FTIR Results for the Binary Blends of P2VP and Pvph

MORPHOLOGICAL STUDY OF COMPATIBILIZATION OF IMMISCIBLE POLYMER BLENDS USING A FUNCTIONALIZED BLOCK COPOLYMER

A Thesis Presented to The Graduate Faculty of The University of Akron

In Partial Fulfillment of the Requirements for the Degree Master of Science

Roungrong Thongtan December, 2006 MORPHOLOGICAL STUDY OF COMPATIBILIZATION OF IMMISCIBLE POLYMER BLENDS USING A FUNCTIONALIZED BLOCK COPOLYMER

Roungrong Thongtan

Thesis

Approved: Accepted:

______Advisor Department Chair Chang Dae Han Sadhan C. Jana

______Faculty Reader Dean of the College Thein Kyu Frank N. Kelley

______Faculty Reader Dean of the Graduate School Mark D. Soucek George R. Newkome

______Date

ii ABSTRACT

The effectiveness of a functional diblock copolymer in compatibilization of two immiscible homopolymers was investigated. For this study, polystyrene (PS) and poly(2- hydroxyethyl methacrylate) (PHEMA) were used as a pair of immiscible homopolymers, and polystyrene-block-poly(2-vinylpyridine) (S2VP diblock) copolymer was used as a functionalized compatibilizer. The rationale behind the choice of the ternary blend system lies in that the hydroxyl groups in PHEMA and the pyridine groups in the poly(2- vinylpyridine) (P2VP) block of S2VP diblock copolymer are expected to form hydrogen bonds, enhancing the miscibility between the PS and the PHEMA in the ternary blends. Indeed, the presence of specific interactions (hydrogen bonding) between PHEMA and S2VP diblock copolymer was observed via Fourier transform infrared spectroscopy, inducing a diffuse interphase between PS and PHEMA in the ternary blend during melt blending above the order-disorder transition temperature of the diblock copolymer. However, the interphase was not thermally stable. The diblock copolymer in the interphase was driven into the PHEMA-rich phase due to the specific interactions. Micelles of the diblock copolymer formed when its local concentration exceed the critical micelle concentration threshold.

iii ACKNOWLEDGEMENTS

I would like to express my deepest gratitude to my advisor, Professor Chang Dae Han, for his patience, his never-ending encouragements, and numerous philosophical lessons I have received during the course of my study. I would like to thank Professor Thein Kyu for kind advices and generosity allowing me to use some of his equipments. Also, I want to acknowledge Dr. Weibin Zha for synthesizing the block copolymer used in this research and for his time spent on training me polymer synthesis and microtoming skills. Thanks also go to Cameron Fraser for always be available and promptly helpful on equipment repair and trouble-shooting. Lastly, I would like to thank my family, especially my mother Kunya Thongton, and my dear friends both in Akron and in Thailand for their love and inexhaustible encouragement during the time of happiness and the moment of despair.

iv TABLE OF CONTENTS

Page

LIST OF TABLES ...... vii LIST OF FIGURES ...... viii CHAPTER I. INTRODUCTION ...... 1 II. BACKGROUND 2.1 Polymer Thermodynamics of Binary Systems ...... 4 2.2 Microphases of Block Copolymers ...... 7 2.3 Assessment of Miscibility in Polymer Blend Systems ...... 8 2.4 Compatibilizers ...... 14 III. LITERATURE REVIEW 3.1 Glass Transition Temperature of Blends with Specific Interactions ...... 16 3.2 Diblock Copolymer as an Emulsifier ...... 17 3.3 Nonfunctionalized Diblock Copolymer as a Compatibilizer ...... 19 3.4 Functionalized Diblock Copolymer as a Compatibilizer ...... 21 3.5 Characteristics of Block Copolymer/Homopolymer Blends ...... 22 3.6 Theories of Phase Structure of Homopolymers/Diblock Copolymer Blends ...... 24

IV. EXPERIMENTAL METHODS 4.1 Materials ...... 32 4.2 Sample Preparation ...... 33 4.3 Characterizations ...... 34

v V. RESULTS AND DISCUSSION 5.1 Phase Behavior of Binary Blends of Homopolymers ...... 37 5.2 Morphology of Binary Blends of Homopolymer and Block Copolymer ...... 48 5.3 Morphology of PS/PHEMA/S2VP Ternary Blend System ...... 58 VI. CONCLUSIONS AND RECOMMENDATIONS ...... 69 REFERENCES ...... 71 APPENDIX ...... 78

vi LIST OF TABLES

Table Page

1 The sensitivity limit of various polymer characterization techniques for phase separation detection...... 9

2 Summary of the molecular characteristics of the polymers investigated in this study characterized by GPC method in THF solvent. Note that PHEMA is insoluble in THF solvent; its Mv is characterized by Aldrich...... 34

3 List of blend samples prepared by melt-blending investigated in this study...... 35

4 List of the thermal properties of the homopolymers measured by DSC and TGA...... 35

5 The characteristic interlamellar domain distance of neat S2VP diblock copolymer and the S2VP diblock copolymer in 30/70 PHEMA/S2VP blend...... 57

6 List of dimensions: average domain size of PHEMA-rich phase, the average micelle size of the S2VP diblock copolymer, and average thickness of S2VP accumulated on the interfacial region along PHEMA-rich phase measured from TEM images via an image analysis program...... 64

vii LIST OF FIGURES

Figure Page

1 Chemical structures of the polymer components in the ternary polymer blend studied in this research...... 3

2 TGA analysis of materials in air at a heating rate 10 ºC/min without antioxidant...... 35

3 DSC thermograms of PS20/PHEMA binary blends during the second heating scan. The tick marks indicate the values of the initial points, midpoints, and the final points of the glass transition...... 38

4 Optical micrograph of 30/70 PS20/PHEMA blend at 23 – 245 °C. There was no change of morphology in such temperature range...... 38

5 (a) Optical micrograph of 20/80 PS1.5/PHEMA blend at 245 ºC, and (b) TEM micrograph of 70/30 PS20/PHEMA blend prepared by solvent casting...... 39

6 A schematic representing a formation of an aggregate in the binary blend of PS and PHEMA homopolymers, when PHEMA is a minor component...... 40

7 Optical micrograph of 50/50 PS20/P2VP20 blend at 245 ºC...... 41

8 The turbidity curve of PS3/P2VP3 blend system observed under 500x using an optical microscope...... 42

9 Thermal analysis results of P2VP20/PHEMA blends. (a) Thermograms of the blends, (b) Plot of midpoint Tg as a function of weight composition of P2VP20...... 44

10 IR spectra between 3800 – 2600 cm-1 of P2VP/PHEMA blends measured at 27 °C. Samples prepared by powder method...... 46

11 IR spectra between 3800 – 2600 cm−1 of P2VP/PHEMA blends measured at 180 °C. Samples prepared by powder method...... 47

viii 12 IR spectra showing carbonyl stretching region between 1790 – 1650 cm−1 of P2VP/PHEMA blends measured at 27 °C. Samples prepared by powder method...... 48

13 An optical micrograph showing the bicontinuous structure for the blend of PS20 and S2VP block copolymer at the composition 30/70 after being dried. The micrograph was taken at room temperature. The dark area represents S2VP block copolymer, and the white area represents PS20...... 49

14 TEM image of 30/70 PS20/S2VP blend prepared by solution blending. The dark area represents P2VP block component, and the white area represents polystyrene...... 51

15 TEM image of 70/30 PS20/S2VP blend prepared by solution blending from DMF. The dark area represents P2VP block component, and the white area represents polystyrene...... 52

16 TEM image of neat S2VP showing lamellar microdomains have the interlamellar domain distance of about 12.5 nm. The dark area represents P2VP block component, and the white area represents polystyrene...... 52

17 The schematic describing molecular arrangement in PS20/S2VP blends of our investigation...... 53

18 TEM image 70/30 PHEMA/S2VP blend...... 54

19 TEM image showing the interface region of 70/30 PHEMA/S2VP blend. The dark phase represents the P2VP block component in S2VP diblock copolymer...... 55

20 TEM images of 30/70 PHEMA/S2VP blend. The dark phase represents the P2VP block component in S2VP diblock copolymer. The white area represents PHEMA droplets...... 55

21 Enlarged TEM image of 30/70 PHEMA/ S2VP blend. The dark phase represents the P2VP block component in S2VP diblock copolymer. The white area represents PHEMA droplets...... 56

22 Optical micrographs of ternary blends consisted of PS20/PHEMA/S2VP taken at 23 °C prepared by solvent casting at compositions: (a) 30/70/5phr, (b) 70/30/5phr. Noted the 10-fold difference in scaling...... 59

23 Morphology evolution upon increasing temperature observed via an optical microscope of a 70/30/5phr PS20/PHEMA/S2VP ternary blend prepared by solvent casting. (a) as-dried sample at 23 °C, (b) 135 °C, (c) 155−245 °C...... 60

ix 24 TEM images 70/30/5phr PS20/PHEMA/S2VP melt-blended at 180 °C for 10 min (bef-20-5-180)...... 62

25 TEM images 70/30/5phr PS10/PHEMA/S2VP melt-blended at 180 °C for 10 min (bef-10-5-180)...... 62

26 TEM images 70/30/5phr PS10/PHEMA/S2VP melt-blended at 130 °C for 10 min (bef-10-5-130)...... 63

27 TEM images 70/30/5phr PS10/PHEMA/S2VP melt-blended at 130 °C for 10 min, and simultaneously annealed at 130 °C for 48 h (aft-10-5-130)...... 63

28 TEM images 70/30/5phr PS10/PHEMA/S2VP melt-blended at 180 °C for 10 min, and simultaneously annealed at 130 °C for 48 h (aft-10-5-180)...... 64

29 Schematic presentation of morphological development (a) before, and (b) after annealing at 130 °C in 70/30/5phr PS10/PHEMA/S2VP...... 68

x CHAPTER I INTRODUCTION

Polymer blends have become the materials of interests for many applications. Blending two or more polymers allows us to tailor the mechanical and/or electrical properties suitable for particular applications. The applications vary for functional and nonfunctional polymer blends including engineering materials, reagents and photoresists. However, in preparing polymer blends the compatibility of the polymers has to be taken into consideration. Blending of two immiscible homopolymers will result in a macroscopic segregation forming two phases in the scale of a few hundred nanometers or larger. Such immiscible polymer blends are expected to have a very sharp interface giving rise to poor mechanical properties due to poor interfacial adhesion. In order to improve the mechanical properties of such immiscible polymer blends, a compatibilizer, which will modify the interfacial adhesion, should be added. An ideal compatibilizer should form an “interphase” region, which is defined as a region in which the constituent components coexist.1 The interphase differs from an “interface” in that the thickness of interface would be very narrow (less than, say, 10 nm) in the presence of “repulsive” interactions between the constituent components. The concept of interphase is relevant to adhesive bonding.1 The thickness of an interphase is expected to be larger compared to that of an interface. An effective compatibilizer should be able to form a sufficiently thick interphase for the blend to withstand an external mechanical stress.2 The formation of an interphase is restricted by thermodynamic requirements. The added compatibilizer must have a preferential affinity with one or two homopolymers 1 being blended, which will localize the compatibilizer to situate at the interface between the two immiscible homopolymers.2 One way to optimally compatibilize a binary blend of chemically dissimilar homopolymers (hereafter will be referred to as A/B binary blends) is by adding small amount of a diblock copolymer (A-block–C or C-block-D copolymers) such that one or both components of the block copolymer can have specific interactions with one of the A/B binary blends. In this study, we conducted an investigation on compatibilizing immiscible A/B binary blends using a functionalized diblock copolymer. For the investigation, polystyrene (PS) and poly(2-hydroxyethyl methacrylate) (PHEMA) were used as a pair of immiscible homopolymers, and polystyrene-block-poly(2-vinylpyridine) (S2VP diblock) copolymer was used as a functionalized compatibilizer. Poly(2-vinylpyridine) (P2VP) is a Lewis base having nitrogen atom in the aromatic ring amenable to share a free electron pair with Lewis acids such as PHEMA that contains a hydroxyl group. The blends of P2VP and PHEMA have been observed to be miscible having a lower critical

solution temperature.3 The rationale behind the choice of the ternary blend system lies in that the hydroxyl groups in PHEMA and the pyridine groups in the P2VP block of S2VP diblock copolymer are expected to form hydrogen bonds (Figure 1), enhancing the miscibility between the PS and the PHEMA in the ternary blends, thus enhancing the miscibility between PS and PHEMA and thus giving rise to improved interfacial adhesion between PS-rich and PHEMA-rich phases. In this thesis, the morphologies of the ternary blend were investigated for various parameters, namely, the molecular weight of a homopolymer, blending temperature, and effect of annealing. Also, a comparison of morphology was made between blends processed by solvent-casting and melt-blending. This research illustrated the structural development in the system having specific interactions between polymer species, which differed from ternary systems without strong attractive interactions. 2

Polystyrene Polystyrene-block-Poly(2-vinylpyridine)

CH2 CH CH2 CH CH2 CH n xy N H O

(CH2)2

O O C

CH2 C

CH3 n

Poly(2-hydroxyethyl methacrylate)

Figure 1. Chemical structures of the polymer components in the ternary polymer blend studied in this research.

3 CHAPTER II BACKGROUND

2.1 Polymer Thermodynamics of Binary Systems

Blending the same homopolymer species or isomers with different molecular weights does not guarantee miscibility. According to a study by Cohen and Torradas,4

the blend of 1,2-polybutadiene having a molecular weight (MW) of 30,000 and 1,4-

polybutadiene having an MW of 100,000 are immiscible. There are factors, other than the

molecular weight, that are related to thermodynamic of mixing. The terms “compatibility” and “miscibility” are usually used interchangeably. The term compatibility is usually used to describe “good adhesion” that yields an enhanced property and the ease of mixing.5 The term miscibility is usually used to describe a more molecular scale mixing, when the resultant blend behaves similar to a single-phase system.5 However, “miscibility” does not imply an ideal molecular mixing, but rather an adequate level of molecular mixing to yield macroscopic properties like a single-phase

system.5 Generally, most of the polymers can be interpreted as blends due to the nature of polydispersity. However, the term blend is commonly refers to a mixture of two or more polymers with different chemical structures. According to the thermodynamic theory of mixing (eq 1), polymer blends will be

miscible if the Gibbs free energy of mixing (ΔGm) is negative. The free energy combines the effects of entropy and enthalpy of mixing. It is well established that, for polymer- polymer blends, the entropy term has a weak contribution to the reduction of free energy

4 due to their high molecular weight. Therefore, the concern for polymer miscibility is concentrated mainly on minimizing the enthalpy of mixing. The enthalpy of mixing is directly related to interactions at the segmental level between component polymers.

ΔGm = ΔH m − TΔSm (1)

ΔGm φ A φB = lnφ A + lnφB + χφ AφB (2) RT M A M B ΔH m = χ φ φ (3) RT AB A B

For a negative enthalpy, the polymer pair tends to attract each other chemically. The

Flory−Huggins theory provides an expression of a regular solution model (eq 2) which introduces an interaction parameter (χ), which is directly proportional to enthalpy of mixing (eq 3). The interaction parameter is designated to quantify such chemical interaction. Most chemically dissimilar polymers have a positive χ and they do not mix forming a single phase when the component polymers have sufficiently high molecular weights so that the term TΔSm is negligible. Mixing such two polymers will produce a blend which contains phase separation. Hildebrand6 introduced a solubility parameter δ, which can be used to approximate the interaction parameter according to eq 4. However, the Hildebrand solubility theory works well in the absence of specific interactions such as the nonpolar or weakly polar molecules. Based on the Hildebrand’s theory, the χr now becomes an empirical parameter that includes both enthalpic and entropic contributions. Therefore, the parameter χr empirically varies with temperature as expressed in eq 5.

V χ = r ()δ −δ 2 (4) r RT A B b χ = a + (5) r T

5 Polymer pair that has specific interaction (negative χ) tends to mix well and the molecular weight threshold, at which a blend becomes miscible, is high because of the

strong attractive force. In order to include specific interactions, the Flory−Huggins expression (eq 2) has to be modified by including an additional term:7

ΔGm′ φ A φB ΔGH = lnφ A + lnφB + χφ AφB + (6) RT M A M B RT

where the last term on the right-hand side containing ΔGH accounts for the free energy changes that are a result of specific interactions. The parameters in the ΔGH term may be measured by spectroscopy techniques.7 According to Coleman et al.,7 there are three

reasons why the ΔGH term is not included into the overall χ. First, the ΔGH term has strong composition dependence due to the difference in the equilibrium constants

describing self-association and inter-association. Second, the ΔGH term can be determined from infrared spectroscopy. If χ is assumed to reflect only the physical forces, to a first approximation, χ can be estimated by the solubility parameters. From information on ΔGH and χ available, a phase behavior can be predicted. Lastly, the calculation of an overall χ does not provide any fundamental insight into the balance of intermolecular forces. The definition of the overall χ can be ambiguous.

The calculation of ΔGm/RT can provides a prediction of the phase behavior of polymer blend systems. In a phase diagram there are several factors that affect the miscibility of the blend systems. Such factors are the architecture of the polymer, the choice of monomers, composition or volume fraction, degree of polymerization or molecular weight, the asymmetry parameter or the rigidity of polymer chains, temperature, etc.

6 2.2 Microphases of Block Copolymers

Block copolymers comprise of chemically dissimilar components. A covalent connection in a block copolymer joins two incompatible polymers together, forming an amphiphilic molecule. Block copolymers undergo a transition from an ordered state to a disordered state and vice versa. For the upper critical solution temperature system, the blocks in a copolymer are incompatible at room temperature. The covalent connection reduces the scale of the phase separation to typically a few hundreds angstroms. For a long time, neat block copolymers were thought to exist only in a disordered state because the dimension of the block copolymer microdomains is smaller than the wavelength of visible light. The dimension of the microphases tend to vary with the 2/3 power of the molecular weight, which will be in the order of the radii of gyration of the copolymer blocks depending on the molecular weights of the blocks. It has been discovered that the equilibrium morphology in the ordered state can be spherical, cylindrical, gyroid, or lamellar. In general, the phase behavior and the morphology of neat block copolymers depend on the following factors: (1) the molecular architecture (e.g., diblock, triblock,

starblock, etc.); (2) the degree of polymerization of the block copolymer, ZC; (3) block copolymer composition, f; (4) the interaction parameter between the two components, χ;

2 2 (5) the asymmetry parameter, ε = (ρ 0BbB ) /(ρ 0 AbA ) where ρ 0 A and ρ0B denote the

densities, and bA and bB denote the lengths of the A and B segments, respectively, of the

block copolymer.8 The transition from an ordered state to/from a disordered state or the transition from one equilibrium ordered state to another can be achieved by varying volume fraction of a block constituent or by varying temperature.8 The details of phase transition are discussed in a separate section. In order to investigate the morphology, near equilibrium microstructures have to be achieved by slow casting thin films of the copolymer from a diluted solution in a non-preferential solvent, and annealing above the glass transition temperatures of both components. 7 2.3 Assessment of Miscibility in Polymer Blend Systems

It is difficult to determine ΔGm from experiment. Various techniques have been used to determine miscibility and phase behavior of polymer blend systems. Common techniques that are used include the determination of the glass transition temperatures, microscopy, light or radiation scattering, nuclear magnetic resonance spectroscopy, and infrared spectroscopy. Each experimental technique has limited sensitivity to detect a domain size (Table 1). Therefore, the phase behavior determined from different measurement techniques may vary, and thus may not accurately reflect the thermodynamic phase behavior of the blend. All the methods for determining polymer-polymer miscibility categorized by Olabisi et al.5 are summarized below.

(a) Miscibility Determination by Glass Transition Temperatures

The presence of a single Tg intermediate between that of the pure components suggests miscibility of the polymer pair. A few expressions have been proposed for the composition dependence of the glass transition temperature for miscible blends. Note that, eqs 7-9 are not applicable for a system having specific interactions which cause synergic effect.

(i) Fox equation9

1 w w = 1 + 2 (7) Tg Tg,1 Tg,2

where Tg,i is the glass transition temperature of the pure component i, and wi is the respective weight fraction. This equation may be applicable for a copolymer or a plasticized polymer system that is compatible and not too strongly polar.

8 (ii) Utracki and Jukes equation10

lnT lnT g = w gi (8) T ∑ i T g gi

(iii) Gordon-Taylor expression11

w T + kw T 1 g1 2 g2 Tg = (9) w1 + kw2

where wi is the weight fraction of component i, k is a parameter related to the strength of intermolecular interactions between blend components.

Table 1. The sensitivity limit of various polymer characterization techniques for phase separation detection.1

0.1 nm 1 nm 10 nm 100 nm 1000 nm

Microscopy Optical

Scanning Electron

Transmission Electron

Spectroscopy Infrared

Thermal analysis DSC

DMA

Diffraction WAXS

SANS

SAXS

Light

9 (iv) Kwei equation12

w T + kw T 1 g1 2 g2 Tg = + qw1w2 (10) w1 + kw2 where q is a fitting parameter which corresponds to the strength of hydrogen bonding.

The determination of Tg may be done by calorimetric method. The sensitivity of differential scanning calorimetry (DSC) is reported only detecting a domain size approximately larger than 10 nm.13 This is a widely used technique because of its simplicity. However, the technique has a limitation when the Tg′s of the pure components differ by less than 15 ºC. In order to obtain accurate, distinct Tg′s of two polymers which have small difference in Tg, a heat treatment is needed to achieve the enthalpy recovery. The aging eliminates the kinetic effect during heating and cooling scans. The annealing temperature must be close to the Tg of the pure component with the lowest Tg in such a way that the enthalpy recovery peak of this component approaches equilibrium after a short annealing time, whereas the position of the enthalpy recovery

3 peak of the component with the highest Tg increases.

Also, Tg may be determined by observation of significant reduction of modulus and resilience, and dynamic mechanical thermal analysis (DMTA) where the locations of the loss tangent maxima coincide with the Tg′s. If there are two loss tangent peaks, the blend is immiscible. As the electrical properties of polymers are analogous to mechanical properties, dielectric methods can also determine the Tg. The dielectric constant,ε′ , is similar to compliance, the dielectric loss factor,ε ′′ , is similar to mechanical loss, and the dielectric strength is analogous to tensile strength.5 The dielectric loss factor and the dissipation factor, tan δ ()ε ′′ / ε ′ , are commonly used to ascertain polymeric transitions such as the glass transition.5 Broadband dielectric spectroscopy measures the secondary β relaxation 10 which can occur at temperature below Tg. The secondary relaxation occurs at a shorter length scale than the primary relaxation near Tg. It originates from side group rotation around a bond that connects to the backbone or localized crankshaft motions of four chemical bonds in the backbone.14 Other methods are thermo-optical analysis, radioluminescence spectroscopy, and dilatometric methods. The dilatometric methods are based on the theory that two phases do not have the same free volume. Two-phase behavior in a blend of two different polymers can be determined by two discontinuities in the derivative of dV/dT

5 corresponding to the Tg′s of the respective phases.

(b) Miscibility Determination by Microscopy

Visual observation of phase segregates via transmitted-light and phase-contrast microscopy require as a minimum a difference in refractive index between the phases. Staining has also been used to enhance contrast for optical work.5 Electron micrographs taken via scanning and transmission electron microscopy can observe the phase segregates. The contrast of scanning electron microscopy (SEM) depends on differences on surface topography or texture.5 This can be emphasized by breaking the specimen in its glassy state, or etching one phase away. When the domain phase is chemically etched, dimples can be observed by SEM. Thus, the domain size can be measured. Transmission electron microscopy (TEM) is commonly used to observe the morphology of microdomains of block copolymers, but it can be used to observe a macrophase separation as well.4

The methods can determine the average domain size in a blend. Small angle light scattering (SALS) is less sensitive than small-angle x-ray scattering (SAXS), which 11 means SALS cannot detect small particles such as the case of blends with specific interactions. Techniques utilized are cloud-point methods via SALS, conventional light scattering method, phase-induced critical scattering (PICS), neutron scattering methods particularly small-angle neutron scattering (SANS), and SAXS.

(d) Miscibility Determination by Ternary-Solution Methods

These methods include mutual-solvent and inverse gas chromatography methods.

(e) Miscibility Determination by Rheological Properties.

Han et al.15 illustrated that the slope of log G´ versus log G˝ in the terminal region of monodisperse flexible, entangled homopolymers was 2, independent of molecular weight. The deviation of the slope indicates the presence of inhomogeneity.

(f) Miscibility Determination by Volume of Mixing

Miscible-blend densities are generally higher than those calculated from volumetric additivity relationships especially where specific interactions exist.5

(g) Miscibility Determination by Heat of Mixing by Calorimetry

Some studies use DSC to measure the thermal energy released upon blending. They conclude that the exothermic reaction during blending indicate interaction in a blend.

(h) Miscibility Determination by Melting Point Depression

This method applies to the blend systems that contain crystalline polymers. The utility of the melting point depression allows the calculation of the interaction parameters between polymer species.5

12 (i) Miscibility Determination by Nuclear Magnetic Resonance (NMR) Spectroscopy

The spin-spin relaxation time, T2, can be used to obtain information when small domains of low molecular mobility are dispersed. T2 reflects the mobility of molecules directly, so the temporal change of the mobility can be observed. Such information indicates the degree of heterogeneity between phases. Normally, a shorter T2 corresponds to the mobility of a constrained portion such as a glassy or crystalline phase, while a longer T2 reflects the mobility of a rubbery phase with less-constrained molecular motion. By fractionation and monitoring the change of T2 of each fraction, one may observe the mobility of each phase, and even observe the effectiveness of a compatibilizer.16

(j) Miscibility Determination by Mechanical Properties.

Many studies compare the mechanical properties, especially impact strength and tensile strength, of the blend at different composition to the average properties of the homopolymers. According to Garton,1 the properties of a miscible polymer blend will be close to a rule of mixtures calculation, but the modulus may be slightly higher because of reduction in free volume associated with the mixing process. The improvement of the impact strength after blending a rubbery polymer to a glassy polymer, such as the high- impact polystyrene (HIPS), which is a graft copolymer of polystyrene particles onto polybutadiene matrix, does not indicate immiscibility.1 The further investigation of the interface is needed to ensure such a conclusion.

(k) Miscibility Determination by Infrared and Ultraviolet Spectroscopy.

Fourier transform infrared (FTIR) spectroscopy can identify the interaction between polar functional groups. This technique is sensitive to the local environment of the

13 functional group regardless of the structure of the rest of the molecule. Thus, FTIR spectroscopy is suitable to study mixing in a molecular scale. Formation of a complex or specific interactions can be analyzed. The analysis of polymer-polymer blend is difficult because spectral subtraction requires an isolated absorption which is a characteristic of one component in the mixture. Often, there is spectral overlap which makes the subtraction impossible. The technique should used in junction of other techniques.5

2.4 Compatibilizers

There are two types of compatibilizers, by which differ in their mechanisms of compatibilization. The first type is a reactive compatibilizer which is a graft copolymer that reacts with functionalized homopolymers upon heating such as during melt blending.17-19 The second type, which is the focus of this study, is a nonreactive third component that alters the interfacial properties and may induce the formation of an interphase. The formation of an interphase is restricted by thermodynamic requirements. The effectiveness of a nonreactive compatibilizer on improving interfacial adhesion depends on the chemical affinity of the compatibilizer to the blend components. The added compatibilizer must have a preferental affinity with one or both homopolymers being blended, which will localize the compatibilizer to situate at the interface region between the two immiscible homopolymers.2 Some studies attempt to add nonreactive third components that only decrease the interfacial energy between incompatible homopolymers due to emulsification effect. They are categorized as emulsifiers. As the affinity of the third component to the homopolymers is not energetically preferable, an emulsifier still produces a sharp interface. Therefore, an emulsifier does not improve interfacial adhesion and it is not an effective compatibilizer. There are many ways to improve the properties of a blend of two immiscible homopolymers by adding a nonreactive third component. Theoretically, a nonreactive 14 third polymer C, which has chemical affinity with the two immiscible homopolymers, A and B, is an option for a compatibilizer. The disadvantage is the difficulty to synthesize a single-component homopolymer C having such property. Therefore, a modified polymer such as a random copolymer, a graft copolymer, a diblock copolymer, or a triblock copolymer may be used. Some additives only reduce the size of the dispersed aggregate and improve the properties slightly, while some are acting as compatibilizing agents which improve the properties significantly.

15 CHAPTER III LITERATURE REVIEW

3.1 Glass Transition Temperature of Binary Blends with Specific Interactions

Kwei12 found that a miscible blend system with specific interactions exhibits a single Tg which does not follow the additivity rule of the Fox equation. Some systems have the Tg with positive deviation that is proportional to the product of the weight fractions of the components. The Tg that is higher than the average Tg of the constituent components has been observed and it reached the maximum near 50% composition for the P2VP/poly(4-vinylphenol) (PVPh)14 and P4VP/PVPh20 suggesting a possible formation of a hydrogen-bonded complex. The increase of Tg is attributed to a decrease of chain mobility, which may be viewed as effective crosslinks as a consequence of the strong interactions. Some systems exhibit S-shaped curves. Kwei12 speculated that the S-shaped curves derived from the combination of mixing term that is concave toward the weight-average line while the hydrogen bond term is convex toward it. Another study21 observed that the flexibility of the polymer backbone may influence the formation of specific interactions. The study mixed a family of poly(N-1-alkyl itaconamic acids), which had various monomer lengths, with either P2VP or PVPh. The Tg of the blends was found to behave in a sideway-S-shaped curve upon increasing blend composition. The flexibility of the polymer chains may reduce the steric hindrance and enhance the probability of having specific interaction between the functional groups.

16 Cesteros et al.3 studied the miscibility and specific interactions in blends of PHEMA with poly(2-vinylpyridine) and poly(4-vinylpyridine). They concluded that number of hydrogen bonding from intermolecular interaction increases as the amount of polypyridine increases, because there is more possibility of N-OH bonding. They found that the blend of PHEMA/P2VP exhibits a lower critical solution temperature (LCST) behavior at 140-175 °C, while the blend of PHEMA/P4VP does not phase separate in the observed temperature range.

3.2 Diblock Copolymer as an Emulsifier

There are many studies reporting on the addition of a nonfunctional diblock copolymer (A-block-B or A-block-C copolymers) into an immiscible A/B binary blend, which resulted in a reduction of the dispersed aggregates and a slight improvement in mechanical properties.22 The improvement of the impact strength is a result of the smaller phase aggregates due to the reduction in the interfacial tension at the interface between immiscible homopolymers, which was observed via optical microscopy. Many researchers claim that there is also an interpenetration of each block into the homopolymer phase of the identical chemical structure, which contributes to the improvement in mechanical properties. However, there is no proof supporting such assumption, and the interface between the homopolymer-rich phase and the copolymer domain seems to remain sharp. The sharp interface indicates an absence of any kind of attractive interaction contributed to the reduction in entropy of mixing. In such situations, the block copolymer played a role of emulsification which reduced the surface tension by physically wetting the aggregate interface. Hlavata´et al.23 studied the PS/PP blend with addition of PS-block-PB. The interface between the micelles and the domain seems to remain sharp. The interaction parameters

17 in the system are χPS/PP > 0, χ PS/PB > 0, and χPP/PB > 0. An addition of a nonfunctional diblock copolymer into an immiscible PS/PP binary blend results in a size reduction of the dispersed aggregates and a slight improvement in impact strength. The morphology illustrated small droplets of the copolymer wetting partially at the interface of the discrete macrophase. The interface was sharp and the block copolymer did not surround the microphase, which showed that there were no attractive interactions between the copolymer and the homopolymers. Upon annealing, the domain size increases. PS- block-PB did not stabilize the macrophase. An effective compatibilizer should suppress the stability of macrophase and promote the stability of microphase formation.24 Fayt et al.25,26 attempted to use hydrogenated polybutadiene-polystyrene block copolymer to compatibilize PE/PS blends. They found that the addition of the diblock copolymer resulted in the reduction of the domain size (10-20 times) and improve the tensile strength. The higher the molecular weight of the diblock the smaller the domain size. In another work, Fayt et al.27 attempted to observe the location of the copolymer via TEM. They inserted a short segment polyisoprene in between the original diblock copolymer. They observed a continuous layer of the diblock copolymer around the discrete phase in PS/LDPE blend. The layer presented a sharp interface. The continuous layer was possible despite an absence of molecular attractive interactions because the

28 melt blended temperature was higher than the TODT of the block copolymer. Anastasiadis et al.29 observed a large decrease in interfacial tension with the addition of the PS-block-PB in PS/PB blends. The interfacial tension was reduced significantly at the extremely low concentration of the diblock copolymer. The interfacial tension approaches a constant value when the concentration of the copolymer exceeds an apparent critical micelle concentration (CMC). The formation of micelle was also observed by Löwenhaupt and Hellman30 in PS/PMMA/(PS-block-PMMA) blends. They demonstrated that the microphase-to- 18 macrophase transition may be induced by varying the amount of the added block copolymer. The process was apparently reversible. Note that the molecular weight of each block in the diblock copolymer was higher than the homopolymers. The blend system exhibits micelle formation of the diblock copolymer inside PS droplets and in PMMA matrix when the copolymer volume fraction was less than 0.2. They also observed the transition from the micelle to lamellar structure when the copolymer concentration was greater than 0.2. The phase behavior of the A/B/block copolymer ternary mixtures is quite complicated because of the interplay of macro- and micro-phase transitions. Chen et al.31 found that the phase separation kinetics of PMMA/poly(vinyl acetate) (PVAc) binary blends was either retarded or accelerated, depending on the volume fraction of a block copolymer added. However, the block copolymer retarded the phase separation kinetics

32 when the block copolymer was homogeneous at T > TODT.

3.3 Nonfunctionalized Diblock Copolymer as a Compatibilizer

Ruegg et al.33 studied an A/B/(A-block-C) ternary system where block C of the copolymer exhibits repulsive interactions with the homopolymer A but attractive interactions with homopolymer B. Component A was a saturated polybutadiene with 89% 1,2-addition. Component B was polyisobutylene. Block A of the diblock copolymer was chemically equivalent to homopolymer A, and block C of the diblock copolymer was saturated polybutadiene with 63% 1,2-addition. The homopolymers A and B are weakly segregated ( χ AB N ≈ 2 ) having upper critical solution temperature

(UCST) behavior. The study investigated the effect of the addition of A-block-C copolymer at φ A − C = 0.4 and 0.5 into a fixed volume fraction ratio of homopolymers A and B at φ A = 0.493, which is the critical composition of the binary blend. Due to the interplay of the attractive and repulsive interactions, the ternary system exhibited 19 microphase separation at low temperatures, homogeneous state at intermediate temperatures, and macrophase separation at high temperatures. Measurements were done by small-angle neutron scattering (SANS), and small-angle light scattering (SALS) was utilized to confirm the cloud points where macrophase separation occurred. As expected, the homopolymer A penetrated into the A-rich microphase, while the homopolymer B penetrated into the B-rich microphase, since the molecular weight of the diblock copolymer was about 6 times larger than that of the homopolymers. The results were compared with the predictions from the random-phase approximation (RPA) for the homogeneous regime, self-consistent field theory (SCFT) for the microphase-separated regime, and Flory−Huggins theory (FHT) for the macrophase-separated regime. The experimental results mostly quantitatively agreed with the predictions from these theories within acceptable experimental errors. For the reason that RPA only predicts the limit of stability of the homogeneous phase, the metastable disordered microphase (such as microemulsions) which scattered by SANS cannot be described by RPA theory. The authors applied a modified Teubner-Strey analysis34,35 to demonstrate a slight evidence of a transition from a stable lamellar phase to a metastable microemulsion with increasing temperature before the blend becomes homogeneous. A study by Chun and Han2 observed lamellar microdomain of diblock copolymer surrounding discrete aggregates. They compared the interfacial morphologies of three polymer systems that consisted of two homopolymers A and B, and a C-block-D copolymer. These systems are more complicated than the A/B/(A-block-C) ternary system. The A/B/(C-block-D) system is associated with 6 pairs of interaction parameter. The ternary systems investigated in their study were PS/1,2-PB/poly(α-methylstyrene- block-1,4isoprene), poly(phenylene oxide)/polypropylene/(PS-block-PEB), in which PEB denotes poly(ethylene-co-1-butene). The blends were prepared by melt blending. TEM micrographs showed lamellar microdomain of block copolymers surrounding discrete 20 domain when there were attractive interactions between PS and PPO, between 1,2-PB and 1,4-PI, and between PP and PEB. They also found that the block copolymer diffused into a homopolymer-rich phase when the interaction parameter between the homopolymer and one of the blocks had attractive interactions. TEM micrographs also have shown evidence that the diblock copolymer was located at the interface only when the melt blending temperature was higher than the TODT of the block copolymer. The thickness of the lamellar microdomains on the interface was about 30 nm. They concluded that the attractive interactions were insufficient to compatibilize high-viscosity polymers unless they were mixed at temperatures above the TODT of the block copolymer.

3.4 Functionalized Diblock Copolymer as a Compatibilizer

Riemann et al.36,37 compared the effectiveness of the polystyrene-block- poly(methylmethacrylate) copolymer (i.e. PS-block-PMMA) to the poly(cyclohexylmethacrylate)-block-poly(methylmethacrylate) copolymer (i.e. PCHMA- block-PMMA) in a PS/PMMA binary blend. They concluded that the dispersion power or the reduction in size of the dispersed aggregates depended on the entanglement probability of the added block copolymer. Upon annealing, the aggregates in the PS/PMMA/(PS-block-PMMA) system expanded with time; while the study showed that the PCHMA-block-PMMA copolymer stabilized the interface. The dynamic mechanical properties of both systems were similar. The storage modulus was generally improved especially in the low-frequency range with increasing content of the diblock copolymer. This suggested am improvement of elasticity due to a better interfacial adhesion. Since the entanglement probability of PS is higher than PCHMA for the same molecular weight. They concluded that the enthalpic interactions are far more effective to the stability of the dispersed aggregates than the entanglement effect.

21 Kobori et al.38 investigated the surfactant effect which may stabilize the interfaces. The blend system of polyethylene (PE)/PVPh/polyethylene (PE)-block-poly(methyl methacrylate) (PMMA) copolymer was investigated via scanning electron microscopy (SEM) and optical microscopy. The study showed an improvement in dispersion and the size reduction of segregated particles with increasing amount of added block copolymer. The block copolymer also stabilizes the interfaces so that there is no change in the particle sizes even after annealing. Kobori et al.39 compared TEM micrographs of the interfacial region of linear low- density polyethylene (LLDPE)/PVPh/(PE-block-PMMA) ternary system with the nonfunctional system that consisted of LLDPE/PMMA/(PE-block-PMMA) ternary system. For the LLDPE/PVPh/(PE-block-PMMA) system, it was demonstrated by FTIR spectroscopy that hydrogen bonds were formed between PVPh and the PMMA block. The TEM micrographs revealed undulated interface at the interface and belt-like domains for LLDPE/PVPh/(PE-block-PMMA) ternary system, while the LLDPE/PMMA/(PE- block-PMMA) ternary system had spherical domains with sharp interface. Also, the aggregate size for the system containing hydrogen bonds was much smaller. The study suggested a significant morphological difference in the ternary systems depending on the choice of the compatibilizer.

3.5 Characteristics of Block Copolymer/Homopolymer Blends

Sadron and Gallot40 observed structural development of block copolymer/solvent systems. Three structures were seen with increasing block copolymer concentration. In the first region, at very low concentration (< 0.1%), a homogeneous solution in which the copolymer is molecularly dispersed is obtained. In the second region, the concentration of the block copolymer exceeds its solubility limit, and micelles of copolymer are formed. The limit is called the critical micelle concentration (CMC), which is defined as 22 the concentration at which micelles first appear. In the third region, above a certain copolymer concentration, the system is organized in a regular periodic structure that is similar to those found for neat block copolymers. The first and the second regions were reviewed by Tuzar and Kratochvil.41 They presented evidence for the micellization being a similar process characterized by an equilibrium between a single molecularly dispersed molecule and aggregates made up of n (n is on the order of 100) copolymer molecules.42 It was discovered that when the molecular weight of a homopolymers is sufficiently low compared to its fraction in the block copolymer, the homopolymer is soluble in one of the block copolymer microdomains.43-46 This result suggests diffusion of the homopolymer A into the A-rich microdomain, which may thicken the microdomain characteristic spacing. Cohen and Torradas47 found that a homopolymer of low molecular weight can induce microphase separation of an originally homogeneous diblock copolymer by increasing homopolymer volume fraction. The same homopolymer with higher molecular weight does not interact with the diblock copolymer, so that the homopolymer formed discrete domains of several times larger in size. A nonequilibrium bicontinuous structure was observed by Han et al.48 via TEM investigation for (PS-block-PI)/PI blends. Their samples were prepared by solvent casting. Upon heating, the bicontinuous structure disappeared. Having their PS-block-PI copolymer in disordered state at room temperature, they concluded that the bicontinuous morphology evolved from frozen composition fluctuations in the disordered phase near order-disorder transition. Kinning et al.49 obtained TEM micrographs of PS/(PS-block-PB) blends, at near- equilibrium when the copolymer was a minor component. Micelles were formed having a core contained PB block, and the corona region contained PS-block and PS- homopolymer chains. By adding various molecular weights of PS homopolymer into 23 symmetric PS-block-PB copolymers, they observed structural transitions of micelles from spherical to wormlike cylindrical and vesicular upon increasing molecular weight. These transitions are similar to the behavior of block copolymer/solvent systems which relates to the relative volume fraction of the core and the corona of the micelle. When a preferential solvent is added, the effective volume fraction of an attractive block increases; the phase becomes swollen which results in the decrease of the interface curvature from spherical to cylindrical and to lamellar. The higher the molecular weight of PS homopolymer, the lesser its solubility in the corona. The relative volume fraction thus increases which causes the micelle transitions from spherical to nonspherical. By the same token, the core volume fraction may be raised by increasing PB block length, which induces the transition from spherical structure to either multilamellar vesicular or thickened lamellar micelles. In addition, the study demonstrated the existence of isolated spherical and collections of spherical mesophases.

3.6 Theories of Phase Structure of Homopolymers/Diblock Copolymer Blends

Since this study involves in miscibility among several polymer species, the phase behavior of a ternary system, and the interfacial characterization, theories related to these aspects are included for the prediction of the final morphology of blends containing two immiscible homopolymers and a diblock copolymer.

(a) Thermodynamic Theories of Phase Separation

A macrophase separation may be modeled by the Flory−Huggins theory, in which the driving force for the separation largely contributed from the unfavorable interaction between polymer species. The expression is

fυ φi 1 1 χ mnφi, mφi, n = ∑∑∑ln φi + χ mnφmφn − (11) kT inNi 2 mn,,2 i, m φi 24

where f is the Helmholtz free energy per unit volume, ν is a reference volume, φi is the average volume fraction of the polymer i, Ni is the number of reference volume units, χmn is the interaction parameter of the m-n polymer pair. There are several other thermodynamic theories predicting phase separation, which can predict size, shape, distance between domains, the lattice type, and the temperature the phase separation takes place. There are two fields of focus. First is the prediction of domain structure after the phase separation has taken place, such as by the ones by Helfand.8,50-56 Second is the prediction of the domain structure at the onset of the microphase separation, such as the ones by Leibler.57 Helfand and cowokers,51-56 and Meier58-61 developed theories that predict the equilibrium morphology of a neat block copolymer at which the free energy of the system is minimized. The driving force comes from the repulsive interaction related to χ. The domain growth decreases the volume to surface area ratio, decreasing the interfacial tension and so does the free energy. However, the loss of conformation due to the stretching of polymer chain and the loss of entropy from the confining junctions oppose the domain growth. The balance between these contributions will determine the equilibrium structure. The size of the microphase separated domains D, and the interdomain distance dint are predicted (eq 12-13).

a D = K1M t (12)

a dint = K 2 M t (13)

where Md is the molecular weights of the domain forming blocks and Mt is the total molecular weight of the copolymer. Meier predicts a = 0.56, while Helfand predicts a = 0.67.

25 Helfand and Tagami8,51 also developed a theory predicting the interface thickness between two immiscible homopolymers. The interfacial thickness Δl is modeled by

2 2 1/ 2 Δl = 2[(β A + βB )/2χ12 ] (14)

2 2 2 2 with β i = (ρoibi ) / 6 and bi = Ri / Z i where bi is Kuhn statistical segment length, Z i is the degree of polymerization, ρ oi is the density of pure polymer i and Ri is the radius of gyration of i . Equation 14 reduces to eq 15 for the case of identical chains and lattice size, where Δl∞ is the interfacial thickness at the strong segregation limit having Mw → ∞.

2b Δl∞ = 1 / 2 (15) (6χ AB )

For the interfacial thickness between the blocks in diblock copolymers, Helfand and Wasserman53-55 used the narrow interphase approximation to calculate the interfacial thickness Δl between the blocks in block copolymers. A theory based on confined chain statistics predicted that Δl should be a decreasing function of copolymer degree of polymerization Z c . Thus, eq 14 can be modified for the systems with χ AB ≥ 20 as expressed by eq 16. The difference between eq 14 and eq 16 have been found to be 40% in the low χ AB Z c region.

−1 / 2 Δl = Δl∞ [1 − (8 ln 2) / χ AB Z c ] (16)

For a typical value of χ for polystyrene-block-polydiene copolymers at 100 °C, the interface width predicted by the theory is 1.5 nm.42 The interface width determined from SAXS62-64 and SANS65-68 experiments fall in the range of 1.0-2.0 nm, which are in a close agreement with the Helfand-Wasserman theory.

26 Helfand and Tagami also derived the expression for the interfacial tension γ for a blend of homopolymers.

1 / 2 γ = (χ AB / 6) bρ 0i k BT (17)

Broseta et al.50 extended the Helfand−Tagami theory to the finite molecular weight polymers. They concluded that the interface is thicker and the interfacial tension is smaller than predicted by the theory, which assumed the infinite molecular weights. They found that in polydisperse systems, small chains tend to accumulate at the interface in order to lower the interfacial tension. From eqs 15 and 17, one can conclude that the smaller the χ, the larger the thickness and the smaller interfacial tension. Although, the theory was developed for immiscible polymers with relatively large χ, the trend suggests the potential use of block copolymers with small or negative χ as a compatibilization agent. Note that theories developed by Helfand and coworkers were derived based on the assumption that the interface region is small compared to the size of the domains. This assumption is invalid near the phase transition where all the polymer chains are extensively mixed. This leads to the second concept, which approaches from the random phase approximation. Leibler57 developed a theory that phase separation is induced by composition fluctuation. The thermodynamic properties of block copolymers in disordered state were studied by applying the random phase approximation. Leibler developed a relation between the segmental density correlation function and the scattering vector. The stability limit, or spinodal points, can be obtained from incompressible random phase approximation

27 (RPA) which predicts spinodal decomposition of polymer mixtures from a homogeneous phase. The RPA states that the coherent scattering profile may be modeled as

I(q) = BT S(q)B (18) where B is a column-vector describing the contrast and S(q) is the 3 x 3 structure factor matrix. The elements of B are related to the scattering length density of each component j.

2 −1 (bA − bB ) ⎡ 1 1 ⎤ I(q) = ⎢ + − 2χ AB ⎥ (19) υ ⎣ N Aφ A PA (q) N BφB PB (q) ⎦

where the Debye function, Pm (q ) , is defined by

2 Pm (q) = 2 []exp(−xm ) + xm − 1 (m=A, B, C) (20) xm

2 2 2 2 where xm = q Rg, m = q N mlm / 6 .

Leibler predicts that the mesophase structure of a neat block copolymer is highly dependent on χN and the volume fraction of the constituent (f).

(b) Models Predicting the Localization of the Diblock Copolymer Emulsifiers

Noolandi and coworkers69-72 proposed a theoretical model derived from the statistical thermodynamic mean field approximation to calculate the interfacial properties such as the interfacial tension and the width of the interface in A/B/(X-block-Y) ternary blend. The authors expressed the free energy density in terms of spatial volume fractions,

φK (r) , and effective mean field interaction potentials, ωK (r) .

28 F (r ) 1 1 2 = ∑ χ KK 'φ K (r )φ K ' (r ) − ∑ χ KK 'σ KK '∇φ K (r )∇φ K ' (r ) k B T 2 KK ' 12 KK ' (21) N ⎛ N ⎞ K ⎜ K ⎟ − ∑∑φ K (r )ω K (r) + log ⎜ ⎟ KKZ K ⎝ Z K Q K ()[]ω K ⎠

where F(r) is the free energy, χKK' are the Flory−Huggins interaction parameters of all

2 components, σ KK ' are the ranges of the interactions (of order of the statistical segments),

Z K is the degree of polymerization of component, K , N K is the number of polymers of type K, and Q K denotes the conformation of the chains.

They proposed a calculation of the interfacial tension when an A-block-B copolymer is present in an A/B binary blend by

⎡φ (x) ⎛φ (x) ⎞⎤ 1 1 γ = γ + dx C ln⎜ C ⎟ − φ (x) − φ − χφ φ (x) − φ (22) 0 ∫ ⎢ ⎜ ⎟⎥ ()C C P ()C C ⎣ Z C ⎝ φC ⎠⎦ Z C 2 where γ is the interfacial tension of the ternary or quaternary block copolymer- homopolymer (with or without solvent) system, γ 0 is the interfacial tension of the homopolymers without block copolymer present, φC( x ) is the local volume fraction of block copolymer at the interface, φC is the overall volume fraction of block copolymer in the system, φP is the overall volume fraction of homopolymer A or B (φA = φB = φP ), χ is the Flory−Huggins interaction parameter, and ZC is the degree of polymerization of the block copolymer, in which the block copolymer is assumed to be symmetric for simplicity. Referring to eq 22, the first term under the integral represents the loss of entropy due to the localization of the large molecules at the interface, the second term is a contribution from the chemical potential, and the last term represents the gain in energy of preferred interaction resulting from the localization at the interface of the block A in homopolymer A, and block B in homopolymer B. Noolandi and coworkers concluded

29 that the loss of entropy is the counterbalancing factor that limits the accumulation of the block copolymer at the interface. According to them, the reduction of interfacial tension can be expressed by

1 Δγ ≈ − Z φ χ 2φ 2 (23) 8 C C P

for Z C χφP << 1. Equation 23 suggests that the interfacial tension reduction mainly arises from the preferred orientation of the blocks into their identical homopolymer at the interface. Another approach of the interfacial tension deduction due to the emulsification effect was proposed by Leibler. Leibler’s theory63 is based on the free energy which includes the contributions of elastic entropy of each block, the entropy of mixing of the block copolymer with the binary blend of homopolymers, and the translational entropy of the copolymers.

Ffilm = γ 0 ⋅ A + Q[g A (∑) + g B (∑)] (24)

⎛ ∂F ⎞ γ = ⎜ film ⎟ = γ − π − π (25) ⎜ ∂A ⎟ 0 A B ⎝ ⎠Q

where π i = − ∂gi ∂ ∑ is the film (Langmuir) pressure.

However, Noolandi72 argued that Leibler’s assumption was incorrect, because it neglected to include the preferred-orientation effect into the interfacial tension reduction

( Δγ = γ − γ 0 ). Noolandi further suggested that the entropy of mixing is not significant compared to the orientation effects; therefore, eq 24 can be corrected by adding the orientation energy of the blocks and the entropy of localization for the block copolymer molecules, but the same result may be obtained by replacing γ 0 in eq 24 with the expression for the interfacial tension with added block copolymer given by eq 22.

30 Dudowicz and Freed74 have developed the lattice cluster theory (LCT) to predict the phase behavior of an A/B blend in the presence of A-block-B, where none of the constituents forms specific interactions. The theory predicts that the phase diagram of an A/B blend exhibiting UCST was shifted toward lower temperatures with increasing amounts of A-block-B, although the decrease was not large. They concluded that the miscibility of an A/B blend is enhanced from the increase of miscibility window, suggesting that A-block-B copolymer acts as an effective compatibilizer. From the theory, the critical composition can also move depending upon the volume fraction (f) of one block in the block copolymer. An experiment by Kim et al.75 suggested that the LCT could not predict the phase behavior of the blends that had lower critical solution temperature (LCST). The LCT predicts that the TLCST decreases with increasing amount of symmetrical block copolymer. According to the experiment for PS/Poly(n-pentyl methacrylate)/(PS-block-

PnPMA), the authors found that the turbidity temperature (Tb) of the LCST does not change with blend composition when the volume fraction of the block copolymer is larger than the critical amount depending on the molecular weight. The results of this study indicate that a symmetric block copolymer does not act as an effective compatibilizer for an off-critical blend composition.

31 CHAPTER IV EXPERIMENTAL METHODS

4.1 Materials

In this study, we have chosen to compatibilize the immiscible PS/PHEMA binary blends using S2VP diblock copolymer as a compatibilizer. For the study we synthesized, via anionic polymerization, a series of polystyrenes (PS) and a polystyrene-block-poly(2- vinylpyridine) (S2VP) copolymer using the standard methods.76,77 For the synthesis, cyclohexane or THF was used as the solvent, and sec-butyllithium as initiator. Poly(2- hydroxyethyl methacrylate) (PHEMA) was supplied and characterized by Aldrich Chemical Company. Table 2 gives a summary of the molecular characteristics of the polymers synthesized, in which the molecular weights were determined using gel permeation chromatograph (GPC) using a series of polystyrenes as standards. The chromatography was performed using tetrahydrofuran (THF) as a solvent, except PHEMA which is insoluble in THF. Owing to the anionic polymerization technique, the polystyrenes and the block copolymer having very narrow molecular weight distributions could be obtained. The TODT of the S2VP diblock copolymer synthesized was determined to be 150 ºC from oscillatory shear measurements.

32 Table 2. Summary of the molecular characteristics of the polymers investigated in this study characterized by GPC method in THF solvent.

Note that PHEMA is insoluble in THF; its Mv is characterized by Aldrich.

Sample Code Mw Mw/Mn

PS20 22,637 1.06

PS10 12,033 1.04

PS5 5,150 1.11

PS3 2,920 1.07

PS1.5 1,600 1.06

P2VP20 19,600 1.05

P2VP3 3,000 1.12

PHEMA 20,000 (Mv)

S2VP 16,500 1.04

4.2 Sample Preparation

Solution Blending Samples of all blends were prepared by solution casting from dimethylformamide (DMF), unless noted otherwise. All blends contain 0.2 wt % of Irganox 1010 (Ciba- Geigy Company) as an antioxidant. The solution-blended mixtures were first dried at room temperature in a fume hood and then in a vacuum oven by raising the temperature very slowly up to 180 ºC until all solvent was evaporated.

Melt Blending The polymers were mixed in a mini twin-screw micro compounder at 70 rpm for 10 min. All blends contain 0.2 wt % of Irganox 1010 as an antioxidant. The samples were 33 prepared by melt blending at 130 °C or 180 °C, and they are summarized in Table 3. A sample code 10-5-130, for instance, refers to a ternary blend with Mw,PS = 10,000 containing 5 wt % S2VP, which was melt-blended at 130 °C. All of the melt-blended samples for TEM investigation were prepared with a fixed blend composition of PS/PHEMA of 70/30 by weight.

Table 3. List of blend samples prepared by melt-blending investigated in this study.

Sample code T Polymers Composition 10-5-130 130 °C PS10/PHEMA/S2VP 70/30/5phr 10-0-180 180 °C PS10 PHEMA/S2VP 70/30/0phr 10-5-180 180 °C PS10 PHEMA/S2VP 70/30/5phr

20-0-180 180 °C PS20 PHEMA/S2VP 70/30/0phr 20-5-180 180 °C PS20 PHEMA/S2VP 70/30/5phr

4.3 Characterizations

The glass transition temperatures of the homopolymers and the blends prepared were measured using a Perkin Elmer differential scanning calorimetry at a scan rate 20 ºC/min. Thermal degradation characteristics of each homopolymer were measured by thermogravimetric analysis (TGA) in air at a heating rate of 10 ºC/min, as given in Figure 2. The onset temperature where two tangent lines intersect is designated to be the degradation temperature, Tdegrad. The thermal characteristics of the materials are listed in Table 4. FTIR spectroscopy was performed on a Digilab Excalibur spectrometer via a powder method. The polymer samples were dried, subjected to melt-blended at 140 °C 34 120

100

60 % Weight

40 PS20 P2VP20 PHEMA20 20

0 0 100 200 300 400 500

o Temperature ( C) Figure 2. TGA analysis of materials in air at a heating rate 10 ºC/min without antioxidant.

Table 4. List of the thermal properties of the homopolymers measured by DSC and TGA.

Materials Tg (ºC) Tdegrad (ºC) PS20 102.9 294.6

P2VP20 105.9 346.6

PHEMA 73.8 289.8

35 without the presence of solvent, and grinded into powder. The powder was dried once again in a vacuum oven overnight. At about 1-1.2 wt% concentration, the powder then mixed with dry KBr powder and packed into a pallet. The pallet sample in a pallet holder was installed in a well-insulated heater assembly. The samples were heated to a set temperature and held isothermally for 10 min before each FTIR measurement. Transmission electron microscopy (TEM) was used to investigate the morphology of the blends containing the S2VP diblock copolymer. The ternary blend samples of PS/PHEMA/S2VP were prepared by melt blending. The blending temperature and the processing conditions are the same of those listed in Table 3. An amount of each sample was separated and annealed at 130 ºC for 48 h in vacuum oven for the investigation of annealing effect. The PS/S2VP and PHEMA/S2VP blend samples were prepared from solvent casting. After the solution was completely evaporated, samples were prepared by compression molding at 180 ºC before further annealing at 180 ºC in order to ensure the homogeneity of the block copolymer. Then, samples were annealed at 120 ºC for 13 h to induce microphase separation from the diblock copolymer. Thin sections were microtomed using a diamond knife at room temperature. Sections were floated onto water before collection, and dried in vacuum oven at room temperature. The films were then exposed to iodine vapor for staining. Iodine selectively reacts with P2VP block; therefore, the dark areas observed under electron microscope is the P2VP-rich phase.

36 CHAPTER V RESULTS AND DISCUSSION

5.1 Phase Behavior of Binary Blends of Homopolymers

(a) PS/PHEMA Binary Blends

DSC thermograms, during the second heating scan, for PS20/PHEMA blends display one broad Tg near 89 − 90 ºC for all compositions, as shown in Figure 3. Since the difference in Tg between the neat components of PS/PHEMA is relatively small (less than 30 ºC), a single Tg from DSC is not sufficient to conclude that the PS/PHEMA blends are necessarily miscible. Figure 4 gives an optical micrograph of 30/70 PS20/PHEMA blend at room temperature, where 30/70 refers to the weight percent of the component polymers. It can be seen in Figure 4 that the size of the phase aggregates varies from a few microns to several hundred microns. The molecular weight of PS was varied (Mw = 1,500, 3,000, 5,000, and 10,000) in order to investigate the tendency of miscibility within the temperature ranging from 23 to 245 ºC. Figure 5a gives an optical micrograph of 20/80 PS20/PHEMA blend at 245 ºC, showing that the aggregates having various sizes (up to several microns) are formed. Figure 5b is an electron micrograph showing the sharp interface between the PS-rich and PHEMA-rich phases. Thus, we can conclude from Figures 4 and 5 that the PS/PHEMA blends are immiscible. Figure 6 show a schematic describing an arrangement of the components in a PS/PHEMA blend.

110 103 91 PS20

99 108 70/30 PS20/PHEMA 84

95 103 50/50 PS20/PHEMA 82

101 74 88 Endotherm 30/70 PS20/PHEMA

94 60 74 PHEMA

0 20 40 60 80 100 120 140 160 180 Temperature (oC)

Figure 3. DSC thermograms of PS20/PHEMA binary blends during the second heating scan. The tick marks indicate the values of the initial points, midpoints, and the final points of the glass transition.

PHEMA

PS20

Figure 4. Optical micrograph of 30/70 PS20/PHEMA blend at 23 – 245 °C. There was no change of morphology in such temperature range.

(a)

(b)

Figure 5. (a) Optical micrograph of 20/80 PS1.5/PHEMA blend at 245 ºC, and (b) TEM micrograph of 70/30 PS20/PHEMA blend prepared by solvent casting.

PHEMA

PSh

PHEMA

PSh

Figure 6. A schematic representing a formation of an aggregate in the binary blend of PS and PHEMA homopolymers, when PHEMA is a minor component.

(b) PS/P2VP Binary Blends

Since the Tg′s of both PS and P2VP are almost identical, the thermal analysis was found to be not useful to determine miscibility. We found that PS20/P2VP20 blends were immiscible, via optical microscopy, over the entire range of blend compositions.

Figure 7 gives an optical micrograph of 50/50 PS/P2VP blend at 245 °C. We found that at very lower molecular weights, the PS/P2VP blends were miscible. A turbidity curve was constructed for the PS3/P2VP3 blend system for the molecular weights of both components having 3,000.

P2VP20

PS20

Figure 7. Optical micrograph of 50/50 PS20/P2VP20 blend at 245 ºC.

41 Figure 8 shows that at an equal molecular weight of 3,000, the cloud point at 50 wt % PS3 is as high as 233 ºC. Since each block component in S2VP diblock copolymer used in this research has equal molecular weight of 8,300, which is two times higher than the molecular weight of the binary blend and the binary blend system exhibits the upper critical solution temperature UCST, it can be concluded that PS and P2VP block in the S2VP are immiscible within the range of the temperatures investigation 200 ºC.

280

260 Homogeneous

240 233 221 227 220 198 197 200 190

180 2 phases Temperature (ºC) 160

140

120

100 0 102030405060708090100 wt% of PS3

Figure 8. The turbidity curve of PS3/P2VP3 blend system observed under 500x using an optical microscope.

42 (c) P2VP/PHEMA Binary Blends

We found that the P2VP20/PHEMA blend was transparent under an optical microscope. DSC thermograms also indicated the blends had a single Tg for all compositions (Figure 9a). Figure 9b displays the Tg of their blends as a function of weight fraction of P2VP20. It is seen in Figure 8 that Tg′s of the binary blend exhibited a positive deviation from the weight-average values predicted by the Fox equation. The behavior suggests a possible formation of a hydrogen-bonded complex between the components indicating that P2VP/PHEMA blends have specific interactions (see eq 10). Infrared spectroscopy was used to investigate the molecular interaction between the PHEMA and P2VP pair. The full spectra were included in the Appendix section. Figure 10 gives FTIR spectra at 27 ºC of P2VP20/PHEMA binary blends at various compositions. The neat PHEMA and the blends contained PHEMA exhibited a broad absorption peak at wavenumbers between 3100 − 3650 cm−1 and a slight shoulder at 3275 cm−1. The broad peak in the wavenumber range corresponds to a characteristic band of stretching hydrogen bond. The spectrum of neat P2VP shows an absence of peak in such a band. Upon mixing PHEMA and P2VP, the maximum of absorbance in this band is shifted to lower wavenumbers. Such absorption may correspond to the increase in the amount of N-OH stretching, which overlaps the hydroxyl stretching band.

(a)

P2VP20

70/30 P2VP20/PHEMA

50/50 P2VP/PHEMA

30/70 P2VP20/PHEMA Endotherm

PHEMA

20 40 60 80 100 120 140 160 Temperature (oC)

(b)

Figure 9. Thermal analysis results of P2VP20/PHEMA blends. (a) Thermograms of the blends, (b) Plot of midpoint Tg as a function of weight composition of P2VP20.

44 The IR spectra of all samples show a similar trend at temperatures under 100 ºC; there is no shift in absorption peak upon increasing temperature, and only the absorption intensity deceases in the absorption band of hydrogen bonding. The observed decrease in absorption intensity is a characteristic behavior of hydrogen bonding; that is the length of hydrogen bonds increases with increasing temperature. However, at 140 ºC and 180 ºC, as can be seen in Figure 11, there is a significant peak shift in the absorption band to a higher wavenumber near 3550 cm−1. Therefore, it can be concluded that the amount of hydrogen bonding was significantly reduced at this temperature. Taking into account the chemical structure of PHEMA, self-association may occur especially by hydroxyl-carbonyl or hydroxyl-hydroxyl hydrogen bonding. Figure 12 shows the carbonyl stretching region for PHEMA and its blends with P2VP. For pure PHEMA, a maximum is observed at 1728 cm-1 with a slight shoulder at 1700 cm-1. The shoulder at a lower wavenumber corresponds to the carbonyl groups hydrogen bonded to the hydroxyl groups. In the presence of P2VP, the shoulder is diminishing and the absorption peak becomes narrower. The self-associated carbonyl-hydroxyl hydrogen bonding becomes less important with the presence of nitrogen in P2VP. This observation indicates the presence of intermolecular specific interactions between P2VP and PHEMA.

100 %wt P2VP 3436 3438 Absorbance

3436 70%wt 3448 50%wt

30%wt

0%wt 3800 3600 3400 3200 3000 2800 2600 Wavenumber

Figure 10. IR spectra between 3800 – 2600 cm-1 of P2VP/PHEMA blends measured at 27 °C. Samples prepared by powder method.

100 %wt P2VP

3546 3537 70%wt 3511 Absorbance 50%wt 3546

30%wt

0%wt

3800 3600 3400 3200 3000 2800 2600 Wavenumber

. Figure 11. IR spectra between 3800 – 2600 cm−1 of P2VP/PHEMA

blends measured at 180 °C. Samples prepared by powder method.

47 1728 1700 Absorbance

0%wt

50%wt 30%wt 70%wt 100 %wt P2VP

1780 1760 1740 1720 1700 1680 1660 Wavenumber

Figure 12. IR spectra showing carbonyl stretching region between 1790 – 1650 cm−1 of P2VP/PHEMA blends measured at 27 °C. Samples prepared by powder method.

5.2 Morphology of Binary Blends of Homopolymer and Block Copolymer

(a) PS/S2VP Binary Blends

An investigation under an optical microscope of the solvent-cast PS20/S2VP blends initially revealed a bicontinuous structure of macrophase separation (Figure 13). However, the as-cast film was not an equilibrium structure. After the as-dried thin films were heated to 180 °C which was above the Tg's of all the constituents and above the

TODT of the S2VP diblock copolymer, the films became clear and maintained clear even

Figure 13. An optical micrograph showing the bicontinuous structure for the blend of PS20 and S2VP block copolymer at the composition 30/70 after being dried. The micrograph was taken at room temperature. The dark area represents S2VP block copolymer, and the white area represents PS20.

49 upon cooling to room temperature. The films were prepared with THF solvent and dried in vacuum oven at 110 °C (near the Tg's of all the constituents) for several days to completely remove the solvent. The bicontinuous structure similar to this solvent-cast PS/S2VP film was observed by Han et al.48. However, their dimension of the morphology was in the order of 10 nm, compared to a few microns observed in this study. The difference in scale may arise from the fact that their block copolymer was homogeneous at room temperature. Han et al. concluded that the bicontinuous morphology evolved from frozen composition fluctuations in the disordered phase near order-disorder transition. However, from this observation, the bicontinuous morphology seems to evolve from the frozen of the copolymer micelles due to the effect of the slightly selective solvent. The frozen bicontinuous structure disappears upon heating to the temperature higher than both the glass transition temperatures and the TODT. That means the rearrangement of morphology requires the disordered state and enough free volume and thermal energy to allow the molecular segments to interdiffuse and approach the equilibrium morphology observed in Figure 14. Since the binary blends of PS20 and S2VP looked transparent under an optical microscope upon heating above 180 °C, a further investigation under TEM was conducted. Figure 14 gives a TEM image of 30/70 PS20/S2VP blend and Figure 15 gives a TEM image of 70/30 PS20/S2VP blend, each showing lamellar microdomains of S2VP diblock copolymer. The interlamellar domain distance of the 30/70 PS20/S2VP blend is estimated to be 20−22 nm, while that of the 70/30 PS20/S2VP blend is estimated to be 36−100 nm, compared with the domain distance of 12.5 nm for neat S2VP diblock copolymer having lamellar microdomains as shown in Figure 16. In this blend system, the morphology evolved from the preferential enthalpy. The relative interaction between each polymer species predetermines the stable morphology.

The thickening of PS-rich lamellae suggests that χ PS20/P2VP-block > 50 χPS-block/P2VP-block > χPS20/PS-block. The PS20 homopolymer has a much weaker repulsive interaction to the PS block than to the P2VP block. Thus, the PS20 homopolymer prefers to diffuse into the lamellar microphase and situates next to the PS-rich microphase, which resulted in an expanding domain distance of the lamellae. Figure 17 displays a schematic illustration of the molecular arrangement. This behavior has been observed78 when the molecular weight of the PS homopolymer is low enough and the volume fraction of the polystyrene in the blend does not exceed a critical concentration before the PS homopolymer coalesces forming its discrete macrophase.

Figure 14. TEM image of 30/70 PS20/S2VP blend prepared by solution blending. The dark area represents P2VP block component, and the white area represents polystyrene.

Figure 15. TEM image of 70/30 PS20/S2VP blend prepared by solution blending from DMF. The dark area represents P2VP block component, and the white area represents polystyrene.

Figure 16. TEM image of neat S2VP showing lamellar microdomains have the interlamellar domain distance of about 12.5 nm. The dark area represents P2VP block component, and the white area represents polystyrene.

PShPSc P2VPc PSc PSh PSc P2VPc

Figure 17. The schematic describing molecular arrangement in PS20/S2VP

blends of our investigation.

(b) PHEMA/S2VP Binary Systems

The samples of 30/70 PHEMA/S2VP and 70/30 PHEMA/S2VP blends looked turbid, and TEM micrographs showed evidence of a coexistence of macrophase separated PHEMA-rich phase and the microphase-separated microdomain of the diblock copolymer. Figure 18 gives a TEM image of 70/30 PHEMA/S2VP blend, showing that the minor component, S2VP diblock copolymer, forms the dispersed phase. Figure 19 gives a TEM image of the same blend, showing a diffuse interphase. This observation indicates some miscibility between the P2VP and the PHEMA at the interface. The core of the block copolymer appears to have lamellar structure. The observed microstructure of the diblock copolymer near the interface appears to be spherical micellar, which indicates that some mixing occurred at the interface between the P2VP and PHEMA. Figure 20 gives a TEM image of 30/70 PHEMA/S2VP blend, in which microphase- separated S2VP diblock copolymer forms the matrix phase in which the minor 53 component PHEMA forms the discrete phase. It is seen in Figure 20 that the lamellar microdomain structure is interrupted by the presence of the discrete phase of PHEMA. Figure 21 gives an enlarged section of the 30/70 PHEMA/S2VP blend, showing clearly the presence of a diffuse interphase, similar to that observed in Figure 19. Again, this observation indicates the presence of specific interactions between the PHEMA and S2VP block copolymer. We conclude that the specific interactions between the two components are attributable to the formation of hydrogen bonds between the hydroxyl groups in PHEMA and the pyridine groups in P2VP block of S2VP diblock copolymer.

PHEMA S2VP

Figure 18. TEM image 70/30 PHEMA/S2VP blend.

54 S2VP

PHEMA

Figure 19. TEM image showing the interface region of 70/30 PHEMA/S2VP blend. The dark phase represents the P2VP block component in S2VP diblock copolymer.

Figure 20. TEM images of 30/70 PHEMA/S2VP blend. The dark phase represents the P2VP block component in S2VP diblock copolymer. The white area represents PHEMA droplets. 55

Figure 21. Enlarged TEM image of 30/70 PHEMA/ S2VP blend. The dark phase represents the P2VP block component in S2VP diblock copolymer. The white area represents

PHEMA droplets.

Table 5 gives a comparison of the interlamellar domain distance of neat S2VP diblock copolymer with the 30/70 PHEMA/S2VP blend, showing that that the domain distance has a negligible increase from 12.5 nm to 14 nm when 30 wt % PHEMA was mixed with S2VP diblock copolymer. The strong tendency of self-association due to the OH-OH hydrogen bonding in PHEMA seems to encourage the clustering of PHEMA rather than driving the PHEMA molecules in between the lamellar microdomains of the S2VP copolymer. This results in a much smaller expansion of the lamellar spacing when compared with the PS20/S2VP blends mentioned in the previous section. The relative attractive interactions play a role in molecular arrangements. The interactions between PHEMA and P2VP block appear to have a short range effect. 56

Table 5. The characteristic interlamellar domain distance of neat S2VP diblock copolymer and the S2VP diblock copolymer in 30/70 PHEMA/S2VP blend.

PHEMA (wt %) S2VP (wt %) Interlamellar Domain Distance (nm)

0 100 12.5

30 70 14

70 30 spherical microdomain/micelle

57 5.3 Morphology of PS/PHEMA/S2VP Ternary Blend System

5.3.1 Macroscopic Observations

Initial investigation of the morphology of the PS20/PHEMA/S2VP ternary blend was via optical microscope. Two ternary blends contained 5 phr of S2VP block copolymer in 30 wt % and 70 wt % of PS20 in PHEMA, which were prepared by solvent casting, are shown in Figure 22. One may observe a kind of corona surrounding dispersed phase for both compositions. It was not clear from here what happened at the interface. At a glance, the size of the aggregates differs tremendously between the two compositions. When PS20 is a minor component, the size of the dispersed phase is about

5 μm. Inversely, the size of the PHEMA-rich dispersed phase is about 100−200 μm. This size difference largely arises from the interactions among polymer species which contributed to the surface-to-volume free energy during phase growth. Despite the size difference between the two ternary blend, their phase sizes are much smaller than the PS20/PHEMA binary blend (Figure 5b), which has an average dimension about 1 mm. The major size reduction upon an addition of S2VP copolymer is one of the evidence indicating that the interfacial tension is reduced significantly from the binary blend of the homopolymers. The morphology development of the ternary blend upon increasing temperature from 23 °C to 245 °C was investigated, shown in Figure 23. At 135 °C, which was above the Tg's but below the TODT of the copolymer, the round interfacial boundary observed at low temperature appeared to deform and modulate. The width of the corona did not change in dimension. At this temperature, the polymer segments become mobile. However, the ordered microdomains of block copolymer hinder much of the diffusion.

At 155 °C, the temperature slightly above the TODT, the corona increased in width and the

(a) (b) PHEMA

Figure 22. Optical micrographs of ternary blends consisted of PS20/PHEMA/S2VP taken at 23 °C prepared by solvent casting at compositions: (a) 30/70/5phr, (b) 70/30/5phr. Noted the 10-fold difference in scaling.

phase boundary gradually vanished, though not completely. No further morphology development is observed upon increasing temperature to 245 °C, or upon cooling from 245 °C to room temperature. The morphology shown in Figure 23c seemed to be in the equilibrium state. This vanishing of the interface indicated that the compatibility was enhanced upon increasing temperature. Compared to the case of PS/PHEMA binary blends, there was no transition and no reduction in aggregate size at all upon increasing temperature in the binary blends of homopolymers even when the molecular weight of PS was 1,500. Further investigation of the microstructure was needed to understand the effect of the presence of the S2VP copolymer in two immiscible homopolymers.

59 (a) PHEMA

PS S2VP

(b)

(c)

Figure 23. Morphology evolution upon increasing temperature observed via an optical microscope of a 70/30/5phr PS20/PHEMA/S2VP ternary blend prepared by solvent casting. (a) as-dried sample at 23 °C, (b) 135 °C, (c)

155−245 °C.

60 5.3.2 Morphology Observation via TEM

(a) TEM Images

The morphology of 70/30/5phr PS/PHEMA/S2VP ternary blend was further investigated by TEM. Two molecular weights of PS having Mw = 10,000 and 20,000 were blended having 70 wt % PS in 30 wt % PHEMA, with an addition of 5 phr of the S2VP copolymer. The samples were prepared by melt blending at 130 °C or 180 °C.

The sample code “bef-10-5-130” refers to a ternary blend with Mw,PS = 10,000 containing 5 phr of S2VP, which was melt-blended at 130 °C, with a fixed blend composition of PS/PHEMA of 70/30 by weight. The “bef-” or “aft-” prefixes of the sample denoting whether the TEM images were taken before or after the blend sample were subjected to 48 h of annealing at 130 °C. Figure 24 gives a TEM image of bef-20-5-180 and Figure 25 gives a TEM image bef-10-5-180, prepared at the same thermal condition, but only differed in molecular weight of PS homopolymer. Figure 26 gives a TEM image of bef-10-5-130, blended at temperature below the TODT of the S2VP copolymer. Figure 27 gives a TEM image of a of sample 10-5-130 and Figure 28 gives a TEM image of sample 10-5-180 after annealing at 180 °C for 48 h. An image analysis program was deployed to estimate the average size of the discrete PHEMA-rich phase, the size of block copolymer micelles formed in PHEMA-rich phase, and the thickness of the S2VP accumulated at the PHEMA-rich phase (excluded the region containing micelles). The results are shown in Table 6. Note that the standard deviations for these dimensions were quite large despite many data points were collected. Thus, the analysis would serve as a qualitative interpretation.

61 PHEMA

S2VP

PS20

Figure 24. TEM images 70/30/5phr PS20/PHEMA/S2VP melt-blended at 180 °C for 10 min (bef-20-5-180).

S2VP PS10

PHEMA

Figure 25. TEM images 70/30/5phr PS10/PHEMA/S2VP melt-blended at 180 °C for 10 min (bef-10-5-180).

62 S2VP

PHEMA

PS10

Figure 26. TEM images 70/30/5phr PS10/PHEMA/S2VP melt-blended at 130 °C for 10 min (bef-10-5-130).

PHEMA

S2VP PS10

Figure 27. TEM images 70/30/5phr PS10/PHEMA/S2VP melt-blended at 130 °C for 10 min, and simultaneously annealed at 130 °C for 48 h (aft-10-5-130). 63

PS10

S2VP

PHEMA

Figure 28. TEM images 70/30/5phr PS10/PHEMA/S2VP melt-blended at 180 °C for 10 min, and simultaneously annealed at 130 °C for 48 h (aft-10-5-180).

Table 6. List of dimensions: average domain size of PHEMA-rich phase, the average micelle size of the S2VP diblock copolymer, and average thickness of S2VP accumulated on the interfacial region along PHEMA-rich phase measured from TEM images via an image analysis program.

Avg Domain Size Avg Micelle Size Avg thickness of interphase Sample code (nm) (nm) (nm)

bef-10-5-130 254 ± 98 15 ± 9

aft-10-5-130 295 ± 200 26.8 ± 8 22 ± 18

bef-10-5-180 337 ± 222 36 ± 18

aft-10-5-180 540 ± 335 29.3 ± 11 11 ± 6

bef-20-5-180 460 ± 292 39 ± 24

(b) Effect of Molecular Weight of PS to the Compatibility

Having exposed to the same thermal condition, samples bef-10-5-180 and bef-20-5- 180 only differs in the molecular weights of the PS homopolymer. The PHEMA domain size was 337 nm for PS10, and 460 nm for PS20, respectively. The dimensions differ by 64 25%. The fact that the two samples had the same composition of PS, PHEMA, and S2VP, points out that the volume fractions of the PS in ternary blend system were equal among the two samples. Only the χNPS was different. For the two-fold increase of molecular weight, the phase-separated size will increase by ¼. The size increase indicated the reduction in compatibility with the higher molecular weight due to the increase in χN. The variation of the molecular weight did not influence the interphase thickness.

The introduction of a block copolymer as a compatibilizer into two immiscible homopolymers complicates the morphology of the resultant blend. This is because the block copolymer has unique characteristics, such as microphase separation. A block copolymer will have a homogeneous morphology at temperatures above its TODT. At temperatures below TODT, the block polymer will spontaneously phase-separate into a spatially order morphology at thermodynamic equilibrium.

The block copolymer S2VP has the TODT of 150 °C, but the strength of specific interactions reduces dramatically at temperature higher than 140 °C. Since the copolymer was disordered during blending, the interphase appeared to be less sharp for 180 °C when compared to the sample blended at 130 °C. When compared the interphase thickness, the diblock copolymer accumulated around the PHEMA-rich phase for the sample blended at 180 °C was twice thicker than the one blended at 130 °C, i.e. the thickness of 36 nm versus 15 nm respectively. Although, we assumed that the block copolymer was in an ordered state during blending at 130 °C, the morphology of the accumulated block copolymer did not appear to have an ordered microstructure and the wetting was completely covered the entire macrophase. Therefore, it suggested that the block copolymer was partially in the disordered state during blending. It was possible 65 that heat generated during blending might cause the surge of the melt temperature to near 150 °C. Thus, the complete compatibilization of the S2VP occurred on the interfacial region. However, the partially ordered microdomains reduced the degree of freedom of the diblock copolymer, which resulted in a smaller thickness of the interphase thickness since the wetting of the ordered copolymer was less effective than the copolymer in a disordered state. The diblock copolymer, therefore, prefers to accumulate more at the interface at the blending temperature higher than the TODT. From the image analysis, the domain size of the PHEMA-rich phase was 254 nm for 10-5-130, while the size was 337 nm for the 10-5-180. The interaction parameter χ generally is inversely proportional to temperature. However, the specific interactions become weaker as temperature arises. The specific interactions were stronger at 130 °C. Thus, the resultant macrophase size was slightly smaller due to the reduction of interfacial tension between the PHEMA-phase and the S2VP copolymer accumulated around it.

(d) Effect of Annealing

After annealing at 130 °C for 48 h, some of the S2VP diblock copolymer diffused into the PHEMA-rich phase, which resulted in the swell of the PHEMA-rich phase, and formed spherical or ellipsoidal micelles inside the discrete phase, even though the S2VP copolymer should be in an ordered state at 130 °C and the annealing would induce the formation of the lamellar microstructure which was observed in other studies2,28. The micelles were similar to the results found for an amphiphilic block copolymer in a selective solvent. If there were no attractive interactions between the polymer species in the ternary blend, the added block copolymer would form micelles in the bulk phases, some would solubilize in the blend, and the rest would locate at the interface between homopolymers as an emulsifier.23 The preferential interactions from the intermolecular

66 hydrogen bonding between the PHEMA and P2VP block in the copolymer are the driving force to bring S2VP into the macrophase of PHEMA. The thermal energy from annealing enhances the diffusion. When the concentration of the block copolymer exceeds the critical micelle concentration, the micelles then form. Figure 29 shows a schematic representation of the morphological development before and after annealing. As the system always seeks to obtain the morphology that has lowest free energy, the enthalpy of mixing plays a large role to the morphology of ternary blends. The PS/PHEMA/S2VP ternary blend system consists of four constituents. The examination of the binary blend suggests that the polymer interactions vary. The intramolecular interaction in the PHEMA homopolymer seems to be the strongest, as it tends to self- associate. The second strongest attraction is PHEMA and P2VP block since the species form hydrogen bonding. The interaction between the same species, e.g. P2VP- block/P2VP-block and PS homopolymer of the same molecular weight, seems to be the next strongest attraction since their χN values are equal among the same species of the same molecular weight. The next mildly repulsive interaction pair is between PS homopolymer and PS in the block copolymer. Their small difference in χN contributed to a slight increase in free energy. The cloud point measurement of the P2VP3/PS3 blend suggests that the repulsive interaction between P2VP block in S2VP and PS homopolymer is relatively less strong than the PS/PHEMA pair, which do not have any tendency to mix. These interactions determine the arrangement of the polymer chains as the most negative or the smallest χN pair will reduce the local free energy the most. Therefore, the molecules arrange in space in the order of which species are attracted the most, followed by the species that attracted the second strongest.

(a) (b)

PHEMA

P2VPc PSc

PSh

PHEMA

S2VP PSh

Figure 29. Schematic presentation of morphological development (a) before, and (b) after annealing at 130 °C in 70/30/5phr PS10/PHEMA/S2VP.

68 CHAPTER VI CONCLUSIONS AND RECOMMENDATIONS

This thesis presents the morphology development in immiscible PS/PHEMA blends using functional S2VP diblock copolymer as a compatibilizing agent. The introduction of the S2VP diblock copolymer significantly reduced the size of the aggregates resulting from the significant reduction in interfacial tension. After melt blending PS and PHEMA for only 10 minutes, the functionalized S2VP diblock copolymer accumulated completely around the discrete macrophase, unlike the case of emulsifier block copolymer where there was no preferential interaction between polymer molecules. Even though the blending temperature may have been lower than the TODT, the S2VP diblock copolymer can form an interphase over the entire the interfacial boundary due to the presence of the specific interactions. The specific interactions help overcoming the high viscosity of polymers that may impede the diffusion rate, which limits the ability of each block component to be situated at the interphase. The present study indicates that the strong attractive interactions, via hydrogen bonding, between the P2VP block of S2VP diblock copolymer and PHEMA homopolymer has an advantage of inducing the growth of the interphase over a ternary blend in which a nonfunctional block copolymer is used as a compatibilizing agent. The interphase formation suggests the enhancement in interfacial adhesion, although mechanical testing was not implemented in this thesis because the model polymers were too brittle to prepare specimens. The interphase, however, is not thermally stable because the preferential attraction drives the block copolymer into the discrete macrophase. The processing time and 69 temperature has to be carefully chosen for the compatibilizing agent in order to obtain the thick interphase for the blend. The factors influencing the interphase formation in the PS/PHEMA/S2VP ternary blend should be further investigated by varying the block length, the block volume fraction, the concentration of S2VP diblock copolymer in the blend, the duration of isothermal annealing, and the blending temperature. The morphology evolution during melt blending of two immiscible polymer bends should be investigated further using other functional block copolymer as a compatibilizing agent. The efficiency of a functional block copolymer as compatibilizing agent would depend on the extent of attractive interactions between the homopolymer and block copolymer and also on the TODT of block copolymer.

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77 APPENDIX