCL(30) cured at 180 o C and (b) DGEBA/DDS/(80)/PCL(30) cured at 150 o C 49

62 Chapter 3. Thermomechanical Study of Cured Epoxy/PCL Systems 3.1 Scope This chapter focuses primarily on the thermomechanical properties of cured epoxy/pcl blends. It is important to understand the morphology-property relationships from a technical perspective, to determine end-use applications of the composite materials, and from a scientific perspective, to reveal information on the details of the micro-structures. Various techniques, including thermogravimetry (TGA), differential scanning calorimetry (DSC), and dynamic mechanical analysis (DMA) were used to study the cured systems in detail. 3.2 Thermal stability Thermal stability is an important material parameter that largely determines the possible end-use applications. For multi-component composite materials, thermal stability cannot be simply predicted on the basis of the behavior of each individual component and its relative proportion. 63 Interactions between the individual components play a vital role in determining the thermal stability of the final material. 10 Thermal stability is also scientifically interesting because it reveals useful information on the compatibility of the components. 50

63 The thermal stability of cured epoxy/pcl systems was studied by thermogravimetric analysis (TGA) using a TA Q-500 TGA instrument under a nitrogen gas purge flow environment. For all samples, a small amount of material ranging from 2-10 mg was loaded into the TGA pan and heated at a rate of 20 o C/min to 1000 o C while recording the weight loss as a function of temperature. The results for all stoichiometric DGEBA/DDS/PCL samples are shown in Figure 3.1 along with pure PCL for comparison. All cured epoxy/pcl samples show only a single degradation transition, as shown in Figure 3.1. The onset temperatures of degradation for cured epoxy/pcl samples were higher than pure PCL, consistent with a previous report. 52 The final material, even though phase separated, degraded in a single transition. Conversely, uncured mixtures show multiple degradation steps, corresponding to the individual components. 52 Experimental observations indicate that there is good compatibility between cured epoxy and PCL, possibly due to chain inter-penetration and specific intermolecular interactions such as hydrogen-bonding. Another observation is the noticeable char formation for all cured epoxy/pcl samples. The residue char at 1000 o C increased with decreasing PCL weight fraction (or increasing epoxy fraction). Pure PCL had little char formation at all. Increased char formation is generally considered desirable and can limit the production of combustible carbon-containing gases, decreasing the exothermicity due to pyrolysis reactions, as well as decreasing the thermal conductivity of the surface of a burning 51

64 material. 52, DSC study Differential scanning calorimetry (DSC) was used to measure the thermal transitions of cured epoxy/pcl blends using a TA Q100 DSC instrument in a nitrogen environment. For each sample, 3 to 5 mg of material was weighed, encapsulated in an aluminum hermetic pan, and loaded on the DSC instrument. A first heating run at a rate of 10 o C/min to 200 o C was conducted to remove the thermal history. The sample was then cooled down to -80 o C at the same rate. Finally, a heating scan at 10 o C/min to 240 o C was carried out to measure thermal transitions in the given temperature range. For comparison, an epoxy/pcl sample cured with another curing agent, 4,4 -methylene-bis(2-chloroaniline) (MOCA, chemical structure shown in Figure 3.2), with 30 wt% PCL loading was prepared and examined. In contrast to DDS-cured epoxy/pcl systems, MOCA cannot form intramolecular hydrogen bonds with the pendant hydroxyl groups on the DGEBA backbone 54 ; therefore the cured blend remains miscible and no PIPS occurs. As visually observed, the MOCA cured sample appeared highly transparent, (compared with DDS cured samples which were all opaque) indicating that no phase separation had occurred, at least at the length scale of the wavelength of visible light. DSC curves are shown in Figure 3.3a and b for the cooling and 2 nd heating scans, respectively, for all the above-mentioned epoxy/pcl as 52

65 well as a pure PCL sample. Results from analyses of the DSC data are summarized in Table 3.1. The MOCA cured sample, DGEBA/MOCA/PCL(30) showed only a single, broad glass transition and no PCL crystallization or melting peaks, further indicating the miscibility of PCL and cured epoxy at molecular level. The T g, determined as the mid-point of the transition, was o C, much lower than conventional aromatic amine cured epoxy resins due to the plasticizing effect of PCL. Essentially DGEBA/MOCA/PCL(30) formed an interpenetrating network (IPN). 58 On the contrary, all DDS-cured epoxy/pcl samples showed PCL crystallization on cooling and melting on 2 nd heating. The heats of crystallization (ΔH c ) and heats of melting (ΔH m ) all increase with increasing PCL weight fraction (Table 3.1, columns 3 and 5). All melting temperatures are within a small range around 57 o C, similar to the melting point of pure PCL; however, crystallization temperatures show deviation from pure PCL. PCL in cured samples crystallizes at a higher temperature of 5 to 10 degrees than in the bulk, indicating that the PCL component behaves differently in phase separated blends than in the bulk, possibly due to the restraints resulting from chain penetration and hydrogen bonding. Wide angle X-ray scattering (WAXS) experiments are being carried out to study the chain packing behavior of PCL in cured epoxy/pcl systems and will be discussed in the future work plans in Chapter 4. The glass transition temperatures of epoxy in all DDS cured systems are around 200 o C. Some of the curves show a peak due to physical aging. All epoxy T g s are 53

66 very similar to the T g of pure DDS cured epoxy resin, indicating no or very little plasticizing effect of PCL, or in other words, that the phase separation is quite complete, yielding nearly pure epoxy. Partial solubility of the thermoplastic component in the thermoset-rich phase (even at complete conversion) tends to compromise the T g and mechanical properties and is a common problem associated with many PIPS systems aimed to produce toughened thermoset resins. 1 The high degree of phase separation in the discussed epoxy/pcl systems prevents this material deficiency, making them ideal candidates for many high-temperature engineering applications. 3.4 Dynamic mechanical analysis (DMA) study Basics of dynamic mechanical analysis (DMA) Dynamic mechanical analysis (DMA), or dynamic mechanical thermal analysis (DMTA), is an extremely useful technique for characterizing the thermomechanical properties of materials, especially the viscoelastic properties of polymers. There are different experimental modes, with the forced oscillation mode being the most generally used one. In this mode, a sinuidal strain is applied to the material as ε () t = ε sinwt (1) 0 Considering a viscoelastic material such as polymers, the resulting stress response recorded by the instrument is 54

67 σ () t = σ sin( wt+ δ) (2) 0 where δ is defined as the phase angle, representing the phase lag of the stress response relative to the applied strain. Since σ ( t) = σ sin( wt+ δ) = σ sin wtcosδ + σ cos wtsinδ (3) by defining the storage modulus as σ = δ (4) ε 0 E ' ( )cos 0 and the loss modulus as σ = δ (5) ε 0 E" ( )sin 0 equation (3) can be rewritten as σ () t = E' ε sin wt+ E" ε coswt (6) 0 0 Since all quantities are either known or measurable, E and E can be determined by the instrument. Equations (4) and (5) lead to tanδ = E" E ' which is defined as the damping factor and is also an important material property of interest. It is essentially the ratio of the viscous to elastic response of the material and represents the total energy dissipation behavior. The DMA instrument can determine E, E and tan δ values over a large 55

68 temperature range (temperature sweep mode) or frequency range (frequency sweep mode). Since the DMA curve resembles a spectrum in showing various transitions with changing frequency (or temperature), it is sometimes termed dynamic mechanical spectrum. DMA can further provide very useful information about the structural transitions of the material in the given temperature or frequency range. Major applications include the measurement of T g, various sub-t g transitions, and absolute modulus values It is advantageous over other techniques such as DSC because of its high sensitivity Experimental A TA Q-800 dynamic mechanical analyzer was used for all DMA experiments. The instrument has a force resolution of N, strain resolution of 1 nm, and tan delta sensitivity of There are also multiple deformation modes available, such as tensile, flexural, compression and shear. For all the samples a rectangular bar geometry was used. Mold-processed sample bars (see Chapter 2) were polished to have uniform dimensions of 25 mm (length) * 5 mm (width) * 1 mm (thickness), and loaded in tensile mode on the DMA. The displacement amplitude was set to 10 μm and the frequency was kept constant at 1 Hz. The sample was first equilibrated at -90 o C, and then heated to 240 o C at a constant 56

69 heating rate of 3 o C/min Results and discussion Four samples, each representing a typical morphology type, were chosen for the DMA study. Tensile storage modulus and tan delta are plotted versus temperature in Figure 3.4 (a) and (b), respectively. Different samples, depending on the morphology, showed quite different dynamic mechanical behavior in the given temperature range. DGEBA/DDS/PCL(10), with discrete PCL phase domains dispersed in a continuous epoxy matrix, maintained a high modulus (> 1GPa) until the temperature reached about 200 o C, the T g of the epoxy. DGEBA/DDS/PCL(15), which has a co-continuous morphology, displayed a similar behavior to DGEBA/DDS/PCL(10), but showed a slight decrease in modulus around 60 o C, corresponding to the melting point of PCL ( additive effect ). In both cases, the mechanical properties were dominated by the epoxy component, which is desirable for most high-temperature, high-strength engineering applications. The phase inverted sample, DGEBA/DDS/PCL(40), showed quite different behavior where PCL, as the continuous phase, dominated the mechanical properties. After the melting point of PCL, the storage modulus decreased by 4 orders of magnitude and reached a pseudo-rubbery plateau. 68 The dispersed epoxy particles acted as physical crosslinks with PCL chains penetrating through to form an infinite 57

70 network structure, depicted schematically in Figure 3.5. A simple experiment was conducted to compare the difference in properties between the pure PCL and DGEBA/DDS/PCL(40), shown in Figure 3.6. DGEBA/DDS/PCL(40) and pure PCL sample bars were placed together on a temperature-controlled hot-plate. The temperature was quickly increased from room temperature to 80 o C and held isothermally for 10 to 15 minutes. A picture was taken to visualize the change of macroscopic shapes and appearances of the two sample bars. The same process was repeated for 120, 180 and 200 o C. From Figure 3.6, the pure PCL bar became transparent at 80 o C due to melting of the crystals. As the temperature further increased, the pure PCL started to flow macroscopically under the sample s own weight and the whole mass spread on the glass slide. On the contrary, no visible change was observed for DGEBA/DDS/PCL(40). The epoxy particles, acting as physical crosslinks for the system, were able to maintain the integrity of the macroscopic shape. This result further supports the hypothesis of the network structure as shown in Figure 3.5. Finally, the MOCA cured sample DGEBA/MOCA/PCL(30) showed only a single glass-transition as shown in Figure 3.4a and b, consistent with the DSC results, leading to a decrease of modulus by 2 orders of magnitude. Generally, the DMA results have revealed some quite interesting points: (a) All the samples, irrespective of morphology, showed comparable modulus and 58

71 mechanical properties at room temperature. However, depending on the morphology, their behaviors at elevated temperatures (60 o C to 200 o C) were quite different, ranging from a physically reinforced and crosslinked pseudo-elastomer to toughened, high strength, high temperature glassy polymer to a chemically crosslinked low T g interpenetrating network (IPN). Further mechanical characterization is needed, but it is envisioned that this material system has great versatility and properties that can be tuned by controlling the composition and morphology to fit a large range of applications. (b) All the samples show a rubbery modulus at high temperatures. This feature indicates that, by controlling the morphology, a collection of shape memory materials with different mechanisms can be produced: T g -based, high transition temperature shape memory materials (DGEBA/DDS/PCL(10) and DGEBA/DDS/PCL(15)); T m -based, low transition temperature shape memory materials (DGEBA/DDS/PCL(40)); and T g -based, low transition temperature shape memory materials (DGEBA/MOCA/PCL(30)). The detailed shape memory properties will be characterized in the future. 59

72 3.5 References (1) McNeil, I.; Grassie, N.; Samson, J.; Jamieson, A.; Straiton, T. J. Macromol. Sci., Chem 1978, A 12, 503. (2) Grassie, N.; McNeil, I.; McLaren, I. Eur. Polym. J. 1970, 6, 679. (3) Remiro, P. M.; Cortazar, M.; Calahorra, E.; Calafel, M. M. Polym. Degrad. Stab. 2002, 78, (4) Shau, M.; Wang, T. J. Polym. Sci., Part A, Polym. Chem. 1996, 34, 387. (5) Ni, Y.; Zheng, S. Polymer 2005, 46, (6) Barone, L.; Carciotto, S.; Cicala, G.; Recca, A. Polym. Eng. Sci. 2006, 46, (7) Williams, R. J. J.; Rozenberg, B. A.; Pascaullt, J.-P. Adv. Polym. Sci. 1997, 128, (8) Pierre, A.; Sindt, O.; Thorne, N.; Perez, J.; Gerard, J. F. Macromol. Symp.1999, 147, (9) Goodwin, A. A.; Mercer, F. W. J. Polym. Sci., Part B: Polym. Phys. 1997, 35, (10) Goodwin, A. A.; Simon, G. P. Polymer 1997, 38, (11) Li, Y.; Oono, Y.; Kadowaki, Y.; Inoue, T.; Nakayama, K.; Shimizu, H. Macromolecules 2006, 39,

73 Weight Percent (%) PCL PCL DGEBA/DDS/PCL(10) DGEBA/DDS/PCL(15) DGEBA/DDS/PCL(20) DGEBA/DDS/PCL(30) DGEBA/DDS/PCL(40) Temperature ( o C) Figure 3.1: Thermogravimetric curves for cured stoichiometric epoxy/pcl samples and pure PCL Cl Cl H 2 N CH 2 NH 2 Figure 3.2: The chemical structure of 4,4 -methylene-bis(2-chloroaniline) 61

74 w/g Heat Flow (Exo Up) Pure PCL DGEBA/DDS/PCL(40) DGEBA/DDS/PCL(30) 2 0 DGEBA/DDS/PCL(20) DGEBA/DDS/PCL(15) DGEBA/DDS/PCL(10) DGEBA/MOCA/PCL(30) Temperature ( o C) (a) 0 DGEBA/MOCA/PCL(30) DGEBA/DDS/PCL(10) -2 DGEBA/DDS/PCL(15) DGEBA/DDS/PCL(20) DGEBA/DDS/PCL(30) Heat Flow (Exo Up) -4-6 DGEBA/DDS/PCL(40) Pure PCL -8 1 w/g Temperature ( o C) (b) Figure 3.3: DSC results of cured epoxy/pcl and pure PCL samples (a) cooling runs at 10 o C/min (b) 2 nd heating runs at 10 o C/min 62

75 Table 3.1: Summary of the DSC experimental results Sample Name PCL T c ( o C) ΔH c (J/g) PCL T m ( o C) ΔH m (J/g) Epoxy T g ( o C) * DGEBA/MOCA/PCL(30) / / / / DGEBA/DDS/PCL(10) DGEBA/DDS/PCL(15) DGEBA/DDS/PCL(20) DGEBA/DDS/PCL(30) DGEBA/DDS/PCL(40) Pure PCL / * determined as the middle point of the transitions of the second heating runs 63

76 Tensile Storage Modulus (MPa) DGEBA/DDS/PCL(10) DGEBA/DDS/PCL(15) DGEBA/DDS/PCL(40) DGEBA/MOCA/PCL(30) Temperature ( o C) (a) DGEBA/DDS/PCL(10) DGEBA/DDS/PCL(15) DGEBA/MOCA/PCL(30) Tan Delta Temperature ( o C) (b) Figure 3.4: DMA results of cured epoxy/pcl samples (a) tensile storage modulus vs. temperature (b) tan delta vs. temperature 64

77 Figure 3.5: Envisioned network structure of DGEBA/DDS/PCL(40) Figure 3.6: DGEBA/DDS/PCL(40) and pure PCL sample bars under stepwise heat treatment 65