Photopolymerised Urethane Acrylates

>-5.5.93

PHOTOPOLYME RIS E T) URETHANE ACRYLATES

Anthony Brian Clayton B. Sc. (Hons).

Thesis submitted for the degree of Doctor of Philosophy tn The University of Adelaide

( Faculty of Science ) Table of Contents Summary vl

Statement vt tt Acknowledgements tx

Abbreviations x

Chapter One Introduction 1

Chapter Two Experimental Techniques 7

2.1 Sample sources 7

2.2 Sample preparation 7 2.3 Sample synthesis I 2.4 Sample polymerisation l2 2.5 Dynamic Mechanical Thermal Analysis l2 2.6 Nuclear Magnetic Resonance l4 2.7 Dffferential Scanning Calorimetry 20

2.8 Sorption/Desorption Measurements 22

2.9 FTIR/NIR Spectroscopy 23

2.10 Thermogravimetric Analysis 23

Chapter Three Dynamic Mechanical Measurements 24 3.1 Introduction 24 3.2 Results and discussion 25 3.2.L Homologous series 25

(a) TET homopolymers 25

(b) TPCL homopolymers 34

ll (c) TET(Fr|MO)2 homopolymers 39

(d) Hard segment model compounds 43 3.2.2 Copolymers 53

(a) TET copolymers 53

(b) TPCL copolymers 7L

3.2.3 Water in Urethane Acrylate Networks 79 3.3 Summary 93

Chapter Four DSC Results 95 4.1 Introduction 95 4.2 Results and discussion 96 4.2.I Homopolymer Thermograms 96

(a) TET homopolymers 96

(b) TPCL homopolymers L02

(d) Hard segment model polymers 108

4.2.2 Annealing experiments tr7

4.3 Summary 118

Chapter Five NMR t20 5.1 Introduction 120 5.2 Results and discussion 120 5.2.1 l3C liquid phase NMR 120

5.2.2 SOIid StAtC PEMAS 13C NMR t20

(a)TET homopolymers L20

(b) TET2g0O/methyl acrylate copolymers r30

ut (c) TET29O0/TEGDA copolymers 133

(d ) Vãialion in TEI2ffi ad 290 NMR pææræs wih tirre 136 5.3 Summary 140

Chapter Six So¡ption and Diffusion 142 6.1 Introduction 142 6.2 Diffusion kinetics 142 6.3 Results 145

6.3.1 Homologous series t45

(a) Sorption at 33Vor.h. 145

(b) Sorptio n at 7 9Vor.h. 1s1

(d) Sorption in water. 159

6.3.2 Variation of diffusion coefficienb with water uphke L64 6.3.2 Copolymer sorption in water L66

(a) TET 650 copolymers L66

(b) TET 1000 copolymers L70

(d) TET2900 copolymers t75 6.4 Discussion 177 6.5 Summary 187

Chapter Seven FTIRNIR Measu¡ements 188 7.1 Introduction 188

7.2 Results and discussion 189

7.2.1 ATR results 189

(a) Hydrogen bonding 189

IV (b) Conversron L92

(a) Conversion t96

(b) Water uptake 198 7.3 Summary 200

Chapter Eight Conclusions 20t

References 2t0 SUMMARY

A series of urethane acrylate prepolymers containing toluene 2,4 diisocyanate/ 2- hydroxy ethyl acrylate hard segments and polytetramethylene oxide (PTMO) and poly (e- caprolactone) (PCL) soft segments of various molecular weight were synthesjsed. Polymer samples were prepared by u.v. irradiation of the initiated prepolymers.

Dynamic mechanical properties of the urethane acrylates were measured as a function of

soft segment molecular weight and type. Copolymers of the u¡ethane acrylates with lower molecular weight comonomers were also investigated.

Differential scanning calorimetry was used to measure polymer glass transition and

melting temperatures. Together with dynamic mechanical techniques, this enabled transitions

to be assigned to specif,rc molecula¡ motions.

Proton enhanced magic angle spinning l3C Ntr¿R was used to obtain the time constant

Tlp(C), the relaxation time for l3C in the rotating frame, and the spin-lock cross polarisation

time, T5¡, for ca¡bons in the PTMO soft segments. Tlp(C) increased as the polymers became

more rubbery, indicating the release of mid-kHz components of motion in the soft segments.

These techniques enabled the molecular transitions in the urethane acrylates to be

assigned solely to motions in the polymer soft segment. In this respect these lightly

crosslinked polymers were found to differ considerably from linea¡ polyurethanes, where hard

segment transitions can strongly influence polymer properties.

Polymer samples were equilibrated in closed atmospheres over saturated salt solutions

at different relative humidities. Non-Fickian sorption behaviour in some polymers at low

relative humidities was attributed to clustering of water molecules at specific sites in the

polymers.

Sorption of water into the urethane acrylate polymers produced various changes in

dynamic mechanical response from that of the dry polymer. Mechanical properties of the

VI polymers were found to be affected not by the absolute amount of water in the networks, but by the distribution of water in the polymer. Polymers in which water was not associated with specific molecular units showed ice crystal formation during thermal runs from subambient temperatures.

FIIR spectroscopy, both Attenuated Total Reflectance and Ne-ar Inira¡ed, yielded information on double bond conversion and the state of water in the polymers.

vrl Statement

This work contains no material which has been accepted for the award of any other degree or diploma in any university or other tertiary institution and, to the best of my knowledge and belief, contains no material previously published or written by another person, except where due reference had been made in the text.

I give consent to this copy of my thesis, when deposited in the University Library, being

available for photocopying and loan.

lsr*r Septembe r L992

vlll Acknowledgements

I would like to thank my supervisors, Dr. P.E.M. Allen and Dr. D.R.G. Williams for their assistance and guidance throughout the course of this work.

I would also like to acknowledge Mrs. A. Hounslow for obtaining solid state NMR spectra. Thanks must also go to several people at external institutions. Personnel in the

Polymer and Materials Science Section of Telecom Resea¡ch Laboratories, Melboume, ensured that my visits there were both productive and enjoyable. Particula¡ thanks must go to Alistai¡

Itnpey for his assistance in the operation of the DMTA and Barry Keon who performed DMTA and DSC analysis of ha¡d segment model compounds. The assistance of Dr. Ma¡k Fisher at

Sola Intemational, Adelaide, with subambient DSC is gratefully acknowledged.

The contributions of Darrell Bennett, Chee-HoongLai and Darren Miller through numerous discussions must also be acknowledged.

Finally, I would like to thank my parents for their love, suppof and encouragement throughout my postgraduate studies.

IX ABBREVIATIONS

BD buta¡re diol acp heat capacity change at Tg

AH* DSC peak melting enthalpy

ÂH. DSC peak crystallisation enthalpy

DMTA Dynamic Mechanical Thermal Analyser

DSC differential scanning calorimetry

ED ethylene diamine

EWC equilibrium water content

FTIR Fourie¡ Transform Infra¡ed S pectroscopy

HDDA 1,6 hexanediol diacrylate

HEA 2-hydroxy ethyl acrylate

HEMA 2-hydroxy ethyl methacrylate

MA methyl acrylate

MDI 4,4'-diphenylmethane diisocyanate

NIR Near Infra¡ed Specroscopy

NMR nuclea¡ magnetic resonance

PCL poly(caprolactone) diol

PEMAS proton enhanced magic-angle spinning

PTMO poly(tetramethylene oxide) diol

TDI toluene diisocyanate (2,4 isomer unless otherwise noted)

TEGDA tetra(ethylene glycol) diacrylate

glass temperature Tob transition

TET urethane acrylate with one F fMO chain between two terminal TDVggA

unlts

X TET(PTMO)2 urethane acrylate with two mMO chains connected by a single TDI group,

between two terminal TDVFIEA units

TPCL urethane acrylate with one PCL chain between two terminal TDVFIEA units

XI CHAPTBR ONB

INTRODUCTION

The initial discovery that the reaction of long chain diols with diisocyanates yielded useful polymers, known as polyurethanes, was made by Otto Bayer and coworkers at I.G. Farben in Leverkusen, Germany, in 1937 as a response to investigations into fibre-forming polyamides by Carothers at DuPont, USA 1'2. A typical linear polyurethane structure is shown in Figure 1.1. The polyurethane shown has a "hard" segment consisting of two MDI units linked by butane diol (BD). Other difunctional isocyanates (e.g. 2,4 or 2,6 TDI or isophorone diisocyanate ) can also be used in the hard segment. The BD unit linking the two MDI units is known as a "chain extender" 1,2 and enables a desired number of diisocyanate containing groups to be "built in" to the hard segment. The soft segment shown is a PTMO chain with an unspecified number of repeat units. Other long chain diols (e. g. polypropylene glycol and polycaprolactone glycol) can be reacted with diisocyanates to form polyurethanes with different soft segments. Polyurethanes may be regarded as being block copolymers consisting of alternating blocks of two dissimilar polymer chains 3-6. Electron'microscopic 7 and X-ray studies 8-10 þ¿ve shown that the hard segments of the polyurethane block cluster into separate domains within a rubbery matrix of the soft segments. The hard domains act to reinforce the rubbery segments by functioning as filler particles or pseudo-crosslinks (Figure 1.2) 3' I1,12. The morphology of these systems may vary from two phase, in which there is either a continuous hard phase with dispersed soft phase 2

H HO H

1 I c((cH2)

Figure 1.1 Typical linear polyurethane strucrure.

à) sof t seçm¡n15 \l/,/\/\ lì¿rd gegnì€nt s

b) soft / - plìase

l-srd docr¡ains

Figure L.2 Schematic representation of (a) segmented structure in a linear polyurethane and (b) two-phase structure in the bulk polymer. 3

(at low soft segment content) or a continuous soft phase with dispersed hard domains (at high soft segment content) 13. Intermediate morphology, with both ph ases continuous, is found w hen the concentrations of hard and soft segments are about eQual 14.

The properties of segmented polyurethanes can be varied across a wide range from rigid foams to rubbery elastomers. Compositional variables such as the diisocyanate structure, the chain extender 15, the molecular weight and molecular weight distribution l6 of the soft and hard segments strongly influence the extent of phase segregation and domain organisation in the polyurethane, thus affecting polymer properties. The degree of microphase segregation in segmented polyurethanes has been investigated by small angle X-ray scattering analysis (SAXS) 11-24 and by analysis of the soft microphase glass transition temperature from dynamic mechanical 25-29 and differential scanning calorimetry (DS C) measurements 30-34. Most studies have concluded that microphase separation is incomplete, with an elevation in the soft microphase Tg being found as a result of the presence of "dissolved" hard segment.

The presence of these hard domains in polyurethanes is thought to be responsible for stress softening under cyclic loading conditions. The disruption of domain structure can lead to a decrease in the number of effective' crosslinking sites 35 with consequent degradation in mechanical properties. Stripping of polyurethane from printing rollers, where the elastomeric properties of the polymer are useful in controlling noise and damping vibration 36, has been attributed to this phenomenon 35.

Commercial polyurethane coatings are frequently applied from a solvent carrier, which evaporates after the initial application. Environmental consideration s, coupled with a desire to increase the 4 speed of cure over that attainable with thermal polymerisation led to the increasing use of acrylated urethane oligomers in the late 1970's.

Typically, one mole equivalent of a polyether or polyester diol is capped at both ends with two moles of a diisocyanate such as TDI, then the remaining pendant isocyanate groups are reacted with HEA or HEMA to provide the acrylate functionality. Acrylate terminated urethanes, combining the flexibility in mechanical properties of polyurethanes with the rapidity of acrylate polymerisation by u.v. or electron beam (EB) radiation are currently used in a variety of applications including coatings for optic fibres used in telecommunications 37,38 in the printing industry 39, and in the coating of fabrics in the textile industry 40. Russian workers 41,42 initially investigated structure/property relations in urethane methacrylates as long ago as the late 1960's; extensive literature citations after 1980 coincide with the increasing industrial use of urethane acrylate based coatings. Several of these later studies 43-47 interpreted single glass transition peaks in lower molecular weight soft segment urethane acrylate polymers as being indicative of one - phase morphology, in which soft and hard segments were homogenously mixed. Dynamic mechanical work in this thesis concentrated on acrylated TDI with either PTMO or PCL soft segments with molecular weights ranging from 650 to 3000. This molecular weight range was chosen in order to compare properties of urethane acrylate polymers with one - phase morphology to two - phase networks containing higher molecular weight soft segments. In addition to the polymers with a single continuous soft segment chain between acrylate groups, novel prepolymers were synthesised in which two PTMO chains, interconnected by a single TDI group, were endcapped with two acrylated TDI groups. These prepolymers were made in order to investigate the effect of a large decrease in network 5 crosslink density, while keeping the overall hard segment content nearly con s tant Commercial coating formulations commonly contain other acrylate monomers in addition to the urethane acrylate prepolymer. These monomers (known as reactive diluents) are added to reduce the viscosity of the prepolymer in order to obtain better processability. Previous investigations into the dynamic mechanical properties of urethane acrylate copolymer networks 39 attributed modulus changes on copolymerisation to improved phase separation brought about by preferential association of the reactive diluent with the hard segments in the urethane acrylate. Dynamic mechanical properties of urethane acrylates with three acrylate diluent monomers - methyl acrylate, TEGDA and HDDA were determined in order to compare the effects of the addition of crosslinking and non-crosslinking monomers.

In comparison to sorption studies on polyurethanes 48-55' there have been relatively few investigations into the effects of water sorption in urethane acrylates 56. In this thesis, urethane acrylate homopolymers were sorbed in humid atmospheres and in water in order to determine the total water uptake and the sorption kinetics of water in these systems. The water permeability of these materials is an important consideration when they are used as coatings for moisture-sensitive substrates. Tensile tests on glass optic fibres have correlated 57 fibre weakening to stress induced reactions between silica and absorbed water at the glass-polymer interface.

Urethane acrylate copolymers with methyl acrylate, TEGDA and HDDA, \rere sorbed in water in order to assess the effect of copolymerisation on water sorption into the network. 6

The effect of water sorption on the dynamic mechanical properties of the urethane acrylates was also studied. Solid State PEMAS l3C t¡tvtR has been applied to the study of polycarbo¡¿¡s 58-60 and epoxj slstem5 ól with, in favourable cases, changes in polymer mechanical properties being correlated with changes in NMR time constants which reflect near static motions and motions in the mid-kHz range for the carbons under observation.

NMR pulse sequences were used to determine these time constants for a series of PTMO based urethane acrylate polymers. Other workers 62-64 have used biexponential decay in carbon magnetisation in polyurethanes as a probe of polymer phase separation. It was hoped that PEMAS 13C NtrrtR techniques would permit examination of similar phase structure in the urethane acrylates. 7 CHAPTBR TWO

BXPBRIMBNTAL TECHNIQUES

2.1 Sample sources HEA, TDI, stannous octoate, PTMO diols of molecular weight 650, 1000, 2000 and 2900, PCL diols of molecular weight 1250, 2000 and 3000, and HDDA were obtained from Polysciences. The number of repeat units, n, in the diols was approximately 9, 14,20 and 28 for the PTMO 650, 1000, 2000 and 2900 respectively. The PCL 1250,2000 and 3000 diols contained 11, l8 and 26 repeat units respectively. HEMA, methyl acrylate and trimethylchlorosilane \4/ere obtained from Fluka. TEGDA was obtained from Aldrich. 2,2 dimethoxy 1,2 diphenyl ethan-1-one (Irgacure 651) photoinitiator was received courtesy of Ciba- Geigy. 2.2 Sample preparation

The PTMO and PCL diols were dehydrated prior to synthesis in a 50oC vacuum oven for at least 48 hours. TDI was purified by vacuum distillation : b.p. 60"C (0.33mm Hg) and frozen prior to use.

HEA and HEMA were d¡ied over activated 4Å molecular sieves for at least 48 hours prior to use. All other monomers, stannous octoate and trimethyl chlorosilane were used as received. 8

2.3 Sample synthesis (i) TET and TPCL prepolymers The synthesis scheme for these prepolymers is shown in Figure 2.L. HEA was added dropwise to an equimolar amount of TDI with stirring under nitrogen, with the temperature kept under 40"C throughout the reaction. v/hen the temperature started to drop, dehydrated PTMo or PCL diol (0.5 molar eq) containing 0.l5Vo w/w stannous octoate catalyst

was added and the reaction continued at 70oC for a further 2 hours. The

viscous product was recovered in about 8O7o yield and was stored in the

dark prior to use. (ii) TET(PTMO)2 prepol_vmers The procedure for synthesis of the prepolymers is shown in Figure

2.2. Initially dehydrated PTMO (2 molar eq) containing 0.15 wtVo stannous octoate was added to distilled TDI and reacted, with stirring, for 2 hours at 70oC. 2 molar equivalents of the 1:1 TDIÆIEA adducr, prepared as in (i), were then added to the reaction mixture, together with l a further 0. 15 wtTo stannous octoate. The reaction was heated for a further 2 hours at 70oC with reacrion completion being verified by IR

spectroscopy from the absence of the -NCO stretch at 2273 cm-1 in the prepolymer.

liiil Hsrd seørnent rnrìdel r.rì m nnrrnrlc Hãrd segment model compound prepolymers containing only TDI and HEA (Figure 2.3) were made according ro the following procedure. The 1:1 hard segment model compound was synrhesised by adding I mole of dried HEA dropwise to 1 mole of TDI - after the reaction temperature decreased to 25oC the product was decanted and cast within one hour. 9

=C=O (i)

=C=O o -o-cH2cH2-o-c-cH=cH2il + 30"c o il I HO-CH2CH2-O-c-cH=cHz

HO-((CHt4O),¡-H pTMO diol

or (ii) 2r + o il HO- (c-(cHtso)n -H PCL diol

0. I 57o stannous octoate 70'c

o o n lt NH-C-O- cH2cH2-o-C-CH=CH2

o Hrc lt NH- C - O -((CHt4O)n -TDr-HEA o il or (c-(cHr5o),

Figure 2.1 Reaction scheme for TET and TPCL urethane acrylate prepolymer synthesis. N refers to the numhr of repeat units in the soft link. 10

N=C=O (i) 2 Ho-((cH2)4o)n-H +

N=C=O

0.15% stannous octoate 700c

o Ho-((cH2)4o1"-[ * o il c- o-((cHr4o)n-H

o (ii) u 70"c HEA-rDI-O -(cHzlqo)- [ -¡¡',r + 0.157o stannous .octoate 2l NH ö-o

I o(cHr4o),TDr-HEA

Figure. 2.2 Reaction scheme for TET(PTMO)2 urethane acrylate prepolymer synthesis. 11

CH3 {{

NH- c-o.cH2cH2-o-c-cH{H29?

oo ll il NH-C-O -CH2CH2-O -C-CH=CH2

NH_ c-o-cH2cH2-o-c-cH{H29?

III

Figure 2.3 Ha¡d segment model compound prepolymers : (I) 1: I mole rario TDIÆ{EA and (IIÐ 1:2 mole ratio TDIÆ{EA. T2

Synthesis of the l:2 mole ratio TDI/FIEA hard segment model compound involved addition of a further 1 molar equivalent of dried HEA containing 0.15 wtVo stannous octoate catalyst to the 1:1 mole ratio compound. After the reaction had cooled to room temperature the product was decanted and cast within one hour.

2.4 Sample polymerisation Photoinitiator was added to the prepolymers, usually at 3.IVo by weight, and the mixture heated at 100- 110"C to dissolve the photoinitiator. FTIR confirmed that this heating did not cause premature gelation of the prepolymer. Polymer sheets, 2mm thick, were cast by Cowperthwaite's method I with initiated prepolymer being poured between two glass plates separated by Silastic tubing as a gasket. Prior to casting the hard segment model compounds described earlier, the glass sheets were treated with neat trimethyl chlorosilane to facilitate removal of the polymer from the mould. Afçer pouring into the cavity between the two glass sheets, the initiated prepolymer was allowed to cool to room temperature, at which point the sample was irradiated for l5 minutes with a 300W Wotan U. V. lamp.

2.5 Dynamic Mechanical Thermal Analysis After removal from moulds, dynamic mechanical thermal analysis was carried out on 2 x 8 x 40 mm samples cut from the cast sheet. A Polymer Labs DMTA MkII was used in the bending mode, with the sample clamped in the double cantilever geometry In this configuration the sample was clamped rigidly at both ends and its central point vibrated sinusoidally by the drive clamp (Figure 2.4). Samples were usually scanned between -100 and +100'C at 3K/min using a nitrogen gas purge. 13

7 Vibrator D isplacement Transducer

Temperaiure

:, E nclosure Sample t

Liqu id N itrogen

Drive Shaft Clamps

(a)

(b)

Figure 2.a @) Schematic diagram of the Polymer t-abs DMTA head used in the bending mode (b) Detail of double cantilever sample clamping geometry. t4

Samples were scanned at three frequencies :- 1, 10 and 30 Hz. Strain applied to the samples was kept to less than l7o of sample thickness in order to avoid any non-linear effects 2. For thermal scans commencing at subambient temperatures (usually from -l2O ,o -100oC upwards) rubbery samples were reclamped at low temPerature to ensure good contact through the scan. 3'4 The DMTA Analyser solves the relevant equations of motion to yield the values for the in-phase Young's modulus, E', and the out-of- phase Young's moduluS, E", from which values for the loss tangent, tan õ, were calculated.

2.6 Nuclear Magnetic Resonance 2.6(a)Solution NMR A Brüker WP80 Fourier Transform NMR spectrometer was used to obtain the l3C liquid specrra of the TET series urethane acrylate prepolymers. Prepolymers were dissolved in deuterated chloroform (ZOVo w/v) and placed in a lQmm diameter NMR tube (Wilmad 513-1PP)' All chemical shifts were obtained with reference to the central peak in the l3C CDC13 triplet at't1.0ppm. The spectrometer operated at 20.1MHz for 13C observation. Standard techniques were used for measurement [ pulse : 3.5psec, 45o,lH BB decoupling (2W); 8K data table l.

2.6(b)Solid State NMR Proton-enhanced, magic angle spinning (PE/MAS) solid state high resolution NMR can provide direct information on the dynamics of particular carbon atoms in solid polymers, and is sensitive to a wide range of motional frequencies from 0.01 to 1010 Hz. Thus solid state 15

NMR may, in favourable circumstances, be used to relate the molecular dynamics of the groups or segments of the macromolecule to its macroscopic, i.e. bulk properties. The two frequencies used in this work are described by the time constant Tlp(C), the l3C relaxation time in the rotating frame, which is sensitive to molecular motions in'the mid- kilohertz range and the spin lock crosS polarization time, TSL, which is sensitive to near static motions. Tlp(H), the lH relaxation in the rotating frame, can be derived in the course of calculating T5¡. Dipolar interactions between 13C and lg nuclei result in spectral line broadening. Chemical shift anisotropy (CSA), due to the asymmetry of the electron cloud shietding the carbon nucleus must also be removed to enable high resolution l3C spectra to be obtained. A high powered decoupling field, similar to that used in l3C liquid NMR to remove carbon-proton spin-spin coupling, was used to reduce the dipolar line broadening. Removal of the CSA can be achieved by spinning the polymer at the so-called "magic angle" (54.1"). The degree of shielding which a carbon nucleus in a molecule experiences depends on the orientation of the molecule to the magnetic field, B0 , according to the

relation 3cos2Ê-1, where B is the angle between the magnetic field and the bond axis. Substitution of the magic angle into this relation yields a value of zero for this interaction. Usually the polymer sample is spun at frequencies greater than the dispersion of chemical shifts:- either the polymer is contained inside a ceramic rotor or is itself machined into a cylinder which can be rotated at these frequencies. Cross polarisation (CP) or proton enhancement was also used to overcome two further problems associated with l3C tttr,tR. In order not to saturate the signal from l3C nuclei in the experiment a delay time of several l3C spin-lattice relaxation times must be programmed before data sampling can be repeated. This, coupled with the low natural abundance of 13C nuclei l6

(1.l1Vo compared with 99.9Vo for lH nuclei), would necessitate overly long experimental times before the resultant signal was of sufficient intensity. The CP experiment involves bringing the relatively small "hotter" l3C nuclei reservoir into contact with the larger "cooler" lH reservoir. Magnetisation of the carbons is achieved by transfer from nearby protons when the Hartmann-Hahn condition is satisfied:

YC BIC = YH Bls

where yç and yH are the l3C and lH gyromagnetic ratios and Btç and Blg are the carbon and proton field magnitudes. When this condition is met and the protons and the carbons are locked so that their energy levels are matched, energy conserving spin flips can occur between carbon and proton spins. This transfer is a spin-spin (T2) process, generally requiring no more than 100 ps 5 effectively reducing the T1 times for carbon relaxation and enabling a four fold signal enhancement under ideal conditions. The time constant describing the rate of magnetization transfer is T5¡ (spin tock) (also referred to as Tç¡¡ ) and can be determined from the following matched spin-lock, single-contact cross-polarisation procedure. Initially the proton spins are polarised in a field Bg after which they are placed in the rotating frame by a 90'pulse lH followed by a 90o phase shift and continuous application of a strong field. The third part of the experiment involves placing the l3C spins such that the Hartmann-Hahn condition is satisfied. After the contact has

been made for a known time, tçp, the glç field was then turned off with dipolar decoupling of the lH spins being maintained (Figure 2.5 a). The TSI values, which provide information on the near statlc l3C components of motion, can then be calculated from the variation in 17

Til2 90"

Þpln lock decouple

13c

t acqulre recycle cp

Figure 2.5 (a) Pulse progranìme for obtaining spin-lock, cross polarised NMR time constants, T5¡.

tú2 900 900

13c

hold

Figure 2.5 (b) Pulse programme for obtaining relaxation time constants for I3C in the rotating frame, Tlp(C). 18 signal intensity with contact time, t. The following equation 6 relates the observed intensity to r :

I = Io ¡.-1 tl - exp (#rJ exe (rli--H/f (2.1)

wherel,=1* Ttt Tsl (2.2) rl p (C) r1p (H)

and I = the carbon peak intensity after cross polarisation time t. TSL, Tlp(C) and Tlp(H) are as defined previously. Since Tlp(C) is usually much greater than T5¡ the ratio can be regarded as being =:%rl p (C) insignificant and Equation 2. 1 reduces to

I I - expclfril ( ) (2.3) I=+rexprffi;l Tsl r lp(H)

A non linear least squares regression program , DATAFT 7, was used to fit the experimental data to the above equation from which the T5¡ and Tlp(H) values were obtained. T1p(C) values were also measured using a pulse sequence previously described by Schaefer 8. Spin contact was established for a variable time, and then terminated by turning off the lH rotating field. The l3C spins are then held in their rotating field for a variable time and data collected with dipolar proton decoupling. This pulse sequence is illustrated in Figure 2.5b. Magic angle spectra ,were acquired on a Brüker cXP-300 spectrometer with a frequency of 75.41 MHz, a proton decoupling field of 6lkHz (136G), and a carbon spin-lock field of 60kHz (53G)- The recycle time was five seconds and the nlT carbon pulse duration was 4'2ps' All

croSS polarisation experiments were conducted with spin temperature 19

l.lcltrng tndothcrm 2

É. g 1- é o J e 1- o

Tcmpcrarurc (qc)

Figure 2.6 (a) The peak onset method used for determining melting and crysrallisarion temperatures from DSC thermograms.

-,,

r ? Lcp

T )- r o

Figure 2.6 (b) The midpoint method used for determining glass transitlon temperatures from DSC thermograms. T, is taken as being f f þCp) i.e. the tempemture at which tl¡e heat capacity change is half of that of the total change from the glassy to the rubbery state. 20 alternation and received phase cycling (CYCLOPS ) to remove quad images. The probe temperature was 298+3K. For T5¡ measurements r was in the range 0.1 to 30 ms, while for Tlp(C) measurements delay times ranged from 0.2 to 18 ms. TET urethane acrylate polymers and copolymers were filed to produce a granular powder which was used to pack a boron nitride rotor, with rotor spinning rates being between 2 and 3k}lz.

2.7 Differential Scanning Calorimetry A Perkin-Elmer DSC 7 equipped with a dry box was used for the determination of sample glass transition, crystallisation and melting temperatures. Samples were scanned in the range -100 to +l00oC using liquid nitrogen as coolant and high purity helium as the purge gas. N- octane (m.pt. 216.4 K) and water (m.pt. 273.15 K) were used as temperature calibration standards. N-octane (AHn,:43.59 kcal/g) was used as the enthalpy calibrant for glass transitions, crystallisation and melting peaks. Melting and crystallisation temperatures were measured using the peak onset method (Figure 2.6 a), with glass transition temperatures being measured using the midpoint method (Figure 2.6 b). The peak in the first derivative of the thermogram in the Tg region was used to more accurately determine the glass transition temperature for glass transitions with smáll ACp (Figure 2.6 c).

A DuPont DSC 9900 was used for annealing experiments and high temperature peak determinations up to +200oC. Temperature calibration standards used were identical to those used for calibration of the PE DSC l. The DuPont DSC was not equipped with a dry box. Scanning rates 2T

TI -72_ sJ31 'c f2 -20. o00 'c T9 -52.5?1 'c Do Ito cp 0. 353 J/g-C Oneet -58.8t9 'c

-100- o -75.0 -50. 0 -25- 0 o.0 25. 0 50. 0 75.0 loo- 0 Teoperotrre ('C)

Figure 2.6 (c) An example of the use of the fust derivative of the DSC thermogram to calculate Tg for samples with small ACo. 22 for calibrations and sample runs for both instruments were usually 20Klmin.

2. 8 Water Sorption/Desorption Measurements After casting, samples were alternately placed in deionised water at 25"C for 48 hours and a 50oC vacuum oven for at least 48 hours. This four day cycle was repeated twice in order to remove any unreacted prepolymer from the networks. After drying to constant weight in vacuo samples cut from polymer sheets were either immersed in deionised water or exposed in sealed vessels above saturated electrolyte solutions with different relative humidities at 25+loC. Saturated MgCl, NHaCI and PbN03 solutions gave relative humidities of 33,19 and 98Vo respacfively. Relative humidities achieved by saturated salt solutions have recently

been confirmed 9 as being within ! 3Vo of the values specified.

Desorption runs were made by placing the samples in dessicators of similar size to those used for sorption measurements, containing anhydrous calcium sulphate, after equilibrium water uptake was achieved in the wet environment. Prior to weighing, samples sorbed in water had surface water removed by tissue blotting. Samples were weighed at appropriate intervals on a Mettler 4E166 electronic balance accurate to +0. 1mg. Wur", content for the samples was expressed in terms of weight percent relative to the wet weight (W) l0'

WC = 100 (V/-Wù/W Vo 23

W6 being the dry weight Equilibrium water content (EWC) was the V/C determined when the sample had reached equilibrium in a particular wet environment.

2.9 FTIR/NIR S pectroscopy Atl spectra were obtained on a single-beam Perkin-Elmer 1720 FTIR spectrometer at 2cm-l resolution. 50 scans were signal averaged and stored on magnetic disk. Transmission spectra were obtained from thin films on KBr disks. Attenuated Total Reflectance (ATR) spectra were obtained using a Perkin-Elmer Multiple Internal Reflectance Accessory using a KRS-5 Internal Reflection Element (IRE), at a 45" angle of incidence. Samples for ATR spectra consisted of polymer sheets of approximate dimensions 1 x 4 x 0.2cm placed on one, or where sufficient sample was available, on both sides of the IRE to maximize the signal-to-noise ratio. Where there was insufficient sample to cover the IRE, non-coated aluminium foil was used to provide a backing to the polymer sheet. NIR spectra were obtained by placing the polymer (or in

some cases prepolymer) sample directly in the beam path. In the case of prepolymers samples were enclosed in glass sheets, i.e. prior to casting as described earlier. This configuration did not affect the spectral resolution.

2. l0 Thermogravimetric Analysis Thermal stabilities of urethane acrylate polymers were determined using a Mettler TG50. 5-10mg polymer samples were heated at 20Klmin under nitrogen or oxygen atmospheres (50m1/min). 24 CHAPTER THRBE

DYNAMIC MBCHANICAL MEASUREMBNTS

3.1 Introduction The variation in dynamic mechanical properties of poly(tetramethylene oxide) and poly(e-caprolactone) urethane acrylate networks was investigated as a function of PTMO and PCL molecular weight, using a Dynamic Mechanical Thermal Analyser (DMTA). DMTA scans were also performed on urethane acrylate copolymers, containing up to 50 weight per cent of three acrylate based comonomers : methyl acrylare, hexanediol diacrylate (HDDA) and tetraethylene glycol diacrylate (TEGDA).

The dynamic mechanical properties of PTMO based networks prepared with lower crosslink density (designated TET(PTMO)2 with the appropriate PTMO molecular weight in parentheses) were also examined. Samples of saturated urethane acrylate homopolymers equilibrated in water and hard Segment model polymers, containing no PTMO or PCL soft links, were also scanned.

Variations in the log shear Stolage modulus, log E', log shear loss modulus, log E", and tan ô of a polymer sample with tempefature can provide information on the types and relative intensities of transitions occurring in the polymer as different molecular motions occur. 25

3.2 Results and Discussion 3.2.L Homologous series 3.2. I (a) TET homopolymers After 15 minutes irradiation as described in Chapter Two all samples were readily removed from the glass casting sheets and were stored in a vàcuum dessicator until analysed. The homopolymers produced from the TET650 and the TET1000 prepolymers were leathery, while those cast from the TET2000 and the TET2900 prepolymers were rubbery at room temperature. Attenuated Total Reflectance (ATR) techniques (Chapter Seven) confirmed that the acrylate double bond conversion was uniform from one side of the polymer sheet to the other. The results of the DMTA scans on the TET homologous series are given in Figures 3.la-c. These show the change in tan õ, the log shear storage modulus, log E', and the log shear loss modulus, log E", as a function of temperature. All DMTA scans shown were obtained at an applied frequency of 1 Hz, unless otherwise noted. Three peaks occur in the tan ô - temperature plot of the TET650 : at

+35oC, -75oC and approximately - 1 20oC. Apparent activation energies for the largest tan õ peak were determined from the plot of frequency versus the reciprocal of the absolute temperature at the loss peak maximum (T*u*) using the following equation :

ÂH* = 2.303 * qffm (3.r)

where AH* is the apparent activation enthalpy, fmax the frequency at the loss peak, and R is the gas constant. 26

0.5

650

0.4 o 1000

2000

oo oo tr 2900 cflr o a 0.3 o.'tl ô o !t d o ct EI O sqbo o !0.) o o E I ct ho. % o %b o o o 0.2 o t + o a h a % o a g o q t¡ O a o a E a o ! o a o os o I t a 0.1 E O o f % a a o a ! o ! a o EI + O rh6 .o I t ! a o ¡ I 'ì\*o-

0.0 -r25 -100 -15 -50 -25 0 25 50 15 100 t75

Temperature (oC)

Figure 3.1 a Tan õ - temperature plots for TET urethane acrylates with PTMO soft segments of indicated molecular weighr 27

9.5

s 650

ú o 1000 I o o r O o (, g o r a 2000 8.5 o o I o I a + o 2900 Io fo o O EI +o a + o E rr I rD tr èo I o o e o 7.s e + o g o a a

+ + + \ 6.5 \

\ ++ *1 *t

5-5 -r25 -100 -15 -50 -25 0 25 50 75 100 r25

Temperature (oC)

Figure 3.1^b t og E' - temperature plots for TET u¡ethane acrylates with pTMO soft segments of indicated molecular weighr 28 9

E 650

o 1000 8 4, +E o o ! 2000 o o o o o + (, + 2900 ! o + o F E o + rD E 't + o =t¡l u EI + o èo I o |l 1tr I oo o otr Oq a o 6 ¡E oo otr on a E o #oE a l++ os 5 -q'

4 .r25 -100 -75 -50 -25 0 25 50 15 100 r25

Temperanue (oC)

Figure 3.l^c t og E" - temperature plots for TET urethane acrylates with pTMO soft segments of indicated molecula¡ weighr 29

The activation energies for the largest tan õ peaks for the TET

polymer series were 51, 42,67 and 45 kcal/mole for TET650, 1000, 2000 and 2900 respectively. These activation energies are typical of those found for the glass transition process 1,2. The largest tan ô,peak was

therefore regarded as being due to the amorphous glass transition, and is

designated as cru.

A summary of the dynamic mechanical results for this homopolymer series is presented in Table 3.1.

TABLE 3.1 Dynamic mechanical results for the oligo poly(tetramethylene oxide) based series of urethane acrylates.The Tg (glass transition temperature) has been taken from the tan õ - temperature plot. Tg @") refers to the highest temperature peak in the log E"-temperature plot. Peak height and width at half height flVrd refer to the tan ô glass transition peak. 'W¡ refers to the weight fraction of hard segment (all groups excluding PTMO soft links.)

LogE'(N/m2) PTMO MWt. Tg('c) (25.C) Peak height wp('c) Tg(E") wtr

6s0 +35 7.95 .462 54 -15 0.48 ,] 1000 +13 7.38 329 82 -M 0.37

2000 -50 6.85 302 88* 65 0.23

2900 -50 6.7 | 307 131* -65 0.Ll

* denotes asymmetric peak

The peaks at -15 and -120'C are referred to as p and T respectively as it is usual to designate peaks in order of decreasing temperature. In the TET1000 polymer the glass transition occurred at +13oC, with the p peak occurring as a shoulder to the cu peak at -75"C. The l peak position remained unchanged at -120'C. The tan ô-temperature plots for TET2000 and TET2900 showed

only two peaks - one at -5OoC and the other at -120'C. 30

The aRelaxation The position and intensity of the T peak remained unchanged through the homopolymer series. Kolarik 2 proposed that the y process is a small scale internal motion of the side chains in linear methacrylates - alteration of the length of the side chains in a series of hydrophilic acrylates produced no corresponding change in TT, indicating that the motions are highly localised. In polyurethanes this relaxation has been assigned 3-5 to the local motion of methylene sequences in the soft segment. Kajiyama and MacKnight 3 detected three y relaxation peaks in a series of linear polyurethanes, with the two highest temperature l peaks at -l40oC and -120oC (110H2) assigned to the motions of methylene groups in ether and ester diol sequences respectively.

The I Relaxation The p relaxation was only clearly defined for the TET650 and 1000

polymers, being obscured by the glass transition peak in the TET2000 and - 2900 samples. McCrum et. al.6 attributed the B peak in polymethyl methacrylate (PMMA) to partial rotation of the COOCH3 group about the C-C bond linking it to the main chain. The p peak in PMMA occurs at 280K(lIF.z) 6. The occurrence of P peaks at higher temperatures in polymers with polar side chains has been attributed 2 to polar interactions increasing the activation energy for side chain motion. The extent of the P process in poly HEA has been compared to that of pHEMA 1 ' copolymers of HEMA-HEA showed a reduction in the existing p maxima

and the formation of another peak at 17 8K as the proportion of HEA in the copolymer increased. The position of this peak is some 20K lower than that observed in the TET650 polymer. The side chain in the urethane acrylate polymer series considered here is effectively equivalent in Iength 3l to alternating sequences of hard and soft segments which constitute the main chain in linear polyurethanes. Polymethacrylates with longer side chains than R = C¿H9 may exhibit 2 a T, close to or below that of Tp - the secondary relaxation is overlapped by the glass transition. Even allowing for the fact that p relaxations in polyacrylates occur'at lower temperatures than the corresponding polymethacrylates, rotation of the bulky, polar side chains in these urethane acrylates seems an unlikely mechanism for the observed p peak. The adventitious sorption of minor amounts of water may account for the observed p process - Chien and Rho 8 observed decreases in B peak intensity upon annealing of thermoplastic polyurethane elastomer, while other workers 4'9 attributed the B relaxation peak to water hydrogen bonded to the urethane group.

The g Relaxation

The glass transition temperature decreased as the molecular weight of the PTMO soft segment increased, with the most pronounced drop being from the TETl000 to the TET2000 polymer. As the PTMO chain length increased further in the TET2900 sample the Tg remained unchanged at -5OoC. Koshiba et. al.l0 attributed this large Tg decrease to improved phase Separation between soft and hard segments in the TET2000 polymer. The û, peak for the TET1000 polymer was attributed to combined molecular motions in one homogenous phase, comprising soft segments together with urethane and polyacrylate linkages. The higher temperature shoulder to the tan ô peak at -50oC in the TET 2000 polymer

was taken as additional evidence to indicate that this polymer consisted of two well-defined phases with the shoulder at about +15'C in the TET2000 considered to be due to hard segment relaxations. 32

If the tan õ - temperature plots for the TET2000 and 2900 are compared (Figure 3.1a), it is obvious that the higher temperature shoulder is more pronounced for the higher molecular weight PTMO polymer. Since the TET2900 polymer contains a lower hard segment weight fraction it seems unlikely that this higher temperature shoulder originates from relaxations in the urethane acrylate part of the molecule. Tan ô(Tg) decreased as the PTMO soft link molecular weight increased from 650 to 2000. Increases in tan ô(Tg) have been correlated with decreasing crosslink density in amorphous networks 11,12, however, the development of crystallinity in polymer chains with increasing molecular weight complicates this interpretation. Andrady and Sefcik 12 attributed the decrease in tan ô(Tg) with increasing molecular weight between crosslinks in a poly e-caprolactone system to crystallisation.

Allen et. al. l3 observed both increases and decreases in tan ô(Tg) in a series of oligomeric ethylene glycol dimethacrylates. As the molecular weight of the ethylene glycol chain was increased from 130 to 400 ( three and nine ethylene glycol units respectively) tan ô(Tg) increased from 0.095 to 0.51. The peak height changed little from nine to thirteen repeat units (tan õ (Tg) was 0.55 for the higher molecular weight oligomer) but for the next sample studied (containing 22 oxyethylene repeat units) the tan ô(Tg) was found to be 0.27. The high temperature shoulder which is more prominent in the TET2900 sample may be due to crystallite melting in the polymer. It is believed that, even though the amorphous material in crystalline polymers undergoes its own set of motional transitions, much the same as in the completely amorphous polymer 15, some perturbation of the amorphous transition is inevitable, with commonly observed effects including

shifting of the cr,a process to higher temperatures and diminution of tan õ(Tg). Wadhwa and V/alsh l4 found that increasing the molecular weight 33 of a poly(ethylene adipate) soft segment in a urethane acrylate from 4600

to 6000 increased the Tg from -18 to -6oC, and attributed this increase ro the development of crystallinity in the network with the higher molecular weight soft segment.

The constancy of the peak position at -50"C for both the. TET2000 and 2900 polymers suggests that PTMO crystallinity develops after thermal scanning through the TET2900 ou peak. The dimunition in tan ô(Tg) from TET650 to TET2000 may reflect restrictions on chain motion which occur as a prelude to the development of more extensive

crystallinity in the polymer, which is manifested by both a decrease in tan ô(TS) and an increase in Tg.

The data in Table 3.1 also show that there is a considerable temperature difference between the Tg peak in the tan õ and the log E" - temperature plots, with the maximum in the log E" temperature plots being found at a lower temperature in all cases. A slight increase in this difference occurs from the TET650 to the TET1000 polymer with the difference decreasing sharply for the two polymers of higher molecular

weight. Felisberti et. al. l6 observed a similar difference between loss modulus and tan ô peak positions in a series of copolymers of maleic anhydride crosslinked polystyrene (p(ScoMA)) with linear polyvinylmethylether (PVME). Samples containing small amounts of either component were found to have log E" and tan ô maxima separated by less than 1OoC while in intermediate composition samples (40-10 wtTopScoMA) the maxima were separated by up to 35"C. Measurements made at frequencies from 0. I to 100 Hz showed that the log loss maximum shifted about 10oC with the tan ô maximum being shifted between 25 and 30"C. This was taken as evidence for the occurrence of molecular relaxations with different rate constants. This was not observed in scans of the TET homopolymers at 10, 1 and 0. I Hz where 34 log E" and tan õ peak temperature shifts with frequency wefe found to be roughly similar. Dynamic mechanical data obtained by Miller et.al. l7 revealed that, in a series of linear polyurethanes based on a PTMO of molecular weight 990 , and MDI/BD hard segments the log E" maximum occurred at a lower temperature than the tan õ maximum. As the MDI content increased from

20 to 48 weight Vo the difference increased from about l4 to 40oC. Since these polymers were linear, crosslinking effects cannot be invoked to account for this increased separation of the two maxima, but an increasing interaction of the hard segment with the soft segment may lead to a decrease in soft segment mobility, with an attendant increase in the Tg of the polymer. It is possible that the tan ô maximum is a more accurate reflection of the actual state of the polymer than the maximum in log E" : the TET 650 sample is leathery at room temperature (Tan ô max=+35oC) but the log E" maximum is at -15oC. The log E" maximum may be an indicator of molecular mobility in the PTMO soft segment only, rather than that of the entire polymer.

2 1 lh'ì TPCL homonolvmers The results of the DMTA scans for the PCL soft segment homologous series are given in Figures 3.2a-c- These show the change in tan ô, the loss modulus log E" and the storage modulus log E'as a function of temperature. A summary of the dynamic mechanical properties for this homologous series is given in Table 3.2. 35

0.5

tr 1250

ao 0.4 o 2000 ooo o o oæ o otr"4 o O o 3000 o tr o tr E¡ o o oo E 0.3 o a tr o o GI d q) o -é rt E¡a o ñ (t o Elo o (t 0.2 o o o o a o o o o o O Oa o atr &' O¡ tr o 0.1 o E oo oE- o E ooo o oog o tr o E E ta t" oa o o o aoo 0.0 -100 -75 -50 -25 0 25 50 15 100

Temperature ('C)

Figure 3.2 a Tan ô - temperature plots for TPCL urethane acrylates with PCL soft segments of indicaæd molecula¡ weighr 36

o 1250

9 . 2000 o O¡ 0o oo o o 3000 o oo s o o o tr ti o 8 oo o oèo oo EI oo Oi,

O O a o a

ael o oo¡-ereo oo ot

6 -100 -15 -50 -25 0 25 50 15 100

Temperature ("C)

Figure 3.2 b l,og E' - temperature plos for TPCL urethane acrylates with PCL soft segments of indicated molecular weight. 37

o 1250

8 o o o o o 2000 oa tr o a E o E a E o O tr o 7 o E 3000 E o o o o E o so tr a o tr o g t o E oo Qc tr f¡ì 6 oo E E õo o o E o o o (lo %

ooo EI EI 5 E o tEE¡ .. tr o û

3 100 -'t5 -50 -25 0 25 50 15 100

Temperature(oC)

Figure 3.2 cl-og E" - temperature plots for TPCL urethane acrylates with PCL soft segments of indicatirC molecular weight. 38

TABLE 3.2 Dynamic mechanical results for the TPCL homologous series. The Tg was determined from the tan ô- temperature plot. Tg(E") refers to the largest peak in the log E"- remperaure plot. Peak height and width at hatf height refer to the tan õ glass transition 'l'g("c) PCL LogE'(N/mz¡çZS"C) Peak height 'Wp("C) Tg(E") M.V/t rz50 +8 1.2027 0.3711 55 -27 2000 -26 ó.EE2I o.4045 54 -40 3000 -37 7.6300 0.3870 3E -45

A large decrease in the homopolymer glass transition temperature was observed with an increase in PCL molecular weight from 1250 to 2000 with a smaller decrease as the soft segment MW increased to 3000. The log storage modulii at room temperature did not follow a predictable trend with log E' for TPCL 3000 at 25oC being considerably greater than that for the TPCL 2000 polymer. Examination of the log storage modulus versus temperature plot shows a small plateau between -15 and +10"C, presumably due to the presence of a crystalline phase. Some TET2900 copolymer samples showed an increase in log Storage modulus aS the temperature increased - this was attributed to PTMO soft link crystallisation during the thermal scan. It therefore seems likely that crystallites were present in the TPCL3000 polymer prior to the thermal scan. While there appeared to be a "levelling off" in lower limit to the Tg for the PTMO series ( i.e. Tg's for the TET2000 and 2900 were identical) the Tg decreased further from the TPCL2000 to the TPCL3000. The temperatures of tan ô maxima and log loss modulus maxima again differed, as observed previously for the PTMO based samples, with the magnitude of this difference decreasing with increasing soft segment length. 39

3.2. I (c) TET(PTMO)2 homopolymers The dynamic mechanical results for TET(PTMO)2 polymers are given in Figures 3.3a-c. These show the change in tan ô, the loss modulus, log E" and the storage modulus log E' aS a function of temperature. A summary of the dynamic mechanical properties for this polymer series is given in Table 3.3.

In common with the PTMO series with only one PTMO unit between crosslinks there was a considerable decrease in glass transition temperature as the PTMO molecular weight increased. From the DMTA scan it appears that the TDI group interconnecting the two PTMO 2000 chains in the TET(2000)2 polymer also does not restrict soft segment crystallisation, with this crystallisation being manifested in an increase in storage modulus as the sample was scanned.

TABLE 3.3 Dynamic mechanical results for the TET(PTMO)2 homologous series. The Tg was determined from the tan õ - temperature plot. Tg@") refers to the largest peak in the log E"-remperature plot. Peak height and width at half height flVrfz) refer to the tan ô glass transition peak.

FTTMO lsCC) LogE'(N/m\25"C Peak wtp("c) Tg(8") lv1Wt height -6 6.ó68 0.4987 63 -32

I 1000 t2 -'t2 6.831 0.4215 48 -46

I ¿t )v 2 -54 6.446 0.5343 32 -63 ¿9VU t2 -41 6.3s9 0.1 l3 -57

The dynamic mechanical properties of the TET(2900)2 sample contrasted markedly with those of the TET(2000)2 polymer - with the magnitude of the tan ô peak in TET(2900)2 near -50'C decreasing to about

20Vo of that seen for TET(2000)2. The soft segment Tg peak was also shifted upwards to -41oC, and a new peak developed at +1OoC- 40

0.6 o o o a o o EI (6s0)2 0.5 oPFt a Ettr o o s E + (1000)2 s **+ s o + tr 0.4 + a (2000)2 o tr a + + E a Eg o (2e00)2 E+ E + tr cl a + o q, 0.3 tat + EI

cÉ + a E o.2 EI + E tr + Et a EI + t? E a o + .ofr. 0.1 ao ++ç + EI o + *+ Eb tr ¡ EtsEI GI **++ + + +++ **** 0.0 -100 -15 -50 -25 0 25 50 15 100

Temperanue ("C)

Figure 3.3 a Tan ô - temperature plots for TET(PTMO)2 urethane acrylates with FrIMO soft segmens of indicated molecular weight. 4r

9.0 E a E (6s0)2 tr o a o (1000)2 a o o o (2000)2 o a a a 8.0 o Â (29O0)2 a o aa o o aa o O a a E l¡l o 'â O Eto O Þo o EI o ooo EI O 7.0 aoeoooræo' .U' l-lt o o D-_. .- Â O¡

6.0

5.0 -100 -15 -50 -25 0 25 50 15 100

Temperanne ("C)

[Cule 3.3 b Log E' - temperature plots for TET(PTM )2 urethane acrylates with FTMO soft segmens of indicated molecular weight. 42

I (6s0)2 o GI 8 + E + (1000)2 + E o a (2000)2 a GI a a + Â (2e00)2 oa a + 7 + + c + Ð : + H + r è0 ** 8E o +tr

t+k

5 + +

4 -100 -15 -50 -25 0 25 50 15 100

Temperanne ('C)

3.3 c Log E" plots IEgt --.temperature for TET(PTMO)2 urethane acrylares with PTMO soft segmens of indicãted mole^cular weight. 43

The peak at +10"C is at the same temperature found for the cr" peak in linear polyurethane block copolymers 4 where the peak was attributed to either an interaction between the amorphous and crystalline portion of the soft phase or melting of small imperfect crystallites. The melting endotherm at 17oC noted for TET(2900)2 in DSC experiments (Chapter

Four) tends to favour the latter interpretation. The higher temperature at which this peak was found in the TET(2900)2 sample contrasts with the crystalline phase being evident as a shoulder to the cru peak in urethane acrylates with higher crosslink density. The higher temperature of the cr" peak in the TET(2900)2 may result from the formation of more ordered crystallites, despite careful pretreatment of the TET(2900)2 sample prior to the DMTA scan which involved annealing in a 50oC oven for four hours, followed by rapid quenching. The rate of crystallisation in this polymer may be such that a fully amorphous PTMO phase cannot be readily isolated. 3.2. I (d) Hard segment model polymers Dynamic mechanical plots fo¡ hard segment model polymers cast from TDI/HEA prepolymers with TDI/HEA in mole ratios of 1:1 and l:2 are shown in Figure 3.4. The main tan õ peak in the linear 1:1 TDI/HEA network was near

80oC, with a high temperature shoulder at 110'C. A third peak was observed at about 160'C. Scanning the linear TDI/HEA polymer at a slower rate, (lK/min), resulted in an improvement in peak definition, with three well defined peaks occurring at 600, 108o and 135'C (Figure 3.5). The peak at 60o is undoubtedly that due to the glass transition process, as it is accompanied

by a decrease in the storage modulus of about two orders of magnitude. The tan ô-temperature plot of the crosslinked I:2 mole ratio

TDI/FIEA polymer showed a main peak at 84o with a pronounced high 44

0.5 o 1:1

0.4 o l:2

tro E 0.3 cÉ o tr o q) E O õ a a o cb. GI 0:2 tr o ¡o\' EO "sdÐ 0.1 ho*l@ Oa 0.0 -50 -25 0 25 50 75 100 r25 150 r75 2N Temperature ('C)

Figure 3.4 Tan ô - temperature plots for linear (1:l mole ratio) and crosslinked (1:2 mole ratio) TDI/FIEA hard segment model polymers.

0.6

0.5

0.4

cË tr tr q) E¡ ! 0.3 E E cl tr E o 0.2 E

0.1

0.0 -30 -5 20 45 10 95 t20 r45 170 195 220 Temperature (oC)

Figure 3.5 Tan ô - temperature plot for linear (1:1 mole ratio) TDVHEA hard segment model polymer scanned at a slower rate(1K/min) than in Figure 3.4.(3lVmin). 45

9.5

l¡l 8.5 oèo

7.5 .50 -25 0 25 50 75 100 r25 150 r75 200 Temperature (oC) Figure 3.6 Log E'- temperature plot for crosslinked (l:2 mole ratio) TDVHEA hard segment model polymer.

320 E 300

280 v (u 260 E¡ Tgexp

GI l- o (u E O À 240 Tgcalc E E¡ o q) tr 220 o o oo 200 o

180 0.1 0.2 0.3 0.4 0.5 wh Figure 3.7 Variation in experimental and calculated Tgs with hard segment weight frattion for TET and TET(PTMO)2 polymers. Tgs were calculated using the Fox equation. 46 temperature shoulder between 130 and 150'C. Assignment of the glass transition to one or other of these peaks was not possible, since the decrease in log E' with temperature occurs in two-steps (Figure 3.6), suggesting that both peaks may be due to amorphous motions in the polymer. The apparent activation energy for the higher temperature peak obtained from multifrequency scanning was 63kcal mol-1, typical of that for major glass transitions 1. Tgs for TDI based hard segments have been found to vary widely depending on the other hard segment component (termed the "chain extender") from 107o for TDI-BD hard segment l8 to between 180 and

190o for TDl/ethylene diamine hard segments 19. The temperature chosen for the crosslinked l:2 mole ratio TDI/f{F'A polymer Tg was that of the midpoint of the plateau between the two step decreases in storage modulus:110oC, with a derived Tg value for the TDIÆIEA hard segment in the urethane acrylate networks being 85oC, the average of the linear (60"C) and crosslinked Tg values. Calculated values for Tgs for the TET, TET(PTMO)2 and TPCL polymers ìwere obtained using the copolymer equation :

I wl w2 + (3.2) To-õ Tgl Te2

where w1 and wz are the weight fractions of soft and hard segment, and Tgl = 189K 20 and Tg2= 346K are the glass transition temperatures for PTMO soft segment and the TDI/HEA hard segment respectively. For the PCL soft segment Tgt = 208K 21. Experimental Tgs, together with those calculated from the copolymer equation for TET and TET(PTMO)2 polymers, are ptotted against weight fraction hard segment in Figure 3.7. 47

It is apparent that the copolymer equation underestimates the polymer glass transition temperature for all the urethane acrylates studied. This underestimation is most severe at high hard segment weight fractions with the experimental Tg of TET 650 polymer being some 64K higher than that calculated from the copolymer equation. In general the calculated values diverged more from experimental values in polymers with lower soft segment molecular weights. It appears that Tg for these polymers is dependent to some extent on the network crosslink density. Shefer and Gottlieb 22 in comparing theoretical methods for calculating polymer Tg, identified four main types of crosslinking reactions : (I) Reactions between two or more functional groups attached to the ends of low molecular weight species, resulting in the formation of highly cross-linked thermoset systems. Epoxy resin formation is a typical reaction in this category. (II) End-linking of high molecular weight prepolymers by means of low molecular weight crosslinkers leading to the formation of lightly crosslinked elastomers. Crosslink density is determined by the molecular weight of the prepolymer (Mg). Shefer and Gottlieb arbitrarily defined the transition from type I to type II reactions as occurring after the prepolymer molecular weight exceeded 1000. (III) Networks formed by the copolymerisation of di- and multifunctional monomers (e.g. styrene/divinyl benzene) (IV) Vulcanisation of polymer chains by crosslinking of functional groups on the polymer backbone (e.g. sulfur crosslinking of natural

rubber. ) The crosslinking of the urethane acrylate prepolymers resembles the type II crosslinking reaction, in which the main contributions to increasing Tg were considered to be the formation of branch points and crosslinks. Monomer depletion, reduction in the concentration of chain 48 ends and non-Gaussian chain statistics were considered 22 to be relatively unimportant in determining the polymer Tg for end linking of high molecular weight prepolymers. There are some important differences, however, between type II cros slinking reactions and the photopolymerisation of urethane acrylate prepolymers. The- urethane acrylates contain two double bonds and can be regarded as being "self- crosslinking", forming insoluble networks without additional low molecular weight crosslinker. Equating the crosslink density with the molecular weight of the prepolymer also presents some difficulties. The urethane acrylate polymers considered here are not simply end-linked PTMO chains - the bulky TDI and HEA end groups contribute significantly to prepolymer molecular weight. Despite these differences the crosslinking model for type II reactions offers a useful starting point for theoretical calculations of Tg fo¡ urethane acrylate polymers. A model for Tg behaviour of the urethane acrylate polymers will be developed with a simplifying assumption : that the observed glass transition in the polymers is solely that due to the PTMO units, with the remainder of the network being regarded as relatively immobile tie down points for the PTMO chains.

Stutz, Illers and Mertz 23 modified De Benedetto's equation 24

Ts = rr,uIt + KzX"/(1-x") ] r¡¡l where X, is the crosslink density expressed as the mole fraction of all crosslinks present in the system weighted by functionality. The parameter K2 is related to DeBenedetto's lattice energy ratio and characterises the influence of the crosslinks on a particular system. A Kz value of 0.73, obtained by Stutz et. aI.23 for a crosslinked polyurethane 49

containing poly(propylene oxide) soft segment, was used in the present work. The term Tg,u, referring to the Tg of an un-crosslinked system identical to the crosslinked one in every respect except that the crosslinks are missing, is used in preference to Tr,- the Tg of a linear poly.mer chain of infinite molecular weight, in order to account for restrictions on PMTO chain end motion in the prepolymer. Tg,u for the PTMO based networks was taken as 200K, some 13K higher than that for high molecular weight PTMO in order to account for the decrease in chain mobility due to the TDIÆIEA end groups. Tg,u for the PCL networks was slightly higher, at 208K. The weighted total

number of crosslinks and junctions, X., is given by :

f Xc =u, I i=z (i-2)/ 2 Xi (3.4)

where Xi is the fraction of i-functional a¡ molecules reacted. This terminology derives from determinations of Tg for linear polymers t condensation crosslinked by low molecular weight crosslinkers (e.g. vinyl terminated linear poly(dimethyl siloxane) (PDMS) crosslinked with

tetrakis (dimethylsiloxyl) silane(HS iMezO)¿S t) 22, where for incompletely

reacted systems X¡ is less than 1. In the urethane acrylates considered here X¡ may be taken as being 1, since double bond conversion is essentially complete. A¡ is the mole

fraction of crosslinker in the system.

If the urethane acrylate prepolymer molecule is regarded as a tetrafunctional crosslinker, then a¡ can be taken as the mole fraction of

groups excluding the soft segment in the TET and TPCL prepolymers. It is more difficult to evaluate the contribution of the central TDI group in 50

TET(PTMO)2 polymers - DSC results (Chapter Four) suggest, however, that this group has much less influence on PTMO crystallisation behaviour than the crosslinks at either end of the polymer molecule. In determining Xc for the TET(PTMO)2 polymers, this central TDI unit was omitted from molecular weight calculations.

An example of the calculation of Tg for the TET650 homopolymer

using the DeBenedetto crosslink model is shown : For the TET650 PolYmer: af = uo = ffi 600.5 1 ã4= t254.36 = 0.479 f Xc = 0'a79 Z i=2 (i-2)/ 2

for i=4 (tetrafunctional crosslinker) Xc = O.479

Ts 2oo[1 + 0.73(0. 41s I (1 - o. 47s)) ]

t Ts 334K for TET650 polymer (Tg"*o = 308K)

Experimental and calculated Tgs for the TET, and TET(PTMO)2 polymers are plotted in Figure 3.8 as a function of crosslink density. Reasonable agreement between experimental and calculated glass transition temperatures was found, with some calculated values coinciding with the actual polymer Tg. Equation 3.3 also gave reasonable agreement with the experimental Tgs for the TPCL polymers. 5l

350 +

v o (¡) 300 ¡i a o Tgexp cÉ ¡r (u + Tgcalc q I q) 250 t o + o o + c

200 0.0 0.1 0.2 0.3 0.4 0.5 Crosslink density (Xc) Figure 3.8 Variation in experimental and calculated Tgs versus network crosslink deñsity for TET and TET(PTMO)2 polymers. Tgs were calculated using DeBenedetto's equation. 52

The higher experimental Tg for the TET(290O)2 polymer may be due to the presence of PTMO crystallinity, which was not accounted for in the model. The calculated Tg for the TET650 polymer overestimated the actual value by 26K, which may reflect the transition from type II to a type I crosslinking reaction discussed earlier. For type I crosslinked systems other factors, e.g. chain end disappearance, plasticisation by unreacted molecules and cyclisation may influence the observed Tg. As the molecular weight of the PTMO link decreases other factors apart from crosslink density come into play.

The much improved prediction of Tg by the DeBenedetto model compared to the copolymer equation suggest that the underlying assumption made initially is reasonable : Tgs for the urethane acrylates studied reflect restrictions on soft segment motion imposed by chemical crosslinks. Physical crosslinks formed by hard domains appear to play no part in the dynamic mechanical behaviour of these networks. 53

3.2.2 Co po ly me rs 3.2.2 (a) TET copolymers (i) TET650 copolymers Tan ô -temperature plots for TET65O/methyl acrylate copolymers are shown in Figure 3.9, with DMTA results for all TET650 copolymers summarised in Table 3.4. Tan ô and log storage modulus - temperature plots for poly(TEGDA) and poly(HDDA) homopolymers are shown in

Figures 3.10a and b.

TABLE 3.4 Dynamic mechanical results for TET650 ursthane acrylate copolymers. The Tg (glass transition temperature) was determined from the tan &temperature plot. Tg@") refers to the highest temperature peak in the log E"-temperature plot. Peak height refers to the tan õ glass transition peak. 'lg(8") WtTodiluent I|gdoIVo refc) Peak diluent height methyl acrylate zt) 18.4 +29 7.801 o.661 +I3 50 93.5 +27 1.195 t.234 +7 TEGDA 20 50.8 +43 8.353 0.356 0 5U ðu.5 +44 8.0ó5 rJ.324 +9 HDDA 20 58.0 +50 8.512 0.307 -z 50 84.ó +57 8.770 0.225 I

Only one glass transition peak was observed for both TET650 methyl acrylate copolymers - this is not surprising since there is only a small difference in Tg between the two polymers : poly(methyl acrylate) has a Tg of 25"C 15. The large increase in tan ô (Tg) for the 50 wt%o methyl acrylate copolymer reflects an increase in segmental motion resulting from a large decrease in crosslink density. Copolymers of TET650 with 20 and 50 wtVo TFGDA show little difference in dynamic mechanical properties (Figure 3.11). A slight

narrowing of the glass transition peak width was apparent for the 50 wt%o 54

1.5

o E 20 o o s0 1.0 cq a {) õ o

ct Ett E¡ g tr tr 0.5 o E E O q oE o rD 0.0 -100 -75 -50 -25 0 25 50 75 100 Temperature (oC)

Figure 3.9 Tan ô - temperature plots for TET 650 copolymers containing the indicated wt%o methyl acrylate.

0.4 O TEGDA . HDDA 0.3 E

c€ q) õ 0.2 GI C! oooo to o o O o 0.1 o o o'r o o o o t E o s o s tr ".".r.råTf.t l".".tåt E O o6 oEEo 0.0 100 -15 -50 -25 0 25 50 75 100 Temperature (oC)

Figure 3.10 a Tan ô - temperature plots for TEGDA and HDDA homopolymers. 55

l0 S TEGDA O HDDA EEtr ry cooote

9 oooooooaoo tr frl ¡t oè0

8 E Egg EggE

7 lm -15 -50 -25 0 25 50 75 100 Temperature (oC)

Figure 3.10 b Log E' - temperature plots for TEGDA and HDDA homopolymers

0.4 E 20

O 50 0.3

ctl a o) o c 'tt O 0.2 O f GI tr o o E o 0.1 oEI o Og a o

0.0 -100 -15 -50 -25 0 2s 50 75 r00 Temperanre (oC)

Figure 3.11Tan õ - temperature plots for TET 650 copolymers containing the indicated wtTo TEGDA. 56

TEGDA copolymer, pres umably due to decreased network heterogeneity t:. The effects of increasing crosslink density and an upward shift in glass transition temperature are apparent for the TET650 co HDDA nerworks (Figure 3.12). Increased crosslink density has previously been correlated l2 with decreases in tan ô peak height. (ii) TET1000 copolymers

TABLE 3.5 Dynamic mechanical results for TET1000 urethane acrylate copolymers. The Tg (glass transition temperature) was determined from the tan &temperature plot. Tg@") refers to the highest temperature peak in the log E"-temperature plot. Peak height refers to the ta¡r ô glass transition peak.

WtTodiluent lv1ol7o 'IsCC) l-ogE'(N/m¿)l¿5"c) Peak Tg(8") diluent height methyl acrylate 10 68 +l'l 7.278 .3922 -17 ztJ E2 +21 7.214 .4835 -2 30 89 +17 7.160 .6274 -2 40 93 +IÓ 6.933 .7942 +l 50 95 +ZtJ 6.878 1.0088 +4 TEGDA 10 37 +25 7.602 .2958 -33 20 57 +28 7.859 2662 -21 30 70 +36 8.027 .258t -17 40 18 +39 7.80ó 2t59 -24 50 84 +4I 8.295 .2638 +9 HDDA l0 44 +23 1.664 .2708 -zr 20 64 +35 7.938 .230t -18 30 75 +43 8.218 .1984 -6 40 83 +53 8.102 1788 -I 50 88 +52 8.376 l6E I +8

The tan ô-temperature plots for copolymers of TET1000 containing l0 to 50 wfVo methyl acrylate are shown in Figure 3.13. With the addition of methyl acrylate there is an immediate narrowing of the tan ô peak with the glass transition temperature given by the tan ô maximum 57

0.4 o20 o50 0.3 g E' GI tr õq) 0.2 cl Eo trO Eo 0.1 E o Eoa EO tr % 0.0 -100 -75 -50 -25 0 25 50 75 100 Temperanue (oC)

Figure 3.12 Tan ô - temperature plots for TET 650 copolymers containing the indicated Wt% HDDA. t.2 .10 a 1.0 a .20 o30 0.8 a ra .40 CI .50 q) ro o õ 0.6 a t o

cÉ a .o a o þta. 0.4 . It a. a .lo rf o(¡ .O a+ 0.2 a a a Q+ aa ;.{

0.0 100 15 50 -25 0 25 50 75 100 Temperature ('C)

Figure 3.13 Tan ô - temperature plots for TET1000 copolymers containing the indicated wt% methyl acrylate. 58 remaining approximately constant once the methyl acrylate content exceeds 2O wtVo. The maximum in the log E"-temperature plot reflects a more gradual change, however, with the log E" maximum at SïVomethyl acrylate still being some 16"C below that of the tan ô peak. Table 3.5 summarises the dynamic mechanical results for these three copolymer series, with the diluent content also being expressed as a mole percentage.

Dynamic mechanical results for copolymers of TETl000 with TEGDA are presented in Figure 3.I4. Only a single o peak was found for this copolymer series with the tan ô (Tg) decreasing as the temperature of the glass transition increased with increasing TEGDA content. The differences between the loss maxima for tan õ and log E" were generally greater than that observed for the TETl000/methyl acrylate copolymers. Tan ô -temperature plots for copolymers of TET1000 with HDDA are shown in Figure 3.15. The changes in mechanical properties showed similar trends to those for the TEGDA copolymers i.e. a decrease in tan õ (Tg) and a broadening of the main transition peak with an increase in HDDA content. (iii)TET2000 copolymers Tan õ -temperature plots for copolymers of TET2000 containing 10 to 50 wt7¿ methyl acrylate are shown in Figure 3.16. V/hile the narrowing of the largest tan ô peak on copolymerisation was similar to that seen in TETl000 copolymers, distinct low temperature shoulders could be discerned in the copolymers containing 20 and 30 weight Vo methyl acrylate. 59

0.4

E¡ 10 +20 0.3 stros o30 *Jq3 EI + a40 tr + aO cll + o50 a A Ff õq,) 0.2 +e El,¡ cl 'o

0.1

o o o

0.0 lm -75 -50 -25 0 25 50 75 100 Temperature ("C)

Figure 3.14 Tan ô - temperature plots for TET 1000 copolymers containing the indicated wtTo TEGDA.

0.3 tr10 at"" .20 t tf,. .t ct !30 o 0.2 q) a ¡EE Ð õ o40 E o o a O ! ao ;% CE o å. .50 a o E E o.oto E O+ a o E Eo E o e o o. o a ! 0.1 EO õ O I' õ a o' a o a E ro. O Ia+ E oio' Eo ! o "tt 0.0 -100 -75 -50 -25 0 25 50 75 100 Temperaure (oC)

Figure 3.15 Tan ô - temperature plots for TET 1000 copolymers containing the indicated wtTo HDDA. 60

0.8 ol0 .20 0.6 o C! o E30 q) õ to+ o o40

cg 0.4 ¡!¡ o .50 n a !9 o O! 0.2 o o' "".,.l".:.TfË, Es o E E a " E 0.0 "*iËii'.ãTi:'o' 100 -15 -50 -25 0 25 50 75 100 Temperatue (oC)

Figure 3.16 Tan õ - temperature plots for TET2000 copolymers containing the indicated wt%o methyl acrylate.

0.3

Er 10 .20

GI r30 0.2 q) ! "EtroEEooErE o40 tr aaaaoooooo 6 o .50 !. o 0.1 OE a a E Þ ¡iïiiii:::: E} "iþ

0.0 -100 -75 -50 -25 0 25 50 15 100 Temperature CC)

Figure 3.17 Tan ô - temperature plots for TET 2000 copolymers containing the indicated wtTo TEGDA. 61

TABLE 3.6 Dynamic mechanical results for TET2000 urethane acrylate copolymers. The Tg (glass transition temperature) was determined f¡om the tan &temperature plot. Tg@") refers to the highest temperature peak in the log E"-temperature plot- Peak height refers to the tan ô glass transition peak.

WtTodiluent MoLTo 'lg("c) LngE'(N/m¿)l¿J"v) Peak herght :l'g(ts") diluent

10 77 -38 6.6t:2 .25'37 -58 ZtJ 88 I 6.723 .3083 -59 30 93 +7 6.74tJ .4016 -54 40 95 +14 6.688 .5409 -58 50 97 +1ó 6.135 .7338 +l TEGDA IO 49 -48,-lU 7.157 .L937,.182 -58 zrJ 68 -52,+ZI 7.428 .15'39,.t72 -ól 30 79 -57,+30 1.715 .rt70,.t625 -63 40 85 -51,+42 7.925 0948,.1977 -62 50 90 -56,+45 8.169 .066I,.2019 -61 HDDA l0 56 -50,-I2 7.189 1911 -59 20 74 -50,+10 7.596 .1603 -59 30 83 -50,+24 7.90r r37 5 -59 40 89 -50,+53 8.018 .1283 -59 50 9Z -50,+59 8. IÓÓ r304 -59

Interesting trends were also observed in the log E" maxima for this copolymer series. For those copolymers up to and including 40 wt Vo methyl acrylate the log E" maxima is at about -58oC. While this is some 7'C higher than the maximum for the TET 2000 homopolymer, the fact that it remains unchanged over this copolymer range would seem to indicate that this peak is still attributable to motion in the PTMO 2000 soft segment. It is not until the methyl acrylate content reaches 5lwtVo (97 moleVo) that the log E" maximum shifts to near that observed for the 50wtVo TET100O/methyl acrylate copolymer, hence presumably arising from motions in the poly(methyl acrylate ). Table 3.6 summa¡ises the dynamic mechanical results for these samples, together with those for copolymers of TET2000 with TEGDA and HDDA. 62

Dynamic mechanical results for copolymers of TET2000 with from

10 to 50 wt%o TEGDA are shown in Figure 3.L7. The tan ô- temperature plot clearly shows the presence of two peaks for all TEGDA weight percentages. The higher temperature tan õ peak for this copolymer series shifts from -l0oC to +45oC as the TEGDA content increases from 10 to 50 wt%o. The higher temperature peak found for the 5Ùwt%o TEGDA sample is very close to that obtained for photopolymerised TEGDA homopolymer (+50'C). For tan ô -temperature plots with two tan õ maxima the low temperature peak will be designated as Tg1, with the higher temperature glass transition process referred to as Tg2. The value for tan õ (Tg) for this peak decreased initially from 10 to 20 wt%o TEGDA, but subsequently increased for the 40 and 50 wt%o TEGDA samples. The corresponding half height peak widths for Tg2 increased for intermediate values of TEGDA incorporation but decreased as the 50 wtVo level was approached. Similar trends were noted by Bennett 26 in a study of hydrated p(HEMA)/oligo(ethylene glycol) dimethacrylate copolymers. Intermediate water contents yielded tan ô peaks which were broader than those of either dry or fully hydrated samples. The explanation offered was that at intermediate water contents the mobile kinetic units which contribute to the glass transition existed in a greater range of environments than in either the dry or the hydrated samples. If this argument is extended to cover the TET 2000/TEGDA copolymer series, then it appears that the 30 wt Vo TEGDA sample is the most heterogenous. Examination of the DMTA scans for copolymers of TET2000 with HDDA (Figure 3.18) generally yielded only one clearly defined peak in the tan õ plots with a lower temperature shoulder to this peak in the region of that previously observed for the PTMO glass transition. 63

Tg2 increased from -12"C to +59"C as the HDDA content was increased to 50 wt%o, with the peak in the 50wt7o HDDA copolymer being some 17oC less than the Tg of HDDA homopolymer. This may account for the fact that the tan ô (Tg) did not increase after an initial decrease as was the case for the TET2000/TEGDA copolymer series. The tan ô (Tg) value for the 5Ùwt%o sample was also less than that for the HDDA homopolymer, leading to the conclusion that increases in tan ô (Tg2¡ may occur at higher HDDA contents.

(iv )TET29 00copo lymers The dynamic mechanical plots for copolymers of TET2900 with from 10 to 50weight 7o mefhyl acrylate are presented in Figure 3.19a. It is immediately apparent that for several of these copolymers there are two separate glass transition peaks. The Tg1 peak near -50oC is rather poorly defined after the methyl acrylate content is greater than about 2O wt 7o. Generally, even at low levels of methyl acrylate incorporation,the Tg2 peak is narrower than that of the PTMO glass transition. The position of the Tg1 also appears to be relatively constant while there is a progressive temperature increase for the Tg2 peak. The data from this copolymer series in addition to that for copolymers of TET2900 with TEGDA and HDDA is summarised in Table

3.1 . 64

0.2 o10 otrtrotretrtr"t"" tr .20 ooot_I{.^"t E30 _...E c! ._:.,.',,:.^...tr q.) o40 E ir.ri, e .50 t o c! 0.1 _ooo ooo "":::id,..*t o ¡ C Eô "::;..::1""' '" t¡ê o OI. .+ Et E ls oE " .I I o tút s

0.0 -100 -75 -50 -25 0 25 50 75 100 Temperarure (oC)

Figure 3.18 Tan ô - temperature plots for TET 2000 copolymers containing the indicated wtTo HDDA.

0.8 a o10 a .20 t 0.6 GI o30 q) ! o e40 ñ 0.4 oo o '50 o o """" o + oo o tr .ttt9€ d Etr o 0.2 o o a o o o o tr O o + tr t + +e t +t + a aa 0.0 -100 -15 -50 -25 o 25 50 75 100 Temperature (oC)

Figure 3.19 a Tan ô - remperature plots for TET2900 copolymers containing the indicated wtTo methyl acrylate. 65

TABLE 3.7 Dynamic mechanical results for TET2900 u¡ethane acrylate copolymers. The Tg (glass transition temperature) was determined from the tan &tempe¡ature plot. Tg@") refers to the highest temperature peak in the log E"-temperature plot. Peak height refers to the tan ô glass transition peak.

MoIVo Tg("c) tngE'(N/m'¿)('¿5"c) Peak height 1g(tj') diluent methyl acrylate l0 82 -50,-7 6.5944 .3t84,.2r72 -6r,-26 20 9I -41,+2 6.s990 .2529,.2439 -6r,-23 30 95 -40,+8 6.561 .1438,.35I9 -63,-18 40 9ó -56,+8 6.5390 .1394,.3957 -60,* 98 -55,+19 6.6470 055,.7224 -55,-1 TEGDA 10 56 -53,-4 6.861 .2143,.r946 -62 20 74 -57,+'J 7.12r .1969,.r141 -62 30 83 -61,+8 7.373 .t524,.156 -ó1 40 89 -59,+40 7.717 l13l,.tgz -63

25 20 79 -58,0 7.288 .1785,.178U -63 30 87 -59,+5 7.620 .1301 ,.1512 -& 40 9r -62,+47 7.819 .1087,.1391 -67 50 94 -58,+53 8.r23 .0762,.L436 -63

The log storage modulus - temperature plots for the TET2900 methyl acrylate copolymer series (Figure 3.19b) showed a decrease in log E' from a glassy plateau commencing at about -65oC. For those copolymers containing up to 30wtflo methyl acrylate there was a subsequent increase in the storage modulus as the temperature increased up to -30oC, after which the modulus again decreased to give a rubbery plateau after +30oC. Illinger et. aI.27 observed similar behaviour in a linear polyurethane system consisting of MDI and butane diol (BD) hard segments and PTMO 2000 soft segments in the molar ratio MDI :BD: PTMO 2.2:I:1. The increase in storage modulus around -30oC was attributed to the PTMO units partially crystallizing after they had acquired the necessary mobility once the soft segment Tg had been exceeded' Subsequent softening and melting of these crystallites generated in situ 66

9 ooooooo El0 I a.+ OO 6" t a a a a .20 oo a oo a O o30 8 o d o r¡ t àD t.ots o40 trtr o EI '50 oO o 7 og'

6 -r00 -15 -50 -25 0 2s 50 75 100 Temperanre (oC)

Figure 3.19 b Log E' - temperature plots for TET2900 copolymers containing the indicated wtTo methyl acrylate. 6l accounted for the decrease in storage modulus once this temperature was exceeded. Somewhat surprisingly, this crystallization behaviour was not observed for the TET2900 homopolymer. Illinger found that for a polyurethane sample consisting of MDI, BD and PTMO 2000 in'the molar ratio 3 .2:l:I there were no anomalous trends in storage modulus with temperature. It was thought that the smaller content of urethane in the sample which crystallized during the thermal scan was insufficient to form a well developed domain stlucture which could act to reduce the mobility of the PTMO units. Some parallels may be drawn between the effect of domain structure in linear polyurethanes (which can act as physical crosslinks) and that of a three-dimensional crosslinked structure of the type considered here for the photopolymerised urethane diacrylates. Initially copolymerisation of the urethane acrylate portion of the prepolymer with methyl acrylate would result in an overall decrease in network crosslink density which may account for crystallization arising from enhanced pTMO unit mobility. As the wt%o of diluent methyl acrylate increased the overall glass transition of the polymer increased (Tg of poly(methyl acrylate) is +25'C). This effect may compete with the decrease in crosslinking densiry so that no PTMO crystallization is observed at the two highest levels of methyl acrylate incorporation' The tan õ -temperature plot for TET2900 copolymers with TEGDA

(Figure 3.ZOa\ shows two well defined peaks through the whole series' Examination of the storage modulus - temperature plots for this series (Figure 3.20b) again shows increases in storage modulus near -3OoC attributable to crystallite formation, followed by melting during the

thermal scan. 68

Tan ô - temperature plots for the TET2900Æ{DDA copolymer series (Figure 3.21a) again showed two peaks, one at -5OoC and the other progressively increasing in temperature as the copolymer becomes richer in HDDA. Log E'- temperature plots for the copolymer series (Figure 3.21b) only showed an obvious increase due to PTMO recrystallisation for the l\wtVo HDDA sample. The crystallisation observed in log E'-temperature plots for copolymers of TET2900 with TEGDA and HDDA cannot be accounted for by a decrease in the number of crosslinks in the network.

An explanation may be that initially, replacement of some prepolymer with lower molecular weight crosslinker may result in a decrease in crosslinks formed by free radical polymerisation of two double bonds in close proximity to bulky TDIÆ{EA terminal groups. The ability of the crosslinks to act as "tie down" points for the PTMO chains may be dilured when the number of inflexible TDI/HEA groups around

each crosslink point is decreased by addition of methyl acrylate, TEGDA or HDDA. Crystallisation would be expected to be absent in networks with a high proportion of crosslinking comonomer, where soft link mobility after the PTMO glass transition temperature would be restricted by the increasingly glassy state of the polymer. 69

0.3 EI 10

o 20

E 30 cÉ 0.2 .. q) uoÞEE" . ! _3"t""""" o 40 ct :,'...... liii:1*;":""., a 50 ", ";;;;T. -t""t..o' 0.1 " ""at:..::...:to.-.' + t ""'Ï Ô + *Û{d Þ t aa 0.0 -100 -75 -50 -25 0 25 50 75 100 Temperature (oC)

Figure 3.20 a Tan ô - temperature plots for TET 2000 copolymers containing the indicated wtTo TEGDA.

E 10

8.5 O 20 è a o a o T 30 tIl o a o o èo o 40 o EI oE tr r oo I o a E o + 50 o a Eo E t 1.5 ! o gO t l I Eo ++ a o + t t o E coooooooo o ! o aaaoa a o E E E¡¡ Eltr O¡ ooo OOO G¡ o Eto EI E tr EEtr 6.5 -100 -15 -50 -25 0 25 50 75 100 Temperature (oC)

Figure 3.20 b Log E' - temperature plots for TET 2000 copolymers containing the indicated wtTo TEGDA. 70

0.3 E l0

O 20

E¡ tr E 30 c{ E o.2 EOstutro c) c ! t:l,,tr_.- 40 .t..... a GI ..... 50 + E o o o o r r' "^'.i".".i"ir.]krtÍ + *' o ir! o. O oÖó 0.1 9a.1*o"ooocoot_-.... t t """"t t t o II ô .r ö i o O! Itr o Ea o tr o c EI oE Es a ¡è o Eltr o I l.?$ trg 0.0 lm -75 -50 -25 0 2s 50 75 100 Temperature ("C)

Figure 3.2L a Tan õ - temperature plots for TET 2900 copolymers containing the indicated wtTo HDDA.

9.5 810 s5queflq .20 !30 iÈ -tql;';";"''''ç.'r..-. ri 8.5 o40 à0 E E 6 '50 t t t a a rì. r, a aa a 7.5 ." ,,:" ";. ooooooo "ttt-ttt;Ï EEIE!!! strtrEoEgog aooaooo soEtrEEtr

6.5 -100 -15 -50 -25 0 25 50 15 100 Temperature (oC)

Figure 3.21 b Log E' - temperature plots for TET 2900 copolymers containing the indicated wtTo HDDA. 7l

3. 2. 2(b)TPCL copolymers (i) TPCL2000 Results of thermal scans for TPCL2000/20wt%o TEGDA and TPCL2000/20wtVo HDDA copolymers are shown in Figure 3.22. The copolymer containing20To TEGDA showed two broad peaks in the tan ô- temperature plot while the corresponding TPCL2OOO/HDDA copolymer showed poorer peak resolution. A slight shift in the PCL soft segment Tg (Tgr) toward a higher temperature was apparent in the HDDA copolymer sample. Although this shift was small (7"C), it may indicate a degree of comparibility with the HDDA diluent not present in PTMO based urethane

acrylate copolymers. The HDDA transition at +1oC also occurred at a lower temperature than that for the TET2000/HDDA copolymer (+1OoC), indicating greater interactions between the two components. The shift in PCL Tg for TEGDA copolymer was only 4oC, with a shift to lower temperature of the TEGDA tan ô peak (TgZ) being of comparable magnitude to that observed for the HDDA peak when compared to the TET2000/2}wtVoTEGDA analog (Tgz=+1OoC for the TPCL 2000 based { copolymer and Tgz=+21oC for the TET2000 based copolymer). (ii) rPcL3000 DMTA scans for copolymers of TPCL 3000 with 20wtVo TEGDA and HDDA are shown in Figure 3.23. Observations made for the TPCL 2000 copolymer plots are generally valid for these copolymers i.e. the peaks in the TPCL3000/TEGDA copolymer were generally better resolved than those for the HDDA copolymer. The Tg2 peak positions are identical to those seen for the TPCL 2000 copolymers. The most significant difference between the TPCL2000 and 3000 copolymers was that tan ô (Tg) for the TPCL3000 Tg1 peak

decreased to a smaller extent than that for the TPCL2000 copolymers. 72

0.5 E HDDA . TEGDA 0.4 oQo o o o o homopolymer o o o o 0.3 o cq o q) o o 0.2 o cl o o o o o ô o O a 0.1 o

0.0 -100 -75 -50 -25 0 25 50 75 100 Temperature (oC)

Figure 3.22 Tan ô - temperature plots for TPCL2000 copolymers containing 2O wtVo TEGDA and HDDA.

0.4 h trO o homopolymer tr O HDDA 0.3 s O TEGDA EI cl q,) 1ó 0.2 o

cg o a q o o 4t 31.ao o o 0.1 o tq, û q#b Eat o 'oa %

0.0 -lm -75 -50 -25 0 25 50 15 100 Temperarure (oC)

Figure 3.23 Tan ô - temperature plots for TPCL3000 copolymers containing 20 wt%o TEGDA and HDDA 73

Minor upward shifting of the Tg¡ peak in TPCL30O0/20wt7oHDDA was observed

a))lc\TtrTIPTMOI 2 nnnnl vrnerq

(i) TET( 1000)2 copolymers. Tan ô plots for TET(1000)2 copolymerised with 2jwtVo TEGDA and HDDA are shown in Figure 3.24. Poor resolution was obtained for both TET(1000)2 co TEGDA and HDDA copolymers with the chief observation being the considerable decrease in tan õ (Tg1) from 0.43 in the TET(I000)2 homopolymer to 0.23 in the TEGDA and HDDA copolymer samples. Small decreases were also observed in the PTMO soft segment Tg1 . from -32"C for the homopolymer to -35 and -34oC for the TEGDA and HDDA copolymer samples.

(ii)TET(2000)2 copolymers Tan ô plots for TET(2000)2 co 20wt7o copolymers with TEGDA and HDDA are shown in Figure 3.25. Good resolution was obtained for the PTMO Tg1 peaks in the range -50 to -60oC, however the copolymer waS extremely rubbery after OoC, leading to some data scatter with increased temperature. The main difference between the TET(1000)2 copolymers and those for TET(2000)2 lies in the fact that the tan õ (Tgr) is affected to a lesser degree in the TET(2000)2 copolymers, decreasing from 0.53 in the homopolymel to 0.35 and 0.21 for the TEGDA and HDDA copolymers respectively. Two factors may be responsible for the observed changes in the value of tan õ (Tg) in the PTMO peak : changes in network crosslink density, or differences in compatibility of the comonomer with the PTMO soft segment. 14

0.5 s homopolymer s% O HDDA 0.4 g E¡ o O TEGDA tr tr 0.3 cË tr €q) E

cq 0.2 Ð b 0.1 j'. o

0.0 100 -15 -50 -25 0 25 50 75 100 Temperature ("C)

Figure 3.24Tan õ - temperature plots for TET(1000)2 copolymers containing 2O wtVo TEGDA and HDDA.

0.6

% El homopolymer 0.5 tr tr O HDDA

o O TEGDA 0.4 E (! o tr oo q) o EI õ 0.3 Gb loo cÉ oo 0.2 a O O oo 0 OO o o 0.1 Eb

0.0 -100 -75 -50 -25 0 25 50 15 100 Temperature (oC)

Figure 3.25 Tan ô - temperature plots for TET(2000)2 copolymers containing 20 wt%o TEGDA and HDDA. 75

The dynamic mechanical properties for this copolymer series appear to lie between those for compatible systems where a single composition dependent combined o peak is observed and completely incompatible systems where each phase exhibits its own characteristic relaxations and the cr process is independent of composition.

Similar decreases in high temperature loss peaks in a two phase system have been observed for rubber toughened epoxy resins 28, with other workers 29 correlating the decrease in the epoxy glass transition temperature from that of the unmodified epoxy resin with the extent of phase separation between the rubber and epoxy phases. Rubber which did not phase separate during thermal curing remained dissolved in the epoxy phase, plasticising the epoxy glass transition. The Tg for higher temperature peaks varied in a non-linear fashion with composition as shown by the Tg2 peaks due to TEGDA and HDDA in TET 2000 and 2900 networks (Figures 3.26 and 3.27).

Lee and Park 30 investigated the dynamic mechanical properties of a series of sequential interpenetrating polystyrene/polyurethane networks, in which styrene was photopolymerised in the presence of a

MDI/PTMO1000 urethane network after the latter had been formed by a thermal condensation reaction. Photopolymerisations were carried out at

OoC and 40oC - in general networks prepared at 40"C had PMTO1000 loss peaks wíth greater tan ô (Tg) and lower glass transition temperatures compared to the networks formed at OoC. This was attributed to the rate of network formation exceeding that of phase mixing at the higher temperature. Jingjiang et.al. 3l examined copolymers of poly(propylene oxide) urethane acrylates with styrene, vinyl acetate and methyl methacrylate. Copolymers showed two tan ô peaks - the lower temperature peak was 76

80 a

60 a a

I' EI 40

è0 tr F a ?r

o g TEGDA 0 . HDDA g

-20 0 20 40 60 80 100 WtTo in copolymer

Figure 3.26 Glass transition temperatures of TEGDA and HDDA copolymers as a function of wt%o in TET2000 neworks.

80 a

60 o o 40 tr o

ào t- 20 E TEGDA EI a E O 0 o HDDA I

-20 0 20 40 60 80 100 WtTo in copolymer Figure 3.27 Glass transition temperatures of TEGDA and HDDA copolymers as a function of wtTo in TET2900 networks. 77 taken to be that due ro the polyurethane phase strengthened by plastic and the other peak was attributed to the plastic phase plasticised by polyurethane. Compositions of both phases were calculated by assuming that the contents of any phase in a two phase system could be calculated.according to the Fox equation:

lWtW2 Tg=Tgt-Tg, (3.2)

where W1+W2=1. Tg represents the glass transition temperature of the phase being considered and W¡ and W2 are the weight fractions of each component in the phase. Tgr and TgZ are the values for the homopolymers of both components. Table 3.8 shows estimates of the molecular composition of each phase obtained by applying Equation 3.2 to a number of urethane acrylate copolymers with TEGDA and HDDA, with W1 referring to the urethane soft segment and V/2 referring to the TEGDA or HDDA copolymer. For all copolymers the copolymerisation had little effect on the temperature of the soft phase Tgt. For the TET2000 copolymers the decrease in Tg2 from the HDDA homopolymer Tg was proportionately greater than that for the TEGDA copolymer. Consequently, according to Equation 3.2, the quantity of PU phase mixed in with the HDDA phase is greater than the PU phase mixed with the TEGDA. The composition of the plastic phase for the TET29QO/TEGDA copolymers is somewhat surprising - up to and including 3Owt%o TEGDA the PU composition in the plastic phase exceeds that found for TET 2000. This behaviour is unexpected, since if TEGDA is less compatible with 78

PTMO chains than HDDA, compatibility would be expected to decrease with increasing PTMO chain length.

TABLE 3.8 Estimates of molecular compositions within soft (polyurethane) and plastic phases for urethane acrylateÆEcDA (I) and HDDA (H) copolymers, calculated from Equation 3.2. W1 refers to the weight fraction of polyurethane soft segment and W2 refers to weight fraction of glassy copolymer present in the phase.

Copolymer Wr Wz V/l Wz TET2Uru/IOT 0.97 0.03 0.ó0 0.4u 207 1.00 0.00 0.78 o.22 307 1.00 0.00 0.85 0.15 407 1.00 0.00 o.94 0.0ó 507 1.00 0.00 0.96 0.04 TET2000/l0H I.00 0.00 0.40 0.óu 20Il 0.99 0.0r 0.s9 0.41 30H I.00 0.00 0.ó9 0.31 40H 1.00 0.00 0.88 o.T2 50H l.0Ll 0.00 0.91 0.09 'rE12900/r0'r' 1.00 0.00 0.55 0.45 207 1.00 0.00 0.62 0.38 30r t.0u 0.00 0.67 0.33 407 1.00 0.00 0.93 0.07 507 I.00 0.00 0.94 0.06 TET2000/10H 1.00 0.00 o.47 0.53 20H I.00 0.00 0.51 0.49 30H 1.00 0.00 0.55 0.45 40H r.00 0.00 0.E4 0.16 50H 1.00 0.00 0.88 o.rz 'ïPCLZ000|20'L' 0.95 0.05 0.54 0.46 TPCLZV¿/aOH 0.91 0.09 0.34 0.óó TPCL3OOO/2OT 1.00 0.00 0.61 0.39 TPCL3OOO/2OH 0.96 0.04 rJ.42 0.58

The observed position of the higher temperature peak in the tan ô-

temperature plots may result not from the compatibility of TEGDA or HDDA with the TET2000 and 2900, but in the effective volume fraction of "pure" soft phase. Important differences exist between the crosslinked networks and SINs studied by Jingjiang and Lee and the urethane acrylates considered here. Jingjiang 3l observed Tg shifts for the poly(propylene oxide) soft phase of up to 22"C, whereas no comparable shifting of the soft phase 79 was observed in either the TET or TPCL urethane acrylates, indicating that the two copolymer components af e es se ntially immiscible. Comparisons of tan õ(Tg) for the soft phase in the TET2000 and TET2900, TPCL2000 and TPCL3000, and the TET(1000)2 and TET(2000)2 copolymers all show that the Tg1 peak height is decreased to a much lesser extent in the TET2900, TPCL3000 and TET(2000)2 copolymers. Iricreasing crosslink density appears to decrease the volume fraction of soft phase - in polymers with a larger soft phase volume fraction, more effective plasticisation of glassy copolymers can occur, Ieading to a larger downward shift in the glassy copolymer Tg2 peak-

3.2.3 Water in Urethane Acrylate Networks The dynamic mechanical properties of wet and dry urethane acrylate homopolymers are shown in Table 3.9, together with the sample EWC. DMTA plots for wet and dry polymers are shown in Figures 3.28a-j. For the TET650 polymer there was a l2"C dectease in Tg for 3-LVo EWC, with the height of the tan õ peak showing a 5Vo increase compared to that of the dry polymer. The p peak at -70oC in the dry polymer increased in magnitude in the wet sample. The origin of this peak in the dynamic mechanical spectrum will be considered later. The Tg of the TET1000 polymer remained unchanged; the tan õ peak height, however,

showed a 257o increase over that of the dry polymer. DMTA traces for wet TET2000 and 2900 homopolymer showed considerable increases in the glass transition temperature. In addition to the shift in Tg an additional higher temperature peak was observed for both wet polymers. Increases in Tg and modulus of polymers with solvent uptake are termed antiplasticisation, with antiplasticising solvents 32 possessing the ability to increase the order in polymer chains, resulting 80

0.6

o dry ooo 0.5 o ùE a wet oEo EI 0.4 o oEl a GI EI tr q) OE o € 0.3 aE E O GI otr tr 0.2 Otr o E a E 0.1 O tr o O o o 0.0 100 -75 -50 -25 0 25 50 75 100 Temperature (oC)

Figure 3.28 a Tan ô - temperature plots for dry and saturated TET 650 polymer

0.5

dry oO 0.4 oo o o wet o o -otr o og Eoqo ñ 0.3 E q) Èo E' E tr GI Ea cl 0.2 o û EIù tr o 0.1 tr 0 o s o ch** oo

0.0 100 -15 -50 -25 0 25 50 15 100 Temperature (oC)

Figure 3.28 b Tan ð - temperature plots for dry and saturated TET 1000 polymer 8r

0.4

oo oo E¡ dry rO O Oo o OOOO O wet 0.3 o F o Eo %" oo¡oo g oa

6É Eþ o q) 0.2 1t E, a cl t O o 0.1 È o .tÈ

0.0 -100 -15 -50 -25 0 25 50 75 100 Temperanrre (oC)

Figure 3.28 c Tan ô - temperature plots for dry and saturated TET 2000 polymer.

0.4

E d.y o wet 0.3 ffi_ o trh^o o o o d tt-*.""%" q) õ 0.2 , ¡oJ"i"*; * o rt cÉ o o o rO o o taa o 0.1 !t oa) ¡O a oooooo E

0.0 100 -7s -50 -25 0 25 50 15 100 Temperature ('C)

Figure 3.28 d Tan ô - temperature plots for dry and saturated TET 2900 polymer. 82

0.4 .ttJff o g'o tr otrotr tr dry 0.3 t"ttt' tEoo' o wet og-o t c! oE¡ q) 0.2 ^E'E " õ a EOE É GI oo o s aE o E 0.1 oEt o "" o o o o o oo O 0.0 100 -15 -50 -25 0 25 50 75 100 Temperanre (oC)

Figure 3.28 e Tan ô - temperature plots for dry and saturated TPCL1250 polymer.

0.5

tr dry 0.4 .q o wet ." %* o G c! 0.3 tr o ag q) õ o ¡E o t"t cÉ 0.2 Otr E o o O O E¡ o 0.1 T o

0.0 -100 -15 -50 -25 0 25 50 7s 100 Temperature (oC)

Figure 3.28 f Tan õ - temperature plots for dry and saturated TPCL2000 polymer 83

0.4 rÈ ooG a E (È¡ dry 5È o wet 0.3 t t cq q) a õ 0.2 cl C

0.1 tra tr tog 0.0 100 -15 -50 -25 0 25 50 15 100 Temperature (oC)

Figure 3.28 gTan ô - temperature plots for dry and saturated TPCL3000 polymer

0.6

aao E dry 0.5 .t"ry" otr o wet a "" 0.4 E o o o cú td E o E (u o ! 0.3 o o Otr E¡ a c! % ¡E o E 0.2 a % EI a E o o oo Eb oE o ft 0.1 ¡tr ¡ otoooo. aaaaoa

0.0 -lm -75 -50 -25 0 25 50 't5 100 Temperature (oC)

Figure 3.28 h Tan ô - temperature plots for dry and saturated TET (650)2 polymer 84

0.5 o a o s%o E dry 0.4 oEE E û o wet 5 c I cl 0.3 T q) E I o \ GI 02 o EI d qq- t oiþ- o^.1r_ 0.1 -.Þq*rff o o htu" t 0.0 -100 -75 -50 -25 0 25 50 75 100 Temperature (oC)

Figure 3.28 i Tan õ - temperature plots for dry and saturated TET (1000)2 polymer.

0.6

% E dry 0.5 E o wet

0.4 tr

cq tr q) o õ 0.3 Eb

cú 0.2 o tt*t' oo ooo oo o o aao o aa OO o 0.1 a o o Eb

0.0 -100 -15 -50 -25 0 25 50 75 100 Temperature ("C)

Figure 3.2S j Tan ô - temperature plots for dry and saturated TET (2000)2 polymer 85 in tighter packing, which in turn results in a decrease in available free volume. It seems unlikely, however, that water, which plasticised TET650 polymer, would be able to act in an opposing manner for the

TET2000 and 2900 polymers.

TABLE 3.9 Dynamic mechanical properties of dry and satu¡ated urethane acrylate polymers. 'lg("c) Polymer EWC(7o) (Tg) wtrz("c) LogE' tan ô (tg+40'C) TET65O 0 +35 0.462 54 6.972 3.I +23 0.5I6 52 6.921 lEt 1000 0 +13 tJ.3Z9 82 7.030 2.8 +14 0.430 68 6.764

'tE12000 0 -50 0.'302 88* 7.392 2.0 -32 0.350 7. r00 +34 0.263

'lï'12900 0 -50 0.307 t3 1* 1.368 2.0 -32 o.23r 7.3t4 -4 0.286

0 +8 0.371 70 ó.854 r.8 -4 0.'373 56 7.058

0 -26 0.405 54 6.97 5 I.6 -26 0.413 46 6.895

0 -31 0.387 44 7.493 1.5 -37 0.363 48 7.168

TET(ós0) 0 -ó 0.498 ó0 6.s33 3 -15 0.542 42 6.114

TET(100( 2 0 -32 0.428 50 ó.981 2.5 :34 tJ.494 36 6.813

0 54 0.534 25 7.935 z.l -'34 0.t79 74 1.424

* denotes assymmetric peak - denotes peak insuff,rciently resolved for W1p to be determined. 86

The Tg for TPCLL250 decreased l2oC for a l.8Vo EWC sample with no accompanying change in peak height. Glass transition temperatures and tan ô peak heights at Tg showed little change from the dry polymer for saturated TPCL2000 and 3000 samples. The Tg of TET(650)2 decreased 9oC for a sample containing 3%oEWC, together with a slight increase in tan ô(Tg). The Tg of

TET(1000)2 remained unaltered on sorption with an increase in the tan ô peak height again being observed. The dynamic mechanical properties of TET(2000)2 polymer \¡/ere drastically affected by water uptake, with the previously sharp Tg peak at -54oC giving way to a broad transition barely one third of the original peak height. Glass transition temperatures in wet polymers can be predicted, either in terms of free volume theory using the Kelley-Beuche equation (3.6¡f f , or the Fox equation(3.Ð 34 '

crpVpTg(p) +aw(1 -Vp)Tg(w) Tg (3. 6) GpVp +crw(l-Vp)

I Wp Ww (3.2) Tg re(p) where Tg(p) and Tg(w) are the glass transition temperatures of polymer and water respectively. V, W and cx, are the volume fraction, weight fraction and expansion coefficients of the two components. The simpler

Fox equation was used to calculate Tgs for those samples in which a decrease in Tg was observed (i. e. plasticisation was apparent), with the glass transition temperature of water being set at - 134K 35. 87

TABLE 3.10 Results of the Fox equation for water plasticised urethane acrylates. The Tg of water was set at -134K. 'I'g(K) Polymer Mass 7o water Experimental equation Tg(K) 296 297

TPCLI25O 1.8 269 275

TET(650)2 3.0 258 260

Calculated Tgs for TET650 and TET(650)2 showed good agreement with those obtained experimentally, while the calculated Tg for TPCL

L25O was 7oC high. Insignificant water plasticisation occurred for several of the urethane acrylate polymers. Despite the low EWC of the netwôrks, absorption of comparable amounts (l-3Vo by weight) of water has been found to cause much larger reductions in the Tg of epoxy resins 36-38. Several possible mechanisms for plasticisation of crosslinked polymer networks have been proposed. One mechanism involves the replacement of polymer-polymer hydrogen bonds between polar functional groups 39 with water-polymer bonds, effectively reducing the network crosslink density. Other workers 40, however, maintain that the plasticising effect of water in epoxy resins can be explained by an increase in free volume decreasing the activation energy for the Tg proces s.

Agreement of saturated polymer Tgs with the Fox equation implies that water is intimately associated with motional units in the polymer, rarher than being isolated in droplets or clusters. The lack of any Tg shift for the wet TET1000 polymer compared to that observed for the saturated 88

TPCL 1250 sample also suggests that the nature of the soft segment has an effect on water distribution in the polymer. A number of studies on the distribution of water in hydrogel networks suggest 4l-41 that the water in the hydrogel exists in a continuum of states between two extremes. Water strongly associated with the polymer network is referred to as "bound" or non-freez\ng water with water unaffected by the polymeric environment being referred to as "free" or freezing water. Dist¡ibution of water in different environments may also occur in polymers with much lower water contents than hydrogels : Moy and Karasz 48 studied plasticisation in wet epoxy resins and observed no water melting endotherms. Such behaviour may be explained either by water in the polymer being present in droplets small enough to supercool, thus giving no melting endotherm, or by water being strongly associated with polar groups in the polymer forming a non-freezing phase. The state of water may also influence the mechanical properties of the polymer. DSC studies on PHEMA 42 showed that above 207o EWC water existed as a "bulk-like" phase capable of freezing in a cooled polymer. Torsion pendulum data on PHEMA hydrated to low water contents 26 showed initial plasticisation of the glass transition in accordance with the Fox equation. The Tg of 0"C found for a fully hydrated pHEMA sample, however, was 48"C higher than that predicted. This change in dynamic mechanical behaviour may be due to the formation of ice in the polymer; the sudden decrease in log storage modulus in hydrated pHEMA samples after OoC has been attributed 49 to the melting of crystalline water. The urethane acrylate polymers most affected by water are those with PTMO soft segment molecular weights of > 2000. Tan ô (Tg) increased for the wet TET2000 polymer while the Tg itself showed an 89

18"C temperature increase. Saturated TET2900 showed both a decrease in tan ô (Tg) and an increase in Tg from the dry polymer, possibly indicating a greater degree of ice formation in the polymer compared to TET2000, either as the sample was cooled to the DMTA run start temperature or through the run itself. The most drastic effects were observed for the sample with the lowest crosslink density of any hydrated sample considered - the TET(2000)2 polymer : presumably ice formation in this sample was extensive and effectively dampened the motions of the PTMO chains. From the results obtained here it is possible to classify water in

these polymers as being in three states:- l.Intimately associated with the motional units, plasticising glass transition temperatures and increasing polymer free volume. 2. Moderately associated with motional units - increasing free volume, but not influencing Tg. 3. Poorly associated with motional units - forming clusters or droplets of differing sizes, which in some polymers are large enough to a permit ice formation either on cooling or during thermal scans.

The increasingly hydrophobic nature of the PTMO polymer chains with increasing molecular weight appears to lead to decreased plasticisation of polymer glass transitions, but ultimately to water clustering, which may result in the formation of a water phase capable of freezing. The decreased hydrophobicity of the poly (caprolactone) diol soft segment compared to PTMO may be manifested by water in TPCL polymers not forming isolated clusters, but being moderately associated with PCL soft links, although clearly the affinity of water for the PCL soft links cannot be great - as evidenced by the low water uptake in the

TPCL polymer series. 90

Solvent Induced Transitions (T¡g) In addition to the low temperature transition at -120oC and the glass transition peak, an additional low temperature peak develops on water sorption into the urethane acrylates. This peak is termed T dil, and appears to be present to a limited extent in the TET650 homopolymer. Copolymers of epoxy bis GMA and tetraethylene glycol dimethacrylate showed a peak at -110oC (lHz) 50 which increased in magnitude for higher tetraEGDMA contents. Dynamic mechanical plots of the subambient temperature region are shown in Figures 3a-c for TET650,

1000 and 2000 polymers. T ¿it peaks also appeared between -80 and - 70'C in wet TPCL and TET(PTMO)2 polymers. Koshiba et. al. 10 observed a similar peak at -60oC in an isophorone diisocyanate/PTMO1000 based polymer despite storing samples over silica gel in a dessicator. The higher temperature of this peak may be due to the frequency of measurement used (110H2). In common with results obtained 5l for a homologous ethylene i¡ glycol dimethacrylate series the T, peak decreased as the T dil peak appeared in saturated samples.

The growth of T ¿¡1 peaks in linear methacrylates has been ascribed to the interaction of the penetrant molecules with ester side groups in the methacrylate. Disruption of polar polymer - polymer interactions on solvent sorption and subsequent replacement by solvent - polymer interactions has been invoked 52 to account for the observed increase in side chain mobility. The temperature of the 1 peak in pHEMA networks has been found to be independent of the length of the side chains 2 confirming that the y peak is due to localised internal motions in the side chains. The urethane group undoubtedly can hydrogen bond to water absorbed into the polymer 4'9 and since the urethane groups occur in thc 91

0.06

0.05 ooooooo o cl 0.04 o o q) õ o o É 0.03 E o G' E oEGtEEtrEg .stcftE E O O 0.02 ¡ OO tr d.y o 0.01 wet

0.00 -r20 -110 -100 -90 -80 -70 -60 -50 Temperarure ('C)

Figure 3.29 a Subambient tan ô - temperature plots for dry and saturated TET 650 polymer.

0.045 e O OE o

cÉ o 0.035 EI q) ! a E GI g o qtr 0.025 ou E dry o o f a g-t o wet o

0.0r5 t20 -u0 -100 -90 -80 10 {0 -50 Temperature (oC)

Figure 3.29 b Subambient tan ô - temperature plots for dry and saturated TET 1000 polymer. 92

0.06 dry

o wet O 0.05 o cÉ q) o It É 0.04 c! a

0.03 o E a E o o o o o oE

0.02 -r20 -r l0 -100 -90 -80 't0 Temperature (oC) { Figure 3.29 c Subambient tan ô - temperature plots for dry and saturated TET 2000 polymer. 93 side chain localised side chain motions are affected on sorption of water into the network.

3.3 Summary Glass transition temperatures showed large decreases.*ith both increasing soft segment molecular weight in the TET and TPCL urethane acrylate series, and decreasing crosslink density, exemplified in the trends in Tg found for the TET(PTMO)2 series. Crystallinity only affected the amorphous glass transition for one sample, with an increase in Tg and broadening of the tan ô peak being

observed in the TET(2900)2 polymer. Predictions of polymer glass transition temperatures, based on the TDIÆIEA hard segment Tg (obtained from DMTA analysis of model polymers) and soft segment Tgs were made using the Fox equation. The calculated values considerably underestimated the actual polymer Tgs, particularly for the urethane acrylate networks with higher crosslink densities.

Much closer agreement of calculated and experimental polymer Tgs was found by using a modified form of DeBenedetto's equation, in which it was assumed that the observed Tg reflected only changes in soft segment mobility with crosslink density. Higher temperature transitions, often fou.nd in linear polyurethanes, were also not observed in any DMTA

runs on the urethane acrylates. This, together with the success of the crosslinking model in predicting polymer glass transition temperatures, suggests that chemical crosslinking in the urethane acrylates inhibits TDI hard segment

aggregation, preventing the formation of a discrete hard phase. 94

Dynamic mechanical behaviour for copolymers of the urethane acrylates with low molecular weight acrylate diluents resembled that for rubber modified epoxy resins, with considerable temperature decreases in loss peaks due to the glassy diluent polymer indicating plasticisation by the rubbery polyurethane. Uniform effects of water sorption on homopolymer glass transitions in saturated urethane acrylates were not observed, with some saturated polymers showing Tg decreases readily explicable in terms of water plasticisation, while others showed no change in dynamic mechanical properties compared to the dry polymer. Drastic changes in mechanical properties for another group of saturated urethane acrylates were attributed to ice formation during the DMTA scan. The differences in the effects of water sorption on the loss peaks in the polymers may be due to water showing a greater tendency to form clusters in networks with a low proportion of polar, water solubilising groups. 95 CHAPTBR FOUR

DIFFERBNTIAL S CANNING CALORIMETRY

4.L Intro duction Differential scanning calorimetry was used to compare the thermal behaviour of crosslinked urethane acrylates with that of linear polyurethane networks. The ability of DSC to detect a wide range of melting, crystallisation and glass transition phenomena with high sensitivity makes the technique particularly applicable to the analysis of polyurethanes, in which a multiplicity of possibte phases (both the hard and soft phases.can be either crystalline or amorphous) may exrs t

In addition to the urethane acrylate homopolymers and their constituents, DS C thermograms of glassy hard segment model polymers, synthesised using the techniques outlined in Chapter Two, were recorded. These polymers were Scanned to determine pure hard segment glass transition and melting behaviour for comparison with homopolymers containing varying hard segment weight fractions.

Po.lyurethanes containing soft and hard segments may exhibit two phase behaviour resulting from thermodynamic incompatibility between the segments. Phase separation is usually incomplete, however, and investigators have detected the presence of both hard 2'3' segments in the soft phase I and soft segments in the hard phase Annealing the polyurethane at high temperature, followed by rapid quenching to below the soft segment glasS transition temperature can 96

"freeze-in" a phase mixed morphology 4,1,2, with the presence of dissolved hard segment in the soft phase resulting in an elevation in the soft microphase glass transition tempe¡ature. One urethane acrylate homopolymer, with a well defined glass transition, was subjected to annealing/quenching cycles in order to determine whether similar phase mixing could be induced.

4.2 Results and Discussion 4.2. L Homo polymer Thermograms 4.2.1. (a) TET homopolyr!çr! Initial DSC runs were made on the polytetramethylene glycols from which the TET series prepolymers were synthesised in order to assist in identification of peaks in homopolymer thermograms. The PTMO diols were quenched from 30'C to -100"C at 200K/min and then thermally scanned (Figure 4.1). No amorphous transitions were detected directly' The PTMO650 diol showed an exotherm neal -7 4"C, which was probably due to the partial crystallisation of PTMO units after they had acquired the necessary mobility once the soft segment Tg had been exceeded. The presence of this exotherm suggests that the PTMO650 diol is not completely crystalline, with an amorphous Tg below -J4"C. The crystallisation and melting peak temPeratures, together with the crystallisation, AHc, and melting, AH¡¡, enthalpies for each of the PTMO

diols are listed in Table 4. 1. It is apparent that there are two melting peaks for each diol, with the lower temperature melting peak appearing as a pronounced shoulder for the PTMO 650 (n=9) and the PTMO 1000 (n=14) samples. While the melting peaks are of similar size in the PTMO 2000 the higher tempefature melting peak dominates the thermogram for the PTMO 2900 sample' The 97

(b)

(c)

m l o- (d) 9 J (D a õ'

I -l 00. 0 -75. 0 -50. 0 -25. 0 û-0 25. 0 50.0 75. 0 loo- 0

.l Tcopcroturc ('C)

Figure 4.1 Thermograms of poly(tetramgqlylene oxide) diols (a) PTMO650 (b) F[MO1000 (c) PTMO2000 and (d) PTMO2900. 98 higher temperature melting endotherm may be due to a slower reorganisation of crystalline regions in the polymer yielding more perfectly formed crystallites which have a higher Tr 5.

TABLE 4.1 Thermal properries of PTMO diols, TET PTMO urethane acrylate prepolymers and polymers.

Sample 1-" ACp(J/g-C) T"EC)T (J/e) T-(OC)Y (íe) Þ ^Hc ^Hm PTMO diols ruMO650 -13.7 8.9 3.0 86.8 rtM0lux) 9.7 89.1 PrrMo2m0 18.9 61.1 37.0 38.4 PTMO29OO 18.9 29.5 40.4 93.3

TETprepolymers IET65O -ó5.ó 0.034 -32.t) 5.1 tEI'1000 -50.6 0.369 -12.8 5.2 14.2 8.0 fET2m0 -70.6 0.45ó -30.3 2r.6 t I.l 38.9 35.6 19.2 IET29OO -68.0 -43.0 14.3 Ió. I 65.5

TETpolymers fET650 tET1000 -53.7 o.24L IETzOOO -72.9 rJ.z59 lb12900 -15.r 0.354 -43 8.6 -0.1 tt.2 20.0 0.3 not present ln f refers to crystallisation peak. Y refers to melting peak

It could be argued that since these samples were quenched from

30oC, a temperature below that of the melting endotherm around 40"C, crystallites existing prior to the thermal scan may inhibit the development of an amorphous phase. The observation that the PTMO 1000 sample, with a melting peak below 30"C, showed no glass transition or crystallisation behaviour, suggests that quenching these diols to a fully amorphous state may be difficult. Schneider et.al-6 obtained the following Tgs by quenching diols from the melt at 99

320K/min: PTMO1000 : 185K, PTMO 2900: 190K, but still cautioned that crystallinity may have affected the observed transition temperature.

Amorphous glass transitions were found for urethane acrylate prepolymers. Thermograms for these are shown in Figure 4.2, with the glass transition, crystallisation and melting peak temperatures and their accompanying enthalpies listed in Table 4. 1. The development of crystallisation peaks for the prepolymers showed that the addition of the TDI/HEA moiety to the terminal OH groups on the diols reduced the tendency of the PTMO chains to crystallise. There was no observable trend in glass transition temperature with soft segment molecular weight - the fact that for all prepolymers AHm >ÂHc indicates that the samples were not fully amorphous at the commencement of the thermal scan. The similarity in Tg for the TET650 and TET2900 prepolymers suggests, however, that there is little variation in Tg in the PTMO molecular weight range considered here. This corresponds to region I in Tg versus log M. plots obtained by Cowie for a number of linear polymers 7 where the glass transition temperature of a polymer chain is constant once it paSseS from oligomer to polymer when the chains become long enough to be able to adopt a gaussian coil conformation. Thermograms for the homopolymers cast from the prepolymers are shown in Figure 4.3. There is a pronounced decrease in manifestations of crystallisation and melting behaviour in the homopolymer series compared to their respective prepolymers but, with the exception of the TET650 homopolymer, in which no Tg was detected, the glass transition 100

(a)

(b) m f o- o J (D 3 c¡

(c)

(d)

100.0 -r00.0 -75- 0 -50_ 0 -25. 0 0-0 25. 0 50- 0 75. 0 Teçorotre ('C)

(a) Figure 4.2 Thermograms of TET P'[MO based ufethane acrylate prepolymers rsTeSO G) TETI000 (c) TET2000 and (d) TET2900- (b) m f o_ o J í) I õ'

(c)

(d)

-100.0 -?5. 0 -50. 0 -25- 0 0.0 2s. 0 50. 0 75. 0 100.0 Tcmponoturc ('C)

Tigfffli Thermograms of TET u¡ethane acrylate polyme.s (a) TET650 (b) TETI000 (c) TET2000 and (d) T8T2900. t02 temperatures of the homopolymers are within 4-6"C of the prepolymer glass transitions. Tgs obtained for the TET1000 and TET2000 are similar to those obtained by other workers 8 with the values obtained in the earlier study being about 6oC lower. While the ÂHm value for TET2900 (ll.2J/Ð is sor-newhat greater than the ÂHç value (8.6J/Ð, the size of these peaks and the fact that a flat baseline between them could not be assigned, means that the amorphous or crystalline nature of the polymer prior to the scan cannot be unambiguously established. NMR results (Chapter Five) obtained for the TET2900 polymer suggest, however, that significant crystallinity does not develop in the polymer within a period of one week. 4.2.1 (b) TPCL homooolvmers Thermal transitions and peak enthalpies for poly(caprolactone)diols are shown in Table 4-2. Thermograms showed no trace of either amorphous Tgs or recrystallisation peaks, with one large peak in the range 55-60'C dominating the DSC trace (Figure 4.4)

T ABLE 4.2 Thermal properties of poly(caprolactone) diols and TPCL uethane acrylate polymers.

Sample rs lCp(J/e-C) T- ¡sn' (J/g) PCt-diols PCLI25O 21.4 t4.42 54.83 78.57 PCL2OOO 59.36 92.37 PCL3OOO 59.9ó 97.'31 IPClpolymers

IPCLI25O -37.3 tJ.245 41.6 o.r7 TPCL2OOO -50.6 0.338 43.9 0.61 IPCL3OOO -54.9 0.509 3t.z 3.34 - denotes peak not present rn samp Y refers to melting peak 103

(b)

m l o_ o J (D 3 r)' (c)

lo0. o -75. 0 -50. 0 -25. 0 0.0 25. 0 50- 0 75. 0 lo0- 0 Tcnpoeoture ('C)

Figure 4.4 Thermognms of poly(caprolactone)diots (a) PCLI250 (b) PCL2000 and (c) PCL3000. 104

Minor peaks were observed in the 25-30'C region; except for the PCLl250 sample the enthalpies of these melting peaks w ere in s ignifican t.

The TPCL polymer series (Figure 4.5) showed glass transition temperatures generally 15 to 20oC higher that the equivalent MW PTMO urethane acrylates, indicative of the less flexible nature of the soft link in these materials. Recrystallisation during thermal scanning was not observed to any great extent in this series, although a very small, broad peak may be discerned in the TPCL3000 thermogram near -10oC. Small melting peaks are present in all homopolymers of this series-the TPCL 3000 homopolymer shows the largest melting endotherm with the decreased melting temperature compared to that for the TPCL1250 and 2000 polymers indicating that the crystallites are poorly ordered.

4.2. 1. (c) TET(PTMO)2 homopolymers. Thermal data for the TET(PTMO)2 series is summarised in Table

4.3, with DS C thermograms for the series excluding TET(2900)2 shown in Figure 4.6.

TABLE 4.3 Thermal properties of TET(PTMO)2 urethane acrylate polymers.

Sample I'Ob ACp(J/e-C) TcCCX (J/g) ¡sn' (J/g) TET(650)2 -34.8 0.08 ^H. rET(r000)2 -60.3 0.39 TET(2000)2 -73.9 0.40 -28.6 32.62 t3.7 28.r9 rET(2900)2 * * + * rl.9 26.5r r09.2 0.32

- denotes peak not present in sample thermogram Y refers to melting peak. t refers to crystallisation peak. * denotes subambient scan not performed. 105

'(a)

I m ol o (b) l- q 3 õ'

(c)

75. 0 100- 0 -t00.0 -75. 0 -50- 0 -25. 0 0.0 25. 0 Teøporoture ('C)

Figure 4.5 Thermograms of TPCL urethane acrylate polymers (a) TPCLl250 (b) TPCL200O and (c) TPCL3000. 106

(a)

m of o J (D l ã

-100.0 -75- 0 -50. 0 -25_ 0 0.0 25. 0 50- 0 75. 0 100. o Temporoture ('C)

Figure 4.6 Thermograms of TET (PTN,ÍO)2 based urethane acrylate polymers (a) rsr(6SO)z (b) TET(1000)2 and (c) TET(2000)2. r07

Reduced crosslink densities for the same PTMO molecular weight in this polymer series resulted in a small decrease in the DSC glass transition temperature for the TET(1000)2 sample from -53.7 to - 60.3oC. For the TET(650)2 sample a Tg, albeit of very small heat capacity change, was detected at -34.8oC, whereas no Tg was observed for the TET650 polymer. The similarity in Tg between TET(2000)2 and TET2900 suggests that the lower limit to glass transition has been reached, with further decreases in crosslink density unlikely to produce a significant Tg decrease. This limiting Tg is some 13oC higher than the literature value for high molecular weight PTMO 9. Similar trends in the soft segment Tg in linear polyurethanes have been attributed to a chain anchoring effect where the Tg of high molecular weight soft segment is increased slightly due to a decrease in free volume associated with the loss of chain end mobility 6.

While the Tg of the TET(2000)2 polymer differs little from that of TET2000, the TET(2000)2 has a far greater ability to crystallise, indicated by cold crystallisation and melting peaks at -28.6 and 13.7oC. This result is significant from a number of aspects.

If the DS C results for the TET homopolymer series were considerêd in isolation, it could be concluded that the length of the PTMO chain dictated crystallisation behaviour. In common with other studies 10 it appears that chain length is not the only factor affecting crystallisation; with crystallisation in shorter chains being enhanced if chain motion is increased sufficiently for nucleation of crystallites to occur. 108

While the molecular weight between acrylate crosslinks has been effectively halved in the TET(PTMO)2 series the weight fraction of hard segment (defined 3 as simply the weight fraction of all components excluding the PTMO soft segment ) has been decreased by only lOVo compared to the TET polymer series. This implies that the TDI group connecting the two PTMO groups to the acrylate crosslinks does not impose any significant restrictions on PTMO chain motions. Lower limits in hard segment content below which the hard segment domain microstructure changes from an isolated to an interconnected morphology have been observed in IR dichroism studies tr,t2 on MDI/BD/PTMO polyurethanes. It could be argued that the low level of hard segment in the TET(2000)2 (l6wt%o) is insufficient to permit hard domain formation This argument could not be sustained, however, for the TET(650)2 and TET(1000)2 samples. For similar hard segment weight fractions, linear TDIÆD/PTMO 1000 polymers l3 showed hard domain formation, as evidenced by an increase in the Tg of the soft segment phase from the value for PTMO1000 diol. A high temperature peak was observed in the initial DSC thermogram of TET(2900)2 polymer (Figure 4.7). This peak was absent in the second DSC scan on this sample. 4.2.1 ld) Hard sesment model oolvmers In addition to the glass transition, crystallisation and melting peaks observed for the soft segment in linear polyurethanes, amorphous and crystalline hard segment phases have been reported

13, 14.

Other workers observed a broad high temperature peak above 125"C in TET2000 homopolymer, which disappeared after quenching

8 109

(a)

o 50 100 150 2C0 i errne.-¿i--,.i^e ('C)

12. 94'C AJ/s

(b)

o2'c

50 c -:0 tOl r50 JO.r [e noe.atu, e ,'C)

Figure 4.7 Thermograms of TET(2900)2 polymer (a) initial scan (b) repeat scan. 110

This was believed to be due to short range ordering associated with urethane linkages in the hard segment domains 15,16. Another study, however, failed to detect any high temperature endotherms for the

same polymer 17.

High temperature scans for the urethane acrylate homopolymers showed an extremely small endotherm near 120qC (Figure 4.8(a)). A repeat scan for the TET2000 polymer showed a decrease in peak enthalpy, but no change in peak temperature. The persistence of this peak after the initial run suggests that the broad peak observed at I25"C 8 ..tuy reflect different endothermic processes in the polymer. In addition to the small peak near 120oC, another peak developed in the DSC thermogram of TET2000 homopolymer after the polymer had been held at 240"C for 10 minutes, then cooled slowly and equilibrated at room temperature for 2 hours(Figure 4.8(b)). While this peak at about 130"C was again small, it was sharper and the + better resolved than the peak at 120oC, and may correspond to broad peak near 125"C seen by othel workers 8. A high temperature peak was also observed in the initial DSC scan of TET(2900)2 polymer (Figure 4.1a) with no peak being found in a subsequent scan on the same sample (Figure 4.7b). These thermograms also showed large crystallisation and melting peaks similar to those for the TET2900 homopolymer, arising from crystallisation and melting of the PTMO

soft segment. In order to clarify the observed behaviour, two hard segment model polymers were scanned under the Same conditions as the

urethane acrylate homopolyme¡s. Considering firstly the 1: I mole ratio (i.e. uncrosslinked) TDIÆIEA polymer (Figure 4.9): four endotherms were found on the first thermal scan, 111

1 17. BB'C

110 120

9 E qJ 1- o :o c uJ (a)

- 150 - 100 -50 0 50 100 150 200 Tempenatune ('c)

131 .32'C 0 . 04363cð ì,/g

100 150

,9 E O 1- o ! c (b) Ll-,

---r -1- - 120 -70 -?o 30 BO 130 1BO Temper-atur-e ('C)

Figure 4.8 Thermogams (b) scan of same sampl crystallisation and melting on the DSC pans. (A dry box was nor used for these DSC runs). l12

(a)

.(J s o, E o D o) c r,!

250 0 50 100 . 150 200 Tempenatune ('Cl

Figure 4.9 Thermogr¿ìms of 1:1 mole ratio TDI,/HEA hard segment model polymer (a) initial scan (b) repeat scan. 113

I, around 40o, II a small peak at L20o, III a broad peak with a maximum at 145" and IV a broad peak near 200"C. Thermogravimetric analysis indicated that rapid polymer decomposition may start when the temperature exceeds 200oC, hence peak IV will not be considered further. The second scan for this sample showed a thermogram..devoid of peaks, save for peak II, which remained unchanged in both intensity and position. DSC thermograms for the crosslinked 1:2 mole ratio TDI/HEA polymer are shown in Figure 4. 10. Peaks I and II are again seen in the first scan, while the second scan resembles that for the linear model polymer in that only peak II remained.

Seymour and Cooper l5 observed multiple endotherms in polyether (PTMO) and polyester (polybutylene adipate) urethanes with a MDIiBD hard segment. Three endotherms were found in an ester soft segment material containing 38 wtVo hard segment : I, around 70o, II a broad peak with a maximum at 160o and III, a small peak at 185o. Annealing samples for 4 hours between 80 and 11OoC caused peak I to move up in tempelature until it merged with peak II to form a single endotherm. Annealing at 175oC caused peaks I and II to disappear, and gave an enhanced III peak, which was accompanied by the

appearance of a distinct ring in the wide angle x-ray patterns absent in samples'annealed at a lower temperature.

The observed x-ray halo was taken to indicate that the III peak was due to microcrystallinity. Hard segment length influenced the development of peak III, with an ester soft segment Sample containing 114

(a)

9 E o) 5 o Ðc tU

(b)

50 100 150 2 250 Ìempenaturc ('C)

Figure 4.10 Thermograms of l:2 mole ratio TDVFIEA hard segment model polymer (a) initial scan (b) repeat scan. r15 two MDI units per hard segment l8 showing no peak and another sample containing three MDI units in an average hard segment showing the III peak.

Since the hard segment in these linear polyurethanes is MDI rather than TDI direct comparisons of endotherm temperatures are not possible. Some comparison of endotherm trends may be made, however. The lower temperature endotherm, peak I, in MDI based polyurethanes has been attributed to dissociation of urethane - soft segment bonds 19,20. Interurethane hydrogen bond dissociation was presumed to account for the peak at 160oC, with Severe annealing producing a single microcrystalline peak (III) in the thermogram 21. Some doubt must be cast on the hydrogen-bond dissociation mechanisms proposed - significant endotherms have been found at high temperatures in piperazine hard Segment polyether polyurethanes 15,22'23 in which no hydrogen bonding capability exists.

Seymour and Cooper l5 proposed that the three peaks are morphological in origin, rather than based on the disruption of hydrogen bonds. Peaks I and II were attributed to disordering of hard segments with relatively short-range order that may be improved by annealing. The lower temperature I peak was attributed to clusters of shorter hard segments, consisting of perhaps only one MDI unit, while the higher remperature II peak was thought to be due to the disordering of longer segments. The lowest temperature peak around 40o in the model compounds considered here may be due to the disruption of short range order in hard segments clusters. The II peak at 120o, although small occurs to the same extent in the model compounds as in the urethane acrylate 116 homopolymers. This peak may be due to short range ordering of some TDI units with adjacent acrylate groups, perhaps involving polar urethane - carbonyl interactions o¡ NH hydrogen bonding with the alkoxy oxygen of the urethane group 24, and appears unchanged on repeat thermal scans.

The high temperature peak between 110 and 130o, which was only seen in initial DSC scans, may be indicative of hard segment crystallinity. The absence of this peak in the 1:2 mole ratio TDI/HEA model polymer may be due to the crosslinked network inhibiting orientation and packing of the TDI/HEA hard segments into microcrystalline domains. The high temperature peak, seen at 131o in the TET2000 polymer after heating at 2400 for 10 minutes, and at 110'C in the TET(2900)2 sample, certainly cannot be associated with any soft segment melting.

Since some depolymerisation had evidently occurred at 240"C in the heated TET2000 sample the hard segments may be able to form crystalline domains more readily when the network crosslink density has been effectively decreased ( as manifested by the similarity of the PTMO crystallisation and melting peaks in the TET2000 polymer after heating to those seen for the TET2900 polymer). The appearance of a high temperature peak in the TET(2900)2 sample supports the hypothesis that very low network crosslink density favours hard s egment crys tallinity. Disappearance of peaks I and III in the hard segment model compounds after initial thermal scanning contrasts with repeated scan results for linear polyurethanes. This is perhaps a reflection of the glassy nature of the networks - reformation of ordered hard segment domains in rubbery networks must take place many times faster than in a polymer glas s w here the approach to equ ilibrium may be t17 infinitesimally small in an analogous manner to the ageing phenomenon in glassy polymers 25. The larger hard segment endotherms observed in the TET2000 urethane acrylate polymer by Speckhard et.al.17 may be due to the considerably lower concentration of photoinitiator used (O.67 wtTo o,f 2,2' -diethoxyacetophenone/N-methyldiethanolamine) compared to the 3wt%o Irgacure 651 used in the samples shown. A slower polymerisation rate, coupled with lower urethane acrylate conversion, may enable more complete hard segment aSsociation to take place, in much the same way as depolymerisation in the heat treated TET2000 polymer produced a decrease in crosslink density, which enabled hard segment crystallites to form.

Glass transitions due to the presence of an amorphous hard phase were not observed, however this cannot be taken as unambiguous evidence for the lack of a hard segment phase. While some workers have observed hard segment glass transitions in linear polyurethanes 6 Leung and Kobe¡stein I determined ACp for the

MDI/BD hard segment to be effectively zero, in agreement with results obtained by Camberlin et. aI.2.

4.2.2 Annealing Experiments A TET 2000 sample was chosen for annealing experiments since this sample showed a well defined Tg process where any increase in transition temperature due to the "mixing in " of hard segments would be apparent. The sample was annealed for 30 minutes at 150oC then rapidly quenched to - 120'C and then scanned through the glass transition. The annealing/quench cycle waS then repeated, with annealing for 30 minutes at 175'C. No change in Tg of the soft 118 segment process, either in terms of temperature or heat capacity change was observed. This cannot, however, be regarded as conclusive evidence for the lack of a hard phase in the polymer. The thermodynamic incompatibility between the hard and soft segment in the TET2000 may be such that microphase separation persists despite attempts to force phase mixing. In this respect the annealing experiment may be regarded as only being capable of providing a. positive result, i. e. the existence of a hard phase could be postulated if the Tg of the TET2000 polymer had increased after annealing/quenching. The lack of any observed shift cannot be used to disprove the existence of a hard phase.

4.3 Summary DSC analysis showed decreases in glass transition temperature with increasing soft segment molecular weight until the soft segment molecular weight approached 2000, after which the polymer Tg did not change significantly. Recrystallisation and melting peaks were observed in PTMO based urethane acrylates w hen the molecular weight between crosslinks exceeded 2000. The presence of a TDI group as a link in the TET(2000)2 sample did not inhibit the development of crystallinity in the PTMO soft link. Minor endotherms corresponding to the development of hard segment crystallinity, seen in hard segment model compounds, were only detected in urethane acrylate polymers with very low crosslink densities. No direct or indirect evidence for the existence of an amorphous hard phase was detected in DSC scans. The soft phase glass transition 119 temperature in one sample remained unchanged after extensive annealing followed by rapid quenching. 120 CHAPTBR FIVB

NMR

5.1 Introduction Liquid phase l3C Ntr¡R techniques ìffere used to verify completion of the synthesis of urethane acrylate prepolymers and assign peak resonances. T1p(C), T5¡ and Tlp(H) measurements were made on solid polymers and TET2900 copolymers with methyl acrylate and TEGDA using the data acquisition techniques and pulse sequences outlined in Chapter Two.

5.2 Results and Discussion 5.2.1 13C liquid phase NMR The high resolution NMR spectrum of TET650 prepolymer is shown in Figure 5.1. The peak assignments given are based on the NMR spectra of the components, namely TDI, HEA and PTMO 650, and the anticipated changes in these resonances when the prepolymer was synthesised. The two major peaks in the spectrum are those due to the inner and outer tetramethylene oxide carbons at 26.4 and 70.5 ppm respectively. The lack of vicinal protons for several of the carbons in the TDI/HEA part of the prepolymer leads to poor Nuclear Overhauser Enhancement of the neighbouring carbon signal, decreasing peak intens ity. 5.2.2 Solid State PEMAS 13C NMR 5.2.2 (a) TET homopolymers The PEMAS l3C NtrrtR specrrum of TET 650 polymer is shown in Figure 5.2 with the peak assignments given based on the previous assignments from the liquid spectrum of TET 650 prepolymer. T2L

8 l5 65 cHl CH.-CH o 9 'l t2 t3 ll rC- ocH2cH2o c _NH il 2 il o NH-c0' CH2CH2CH2O)o-mt_fne 2 3 -(CHz1l 14 14 ll

l0 4

12 13 34567 l2 910 15

60 .¡J

Figurerl,Lfljs! rgsoturion r3c liquid NMR specrrum of rET650 prepolymer (rovo w/v in CDCI3). Peak assignmenrs are as given above. 122

CH, H,C:CH o o il cocH2cH2o-c- ll t zz r NH- CO (CH2CH2CH2CHzO)n - TDI-FIEA o

:00 rao r6n

Figure 5.2 l3C PEMAS NMR sp€ctrum of TET650 polymer 123

The most striking feature of the spectrum is that the only peaks which can be readily distinguished are those for the outer and inner PTMO soft segment carbons, designated as Cl and C2 respectively' The argmatic carbon peaks, weak in intensity in the liquid NMR spectrum' are not discernable above baseline noise. Dickinson et. al. I in a study on the effect of temperature on backbone carbon resonances found that at subambient temperatures cross polarisation waS efficient but, aS the probe temperature increased, the dipole-dipole interactions responsible for magnetisation transfer from the "hOt" prOton poOl to the "cold" carbon nuCleii became progressively weaker as polymer segmental motion increased' The signal to noise ratio in the PEMAS l3C NMR spectrum of poly (2-hydroxy ethyl methacrylate) also increased 2 as water was imbibed into the polymer. This was attributed to plasticisation of the network with a concomitant decrease in polymer T, to below that of the NMR probe temperature.

r 1p(e) In rhe T1p(C) experiment spin-lock contact was established initially, and then terminated by switching off the lH rotating field and allowing the l3C polarisation to decay in its own rotating field for a lH required interval. After the carbon field was switched off, the field l3c was restored and l3c relaxation data acquired. The decay of the magnetization waS assumed to be due to a single relaxation time, which line was evaluated 3 by a 3-parameter fit of the dependence of the intensity on the decaY time. t24

r20 expt

100 calc

80 u) (u 60

GI è0 40 qt

0 0 l0 20 t (ms)

Figure 5.3. DATAFT curve fitting results for Tlp(C) experimental data for the C2 PTMO carbons in TET650 polymer. r25

Figure 5.3 shows the experimental and calculated results obtained for the C2 carbons in the TET 650 polymer. Single exponential fits to the data were used. When attempts were made to fit the data to a double exponential equation, large standard deviations and increased difficuliy in' achieving fits were encountered as the number of parameters increased from two to four. Both spin-spin and spin-lattice processes can contribute to the relaxation mechanisms leading to observed T1p(C) values 4-6. The spin- lattice relaxation is related to the molecular dynamics of the polymer system, while the spin-spin process is caused by dipole field fluctuations generated by changes in the spin state of protons near carbons and is not influenced by molecular motions.

TABLE 5.1 The dependence of PEMAS NMR parameters on the length of the soft link in the TET urethane acrylate series. n is the number of PTMO units in the soft link. C1 and C2 are the outer and inner PTMO ca¡bons defined in Figure 5.2.

Peak n 9 I4 20 28

C1 Tsr /msect 0. 1810.03 0.17610.003 1.22+0.02 2.3+0.5

(70.5ppm) T1p(C)/msecï 2.810.6 6+2 8.610.6 t0+2

T1p(Fl)/msect 3.4+6.6 tt+2 24+4 20+4

C2 T.sr/msect 0.20+0.04 0.5+.1 r.3+0.2 2.3+0.5

(26.7ppm) T1p(C)/msect 4.5!0.4 8+1 14+1 t8+2

Tlp(t-I)/msecf 4.8r0.9 20!7 70+20 70r30

T Errors cited are SE estimated from r.m.s deviation of fit.

The Tl and Tlp(C) relaxation processes for polypropylene and

PMMA at subambient temperatures are dominated by spln-spln lnteractlon s 7'8, other workers 5'9 have concluded, however, that for 126 several amorphous polymers (polystyrene, PMMA, poly(phenylene oxide), polyethylene terephthalate and polycarbonate) T1p(C) at room temperature is largely dominated by spin-lattice processes and can be regarded as an indicator of carbon molecular mobility. PEMAS l3C NtvtR time constants for the PTMO based homologous series are shown in Table 5.1, with T1p(C) values shown in Figure 5.4 as a function of the number of PTMO units in the soft link. Tlp(C) lengthened 3-fold for both Cl and C2 as the molecular weight of the soft link was increased. The absolute difference in T1p(C) between C1 and C2 increased, with values for C1 appearing to level off to a greater extent than those for C2. Allen et. al. l0 observed similar magnitude increases in Tlp(C) for the inner CHZO in tetra (ethylene glycol) dimethacrylate as a function of percent conversion - T1p(C) increased rapidly after the sample had vitrified. V/hile this observed trend of Tlp(C) increasing as the polymer became more glassy may appear to be at variance with the results obtained here, other workers I observed well-defined T1p(C) minima for a series of polymers consisting of ct,tr) dihydroxypoty(propylene oxide) (PPO) chains crosslinked with tris(isocyanatophenyl phosphate) in the temperature range - 10OoC to +100oC. The T1p(C) minima shifted to lower temperatures for longer PPO chains, i.e. as the Tg of the polymer decreased. Presumably a similar situation applies in the case of the PTMO carbons in the urethane acrylate series considered here, with T1p(C) minima shifting to lower temperatures as the polymer becomes more rubbery. For polymers with their Tg near the l3C NtvtR probe temperature ( in this case the TET650 polymer) it could be expected that T1p(C) values for the units responsible for the glass transition would be at or near their minima since at this point the mid-kHz 13C roraring field frequency coincides with their group motion frequency. 127

ch r

CJvl0 o_ F I

Ã I

0 0lon2o 30

Figure 5.4. T1p(C) relaxation time of l3C resonances in a 60-kHz rotating field for Cl(o¡ and C2(Â) PTMO ca¡bons as a function of the number of PTMO units, n, in the soft link of the urethane acrylate. Error bars are shown where the standard error, estimated from the rms deviation of the fit, overlaps the bounds of the data symbol.

u) 2

(t)I F

^ I a 0 0 l0 20 30 n

Figure 5.5. Spin-locked, cross polarisation time constants, T5¡, for C1(o) and C2(A) PTMO ca¡bons, designated as for Figure 5.4. Error bars a¡e shown where the standard error, estimated from the rms deviation of the ht, overlaps the bounds of the data symbol. 128

Tcr and T1o(H) Cross-polarisation occurs through a static-dipolar coupling of proron and 13C spins 4 and is affected by the number of directly bonded protons and isotropic carbon motions which may reduce the cross- polarisation efficiency by disturbing dipolar interactions. It is believed that any increases in near-static components in the motion of a particular group are manifest in lengthened values for the T5¡ time constants of the carbons in that group.

The T5¡ time constants are shown in Figure 5.5 as a function of the number of tetramethylene oxide units, n, in the soft link of the urethane acrylate network. After the TET650 polymer (n=9) the Tg of the polymer dropped below the probe temperature (25+2"C) - the increase in network mobility was accompanied by a slight lengthening in T5¡ for the TET1000 polymer. The T5¡ values then increased by an order of magnitude for both Cl and C2 in the TET2000 and 2900 polymers (n=20 and 28 respectively). Increases of a similar magnitude in T5¡ were observed ll in a series of oligo-ethylene glycol dimethacrylates (P(EG)*DMA) as the number of soft links was increased from x=1 to x=13, thereafter (x=22) the T5¡ of the inner CHZO groups sho¡tened by an order of magnitude with the onset of crystallinity in the ethylene glycol soft segment. For the PTMO based urethane acrylates considered here, further increases in T5¡ for both C1 and C2 as the number of soft links increases to 28 suggests that these groups do not become involved in crystalline domains, or alternatively, that any crystallisation is negligible. T1p(H)s for each carbon are generally averaged by spin diffusion, giving no information about group motions 10,12 but are governed by a single proton "pool" dominating each sample. It is apparent, however, r29 that the Tlp(H) values for C2 are between 50 and300Vo greater than those for C1. Similar ratios were observed l3 for a linear polyurethane (Lycra 146) consisting of PTMO 2000 with MDI hard segments. In unstretched Lycra samples T1p(H) for the inner PTMO carbon was 27.2ms, more than double that for the outer PTMO carbon (12. lms). In a stretched sample (extension ratio l"=5) T1p(H) values decreased to 6.2 and 6.9ms for the inner and outer carbons respectively. The two major contributions to T1p are motional and direct dipolar 13, with the latter becoming more pronounced on crystallisation. In the more crystalline (i. e. stretched) Lycra sample proton spin diffusion can occur more efficiently l4'15 yielding an average value for Tlp(H). Although no explanation for the larger Tlp(H) for the inner PTMO carbon (C2) in the unstretched sample was offered, the differences in Tlp(H) for Cl and C2 for the urethane acrylates considered here again suggest that PTMO soft segment crystallisation is negligible. The changes in Ts¡- and Tlp(C) for both Cl and C2 in this homologous series indicate that for both the Iow frequency and 60-kHz components of motion the mobility of the PTMO link is important. This contrasts with other results I I for a series of oligo ethylene glycol dimethacrylates where 60-kHz motions in the oxyethylene soft links did not change significantly throughout the series and were not relevant to changes in bulk properties (i. e. compression modulus). The damping of methacrylate group motions was found to contribute to the observed changes in gross macroscopic behaviour. The influence of the methacrylate group is perhaps not surprising when the mechanical properties of tetraethylene glycol networks with methacrylate and acrylate end groups are compared. The glass transition temperature of poly(TEGDMA) is 115oC ll while that for poly (TEGDA) is 46"C 130

(Chapter Three). Clearly the methyl group has a considerable influence on mechanical properties.

<',') ¡/hYI.F-|-?Ofìfì/meth.rl qnrr¡ tc

A series of NMR spectra of TET2900 50 wtTo methyl acrylate copolymer are shown in Figure 5.6. Reasonable separation in l3C chemical shifts between the PTMO peaks and those due to methyl acrylate enabled PEMAS time constants to be determined through the copolymer series

containing 10, 25 and 50 wt%o methyl acrylate.

r1p(Ç) The Tlp(C) relaxation times for the inner and outer PTMO carbons of TET2900 copolymerised with 10, 25 and 50 wÍ.Vo methyl acrylate are shown in Table 5.2 and plotted in Figure 5.7.

TABLE 5.2 The dependence of PEMAS NMR parameters onwt%o methyl acrylate in the TET290O I methyfacrylate copolymer series. C1 and C2 a¡e as defined in Figure 5.2. ; Vo mefhyl Peak acrylate 0 10 25 50

C1 Ts/msect 2.3+0.5 |.7!0.2 1.4+9.2 1.7+0.3

(70.5ppm) T1p(C)/msect t0+2 1 1+ 1 I 2+ 1 1 1+11

T1p(H)/msect 20!4 25!3 25+3 28+6

C2 Tsr/msecf 2.3+0.5 2.t+0.3 1.9+0.4 1.8+0.5

(26.7ppm) T1p(C)/msect 18+2 t7+2 t3!2 t2!l

Tlp(H)/msect 70+30 60110 60r30 60+30

t Errors cited are SE estimated from r.m.s deviation of ht. 131

t.o (ms)

18.00

15.00

12.00

10.00

7.00

5.00 a

2.00

1.00

0.50

0.20

Figure r3C PEMAS NMR methyl acrylate rent contact times t n in the .oúting acrylate) peaks are ner and outer carbon peaks. t32

Gls

O o- 10 I F

0 01020304050 60 WtTo methyl acrylate

Figure 5.7. T1p(C) relaxation time of l3C resonances in a 60-kHz rotating field for Cl(o¡ and C2(A) PTMO carbons as a function of weight per cent methyl acrylate in TET2900 co methyl acrylate networks.. Error bars a¡e shown where the standard error, estimated from the rms deviation of the fit, overlaps the bounds of the data symbol.

2.5

2.0

(t) r.5

(t)'ì F 1.0

0.5

0.0 0 l0 20 30 40 50 60 Wt% methyl acrylate

Figure 5.8. Spin-locked, cross polarisation time constants, T5¡, for Cl(o) and C2(A) PTMO carbons as a function of weight per cent methyl acrylate in TET2900/meth.yl acrylate nerworks, designated as foi Fþure 5.7. Error bars are shown where the stañdard error, estimated=from the rms deviation of the fit, overlaps the bounds of the data symbol. 133

Cl showed little change in Tlp(C) on copolymerisation, while C2 showed a systematic decrease in Tlp(C) to 5O Vo of its initial value in the

TET2900 co 50Vo methyl acrylate copolymer. C2 may be a more sensitive indicator of changes in molecular motion than CI - CZ also showed the most pronounced change in Tlp(C) through the TET650- 2900 homopolymer series.

Tlp(H) and Tg¡Time constants. T1p(H) and T5¡values for both C1 and C2 were relatively constant through the methyl acrylate copolymer series suggesting that little damping of low frequency motions occurs on copolymerisation.

\ ) ) In\TFT?QOO/TFGDA conolvmers TlB(C) The T1p(C) relaxation times for the inner and outer PTMO carbons of TET2900 copolymerised with 25 and 50 wt%o tetraethylene glycol diacrylate are shown in Table 5.3 and Figure 5.9. The peak at ?0.5ppm is a composite of the ethylene glycol carbons of TEGDA and C2 in the TET2900, and therefore cannot be taken as an indicator of changes in the molecular mobility çf either component- Simon 16 determined a T1p(C) value for the oxyethylene carbons in TEGDA homopolymer at 298K of 3.5ms : the relatively small decrease in the combined T1p(C) time constant with increasing wt%o TEGDA in the copolymer is interesting, since the molecular motions of the TEGDA component must be ¡eflected to some extent in the dynamics of the 70.5ppm peak. Belfiore l7 et.al. found that mid-kHz components of group mobility for polycarbonate carbons were released by the addition of a plasticiser to the polymer. r34

(t) l5

O o- 10 F I

0 0 l0 20 30 40 50 60 WtTo TEGDA

Figure 5.9. Tlp(C) relaxation rime of l3C resonances in a 60-kHz rotating field for Cllo¡ and C2(Â) PTMO carbons as a function of weight per cent TEGDA in the TET2900 co TEGDA copolymer series. Error ba¡s are shown where the standard error, estimated from the rms dèviãtion of the fit, overlaps the bounds of the data symbol.

(n 2

(t)J F

a 0 0102030405060 WtTo TEGDA Figure 5.10 S n time constants, T5¡, for Cf and C2 PTMO caibons in the er series designated as for Figure 5.8. Error ba¡s are show estimated from the rms deviátion of the fit, overlaps the bounds of the data symbol. 135

The relatively small decrease in the T1p(C) time constant for the composite 70.5ppm peak possibly reflects similar plasticisation of the TEGDA polymer by the rubbery TET2900 matrix, releasing mid-kHz components of motion in the TEGDA ethylene glycol link.

TABLE 5.3 The dependence of PEMAS NMR pafameters on_ wt 7o TEGDA for the TET2pQQ|IEGDA copolymer series. Cl and C2 are as defined in Figure 5.2. The peak at 70.5ppm results from overlapping Cl and TEGDA CH2O peaks.

WtVo Peak TEGDA 0 25 50

CH2O Ts/msecl 2.3!0.5 t.4!0.3 0.08r0.02

(70.5ppm) T1p(C)/msecT 9.8+1.6 1 0J 1 8.5+0.9

T1p(Fl)/msect 20x4 30r10 90130

C2 Tsr/msect 2.3+0.5 2.3+0.5 1.6+0.5

(26.7ppm) T1p(C)/msecT r8+2 tg+2 I4+2

T1p(H)/msect 70+30 150150

f Errors cited are SE estimated from r.m.s deviation of fit.

The C2 carbon Tlp(C) values showed only a slight decrease as the

wtVo TEGDA increased, indicating that the mid-kHz components of group motion in the TET2900 soft segment were not greatly dampened on

copo lymeris atio n.

T5¡ and T1p(H) T5¡ for the combined oxymethylene/oxyethylene peak at 70.5ppm

decreased ro a rhird of the initial (T8T2900 only) value in the TET2900 co

507oTEGDA copolymer (Figure 5.10). Although a T5¡ value for the inner ethylene oxide carbons in TEGDA was not found l6 the magnitude of the T5¡ for the 50wt7o copolymer is close to that found for the TEGDA end carbon CH20 (.O29+.001ms). The copolymer result suggests that any 136 plasticisation of TEGDA group motions by the TET2900 has little effect on the low frequency, near static motions in the TEGDA ethylene glycol groups. The slight decrease in T5¡ for C2 at 26.7ppm indicates that any damping of low frequency motions in the TET2900 PTMO. unit, on copolymerisation is minimal.

tÂra f^- TF'Ftf¡nn ; 'nÞTî onrì < ) ') /Ä\Vqri n l\ll\rfÞ Dqrqmc ^ n homopolymers with time.

TABLE 5.4 The dependence of PEMAS NMR parameters on storage time for TET2000 and 2900 immediately after polymerisation anã after storage at room temperature for the indicated interval. Cl and C2 are as deñned in Figure 5.2.

Peak TET2OOO after 110 TET29OO after 100

days days

C1 Tsr/msecï t.22+0.02 0.7+0.1 2.3!0.5 t.5+0.2

(70.5ppm) Tlp(C)/msect 8.610.6 7.3+0.8 rol.2 8.7+0.5

TloG{)/msect 24+4 29+6 20!4 20!3

C2 Tsr/msecT r.3.0.2 1.110.2 23!0.5 2.O+0.3

(26.7ppm) Tlp(C)/msecï 1411 10.5+0.9 t8!2 10.3+0.3

Tlp(FI)/msecT 10!20 100+50 70r30 40+10

t Errors cited are SE estimated from r.m.s deviation of fit. T1p(C) The Tlp(C) time constants for TET2000 and 2900 polymers for samples run within 3 days of casting and again after Storage at room temperature in a dessicator for between 100 and 110 days are shown in Table 5.4 and Figure 5. il. From Figure 5.11 it can be seen that while any change in Tlp(C) for C1 is small, C2 for both polymers shows larger decreases in Tlp(C). That for TET2000 decreases nearly 30Vo, while t37

20 E TET2OOOC2 . TET2OOO C1 O TET29OOC2 l5 . TET29OO Cl rt) r

r0 E I U T o- F I 5

0 0 20 40 60 80 100 120 Days after polymerisation

Figure 5.11 Tlp(C) relaxation time of 13C resonances in a 60-kHz rotating field for Cl and C2 PTMO ca¡bons in TET2000 and 2900 as a function of time after casting. Error bars are shown where the standard error, estimated from the rms deviation of the fit, overlaps the bounds of the data symbol.

(nJ F I tr TET2OOO C2 X TET2OOO Cl I O TET29OOC2 . TET2gOO Cl 0 o 20 {0 @ .80 100 r20 Days atter polymensatlon Figure 5.12 Spin-locked, cross polarisation time constants, T5¡, for C1 and C2 PTMO caibons in tEî2000 and 2900 a-s a function of time after casting. Error bars are shown where the standa¡d error, estimated from the rms deviation of the fit, overlaps the bounds of the data symbol. 138

T1p(C) for the inner PTMO carbon in TET2900 decreases to close to 607o of the initial value after 100 days. DSC experiments (Chapter Four) showed greater PTMO soft link mobility in the TET2900 polymer compared to the TET2000 polymer. The greater damping of mid-kHz components of motion in the TET2900 sample with time .uy i" due to more extensive crystallinity developing in the PTMO soft link.

Tsr and TlolH) T5¡ values decreased markedly for Cl in both TET2000 and 2900 (Figure. 5.12), while those for CZ decreased only slightly for both polymers. In addition to the single exponential fit a double exponential fitting program (SIMTWO) was used for the T5¡ data. This was only successful in fitting the magnetisation decay data for the TET2900 sample stored for 100 days. Table 5.5 lists the T5¡ and Tlp(H) values given by the double exponential curve fit for C1 and C2 carbons for the stored TET2900 polymer. The T5¡ time constants were resolved into long and short components - with the standard error obtained being of the order of the shorter time constant. Dickinson et.al. 13 were also able to assign two components to the decay curve of a linear polymer (Lycra 146) containing PTMO 2000 soft segments. The longer values obtained for both inner and outer PTMO carbons (2.6-3.5ms) \'r'ere close to those found for the PTMO carbons in the TET2900 polymer and are typical of T5¡ times found for mobile CH2 groups in polymers above the glass transition 6. Five-fold extension of Lycra 146 samples markedly increased the proportion of species with the shorter Tsl values, with reasonable agreement being found in the fractions of inner and outer PTMO carbons with shorter T5¡s (71 and llVo respectively). The proportions of C1 and C2 carbons with a shorter T5¡ 139 component in the stored TET2900 sample conside¡ed here are also in close agreement at 12 and 15Vo respectively.

TABLE 5.5 Relaxation parameters for TET2900 polymer 100 days after polymerisation obtained from double exponential datafit. Cl a¡rd C2 are as defined in Figure 5.2.

Peak

C1 Tsl.t/msect 2.6fl.6(85Vo)

t1 (70.5ppm) I sl2/msec I 0.0810.06(157o)

T1p(Il)/msect 16+3

C2 Tsl.l/msect 3.srO.8(88%)

(26.7ppm) Tsr.r/msect 0.09t0.08(127o)

T1p(F{)/msect 24+6

I Errors cited are SE estimated from r.m.s deviation of fit.

Cory et. al. 18 investigated the double exponential character of the T5¡ relaxation in poly (oxymethylene) (POM) and obtained two T5¡ values of 13 and 290ps. Chemical shift differences enabled the shorter T5¡ value to be conclusively assigned to crystallised POM units- These observations on the effect of crystallinity on decay populations, and DSC data (Chapter Four) indicating that the PTMO units in the TET2900 homopolymer are capable of crystallising near room temperature may account for the observed biexponential decay for the stored TET2900 polymer.

Other workers l9 analysed inversion recovery cross polarisation (IRCP) decay of the outer PTMO carbon (71ppm) in a polyurethane elastomer and were able ro best fit the data to a biexponential two- 140 component model, indicative of a two-phase polymer morphology. Since the molecular weight of the PTMO units in the polymer was not specified, the two-phase nature of the decay cannot be attributed to increased constraint of the PTMO groups due to the development of crystallinity, or to association with hard domains. Biphasic decay has been observed I for networks of crosslinked poly (propylene oxide) (PPO) soft segments with PPO molecular weights ranging from 400 to 2000 The low molecular weight of the PPO units, coupled with their decreased crystallisation tendency compared to PTMO units, suggested that two motional populations resulted from close coupling of a proportion of the PPO units with the phenyl group in the tris (4-isocyanato phenyl ) crosslinks. The extent of microphase separation of MDI polymerised with PTMO and a variety of chain extenders has been studied 20, with some degree of phase mixing being manifested in biexponential magnetisation decay.

Biexponential curve fits were unsuccessful for all the PTMO relaxation decays discussed in this chapter, save for the TET2900 polymer sample stored for 100 days. This suggests that distribution of the PTMO soft segments between soft and hard phases in the TET urethane acrylate polymers considered here does not occur, or is insufficient to be reflected in clearly biphasic decay behaviour. 5.3 Summary T1p(C) times measured for the TET urethane acrylate homologous series showed significant differences in carbon mobility as the soft segment MW was increased from 650 to 2900, with T5¡ times showing a similar trend. Copolymerisation of TET2900 with either methyl acrylate or TEGDA gave relatively little change to either Tlp(C) or T5¡ indicating t4t that both low and mid-kHz components of motion are not greatly damped in the TET2900 PTMO units in the copolymers.

Only one example of biexponential decay behaviour was observed - this was attributable to long term crystallisation in the soft segment of TET2900. Clear evidence for the existence of PTMO units in two or more motional environments was not observed and the presence of more than one phase was not able to be confirmed. 142 CHAPTER SIX

SORPTION AND DIFFUSION

6.1 Introduction The sorption of water in TET, TPCL and TET(PTMO)2 based urethane acrylates has been studied at three relative humidities (33, 79, and 98Vo) and in water, all at 25"C. After equilibrium had been attained (i.e. sample weights remained constant after 3 days) samples were removed from their sorbing environments and placed in a dry atmosphere where desorption was monitored. The sorption behaviour of TET based urethane acrylates with methyl acrylate, hexanediol diacrylate (HDDA) and tetraethylene glycol diacrylate (TEGDA) was also examined. These comonomers were chosen in order to investigate the effect of increasing and decreasing network crosslink density on sorption behaviour. TEGDA and HDDA were included to enable the effects on sorption behaviour of relatively hydrophilic (TEGDA) and hydrophobic crosslinkers to be compared.

6.2 Diffusion kinetics Diffusion of a liquid or gas in a polymer matrix can be described by Fick's laws. The first law states that, in one dimension, the rate of mass transfer gf a penetrant through a unit area of polymer is proportional to the concentration gradient (ôc/õz) normal to the surface l: ôc F D õz

where F is the diffusion flux (mass per cross-sectional area per unit time) and D is the diffusion coefficient. t43

The solution to Fick's laws for sorption for samples of thin sheet geometry 2 is given by Equation 6. 1, in which M¡ and M- are the masses of solvent at times t and infinity respectively, 2l is the sheet thickness, and D is the diffusion coefficient.

8 (2n+l)znzD t - I exp [ I (6. 1) n 2 4î--

For desorption M- is the total weight loss and M¡ is the weight loss at time t.

Data from sorption and desorption runs is normally presented as a plot of fractional mass uptake of penetra", {- agains ,#where t is the run time in seconds and 2l is the thickness of the polymer sheet. These plots are referred to as reduced sorption/desorption curves, since the sheet thickness is included in the abscissa

Systems which conform to the Fickian interpretation of sorption are those for which the initial experimental data (\F(Mt, M-) < 0.5) for a plot ^ Mt trlZ of -'o: versus 7i fit a straight line according to Stefan's approximation 3 : Mt _ru=2( Dt t2 (6.2) fi p )r

For longer sorption times, Bueche's approximation may also be used 4 :

Mt 8 n2D t ln (1 ln ( (6.3) 1fr-) = fiz 4t2

These equations were used to find the initial diffusion coefficient. Sorption and desorption data were also fitted to the first eleven terms of

Equation 6. I (i.e. n I l0) using the DATAFT program 5. 144

Diffusion is regarded as being Fickian 6 if the following criteria are obeyed for the plots of fractional weight increase (or decrease) against ,ll2 i¡ (reduced time) :

1. Sorption and desorption curves of fractional mass change are Mt linear up to _ft- = 0.6 or greater.

2. Above the linear region both sorption and desorption curves are concave to the x axis. 3. The reduced sorption curve lies above the desorption curve.

If the polymer/penetrant system does not follow Fickian (also known as Case I) diffusion kinetics, the values obtained for D are not the true value of the diffusion coefficient but are regarded as "apparent" diffusivities 7. The overall sorption behaviour depends on the relative rates of relaxation motions of the polymeric matrix and concurrent solvent diffusion 8; in Fickian diffusion the rate of diffusion is much less than that of relaxation due to mechanical and structural effects resulting from polymer - solvent interactions. Another possibility for solvent diffusion in polymers is Case II diffusion where diffusion is rapid compared to other relaxation processes. Case II diffusion is commonly characterised by weight gain which is linear with sorption time. When diffusion and relaxation rates are comparable, behaviour intermediate between Fickian and Case II diffusion is observed - in this case the diffusion is described as non-Fickian or anomalous 9.

Frisch l0 proposed that the three modes of diffusion can be 145 differentiated by fitting the fractional solvent uptake data to the following equatron : M¡ M- =ktn (6.4) where k is a constant indicative of polymer - solvent interaction, and the slope of the plot of log # versus log t, n, is an. exponent characteristic of the type of diffusion. If n = j tn"n Equation 6.4 is equivalent to Equation 6.2 and the diffusion mechanism is Fickian, whereas for Case II diffusion n= 1. Log-log plots "t fr- against t yielding values of n between i and 1 characterise non-Fickian or anomalous diffusion.

6.3 Results 6.3. 1 Homologous series

6.3.1 la)Sorption at 337o relative humidity The EWC, sorption and desorption coefficients Ds and D¿ (calculated from Equation 6. 1) and ns and n6 (calculated from equation 6.4) for a series of TET, TPCL and TET(PTMO)2 homopolymers sorbed a,t33%or.h. and desorbed from 33Vor.h. are shown in Table 6. 1.

TABLE 6.1 Sorption results for TET, TPCL and TET(PTMO)2 polymers equilibrateÀat33%o relative humidity. The standa¡d error for D5 and D¿ given by the DATAFT program is shown in brackets. Polymer EWC Ds (x108) D¿ (x10E) flg nd (%) cm2sec-l cm2sec-l TET65O 1.0 2.43(0.04) 2 0.54 0.ó9 'lB'l'100u u.9 5.0r I 6.9r 0.48 0.61 TET2OOO 0.1 12.Ot 1 18.0r 0.38 0.71 TET29OO 0.ó ló.0r 26.0t 0.43 0.57 .TPCLI25O o.l ó.0(0.2) 7.0( 0.4 0.50 0.7 L TPCL2OOO 0.5 I 1 0.8) 1l 0.51 0.62 TPCL3OOO 0.5 r5.0(r.0 0.54 0.77

TET(650)2 1.1 4.8(0.3 5.7 0.44 0.63 TET(1000)2 0.7 13.ó(0.7) 0.58 0.64 'r'E1(2000)2 0.1 2ó.0(3.0) 0.54 0.50 r46

The reduced sorption and desorption curves for these polymers are shown in Figures 6.1a-f. The correlation coefficients from the log-log plots used to calculate n, and n¿ varied from 0.919 to 0.995 for the homopolymers. The observed n5 values varied widely, from close to the value of 0,5 which characterises Fickian diffusion - down to 0.38 and up to 0.58. The lowest values obtained for ns were for TET2000, TET2900 and TET(650)2 homopolymers all of which showed a departure from linearity Mt between -# = 0.7 5 and 0.85.

Sorption coefficients increased by almost an order of magnitude as the soft link molecular weight inc¡eased from 650 to 2900 in the TET polymer series, while the value of D, doubled as the PCL MW increased from 1250 to 3000. Large inc¡eases in Ds also occurred in the TET(PTMO)2 series with the diffusion coefficient Ds=4.8 x 198 çt¡2t""-l for the TET(650)2 polymer increasing to 26.0x108 cm2sec-1 for the TET(2000)2 polymer. Desorption coefficients D¿ generally exceeded the D5 values for the polymers - this was most marked for the TET2000 and 2900 samples. The sorption and desorption curves for the TET2000 polymer are shown together in Figure 6.2. The value of n¿ exceeded that for n5 in all cases except that for the TET(2000)2 polymer. The EV/C for the homopolymers generally decreased to a constant value as.the MW of the soft segment was increased. This was observed for both PTMO and PCL soft segments, with PCL samples of equivalent MW having lower EWCs than the polymers containing PTMO units. r41

t.2

a 1.0 q o o a o EI 0.8 o99 8 ao = o À a g 0.6 o o ¿ ao + 650 o + + + 1000 0.4 a,O + o+* E a tr o 2000 o + ElEt a + 0.2 o + 4 . + E 2900 + otr o 0.0 0 1000 2000 3000 lt2 2l

Figure 6.1a Reduced sorption curves for TET urethane acrylates containing PTMO soft segments of molecula¡ weights given above. Samples equilibrated at 33Vo r.h. -

t.2

1.0 ffis tr"'@ 1T oú 0.8 o ft" 8 o " ¿l3 +. o+ +E ¿ 0.6 O+ f tr 0.4 o+E o 1250 .¡- E + 2000 E 0.2 4 o 3000 EE

0.0 0 1000 2000 3000 4000 Ív2 2l

Figure 6.1b Reduced sorption curves for TPCL urethane acrylates containing PCL soft segments of molecular weights given above. Samples equilibrated at33Vo r-h. - 148

1.2

1.0 o d': ö o ":þ 0.8 + 8 o ¿lr o + Fa 0.6 o + à +

+ EI 0.4 Otr o (6s0)2 o* + (1000)2 0.2 +tr "" EI o (2000)2

0.0 0 1000 2000 3000 ¡112 2t Figure 6.1c Reduced sorption curves for TET (PTMO)2 urethane acrylates containing PTMO soft segments of molecular weights given above. Samples equilibrated at33Vo r.h- -

1.2

1.0 roo CE . t.o g i.€ o o a + 0.8 + 8 o + trtr F a + ¿ + 13 0.6 a ¿ o + tr 650 o + 1000 0.4 +E tr o1 o 2000 + tr 0.2 OO . + E zg00 +tro 0.0 0 1000 2000 3000 fl12 2l

Figure 6.1 d Reduced desorption curves for TET urethane acrylates containing PTMO soft segments of molecular weights given above. Samples initially equilibrated at 337o r. h. 149

t.2

Eto 1.0 tt II o t E t trE 3 t! 0.8 O 8 !g l

Figure 6.1e Reduced desorption curves for TPCL urethane acrylates containing PCL soft segments of molecular weighs given above. Samples initially equilibrated at 337o r.h. .

h- r.2

1.0 ü+o E+ o "oloo o' o ++ + EÞtr o 0.8 + 8 le,( ¿ E + otr àF 0.6 o o + 0.4 + tr (650)2

o + (1000)2 o + 0.2 o (2000)2 + tr E 0.0 0 1000 2000 3000 lt2 2l

Figure 6.1 f Reduced desorption curves for TET(PTMO)2 urethane acrylates containing PTMO soft segments of molecular weights given above. Samples initially equilibrated at33Vo r.h. . 150

1.2

1.0 ooo aoa o oo o o 0.8 r"d4 8 EE ir< oE à E l3 0.6 tr ¿ û I 0.4 E E Er sorption tr 0.2 Eþ o desorption O

0.0 0 1000 2000 3000 4000 rE lt2 2l

Figure 6.2 Reduced sorption and desorption curves for TET2000 homopolymer equilibrated at 33Vo r.h.. 151

6.3. 1 (.b)Sorption at 797o relative humidity

Table 6.2 show s the sorption results for the same samples considered in the previous section, with sorption and desorption taking place to and from 797o humidity. The reduced sorption/desorption curves are shown in Figures 6.3a-f.

TABLE 6.2 Sorption results for TET, TPCL and TET(PTMO)2 polymers equilibrated at 797o rela¡ve humidity. The standard error for Ds and D¿ given by the DATAFT proga¡n is shown in brackets. Polymer EWu^ D5 (xl0E) D¿ (x10E) n5 nd (vo) cm2sec-l cm2sec-l 'rb1'650 2.O 2.671 0.52 0.60 TET1OOO 1.9 4.68 6.8(0.2 rJ.5Z 0.61 TET2OOO 1.6 0.4 0.49 0.63 'rE12900 r.2 I 0.45 0.59

TPCLl25O r.z 6.2t 6.9(]J.2 0.5I 0.5 t .IPCLzOOO r.2 0.59 0.6r 1?CL3000 1.0 0.51 0.51

TET(650)2 2.1 4.9t 5.8(U.2 0.5ó o.62 'rBr(r000)2 t.7 I 0.56 0.ó3 TET(2000)2 1.3 0.52 u.ór { Dr values show the same trend with increasing MW within a polymer series seen for the 337or.h. samples. A slight improvement in the curve fits was evidenced by lower standard errors in Ds for most samples. EWC also showed the same trends within a polymer series as noted previously i.e. EWC decreased as soft segment MW increased. EWC values fo¡ samples equilibrated at 79Vor.h. were roughly double that for

samples equilibrated at 33Vor.h..

Desorption coefficients obtained from the DATAFT program again exceeded sorption coefficients for all samples, except the TPCL3000 homopolymer.

Values of n, in general showed less variation than those found for

the 33Vor-h. samples - with some samples having n, values close to that 152

r.2

1.0 .f a 0.8 8 ¡O ll o ¿ a" o F 0.6 to@ Er E À a 650 oo o + rd + + 1000 0.4 aO +f tr o !o .+ Ed 2000 o 0.2 î*ll' 2900 EF

0.0 0 1000 2000 3Cno fi12 2l

Figure 6.3a Reduced sorption curves for TET urethane acrylates containing PTMO soft segments of molecula¡ weights given above. Samples equilibrated at797o r.h. .

t.2

1.0 ooo úr \*"FJ¿ +++ ?-È "t&F?fl. tr tr Et- 0.8 EI a, E 8 a ¿13 o OO *** lF{ 0.6 + À + tr# *"t 0.4 a E¡ 1250 a + + 2000 ottr 0.2 . 3000 ;ü

0.0 0 1000 2000 3000

¡112 2l

Figure 6.3b Reduced sorption curves for TPCL urethane acrylates containing PCL soft segments of molecula¡ weights given above. Samples equilibrated at'/9Vo r.h. . 153

1.2

1.0 oo I o .t'* 9otf" "*ffi,"uþf:: ooF 0.8 oo- s 8 o ggtr F f4 À + o 0.6 + i¿ + O+ Ef o E 0.4 q (6s0)2 o +tr + + (1000)2 0.2 O+ .1.tr.'tr o (2000)2 E 0.0 0 1000 2000 3000 út2 2l

Figure 6.3c Reduced sorption curves for TET(PTMO)2 urethane acrylates containing ITMO soft segments of molecular weights given above.Samples equilibrated at 79Vo r.h

t.2

1.0 a, il¡Þ ctÉ .¡s .t? Ed aO 0.8 8 o + FT a + À + + àH 0.6 ¡o o + a tr + tr s 650 o 0.4 a + + 1000 o+E +E o 2000 0.2 3 *.t . 2go0 +s E 0.0 0 r000 2000 3000 ¡t12 2T Figure 6.3d Reduced desorption curves for TET urethane acrylates containing PTMO soft segments of molecula¡ weights given above. Samples initially equilibrated at79Vo r.h. 154

t.2

1.0 o dlf o+ þta Ëi. Ela .å+ 0.8 a +Gr 8 + a tr Å + 0.6 ¿F a tr o +o E 0.4 tr+ t250 o + 2000 ß o 0.2 3000 I + 0.0 0 1000 2000 3000 ùt2 2l

Figure 6.3e Reduced desorption curyes for TPCL urethane acrylates containing PCL soft segments of molecular weights given above. Samples equilibrated at79Vo r.h. .

r.2

1.0 o .Ttl o+'o o €t qf + V"" + 0.8 o+ 8 a + o ¿= a+ F 0.6 à o+tr +tr 0.4 tr (6s0)2 ¡* tr +E + (1000)2 O 0.2 tr . (2000)2 + tr 0.0 0 1000 2000 3000

¡L12 2l Figure 6.3f Reduced desorption curves for TET(PTMO)2 urethane acrylates containing PTMO soft segmenis of molecula¡ weights given above. Samples initially equilibrated at797o r.h. . 155

expected for Fickian diffusion. Values of n¿ exceeded those of n, and were usually around 0.6, indicating that desorption at this relative humidity was non-Fickian. Correlation coefficients obtained from the log-log plots .t # versus t varied from 0.943 to 0.997 in a non- systematrc manner

6.3. I (c)Sorption at 987o relative humidity

Sorption results obtained for the homopolymers at 98Vor.h. are summarised in Table 6.3.

TABLE 6.3 Sorption results for TET, TPCL and TET(PTMO)2 polymers equilibrated at 98Vo relatle humidity. The standa¡d error for D5 and D¿ given by the DATAFT program is shown in brackets. Polymer EWC Ds (x108) D¿ (x108) Il5 nd (vo) cm2sec-l cm2sec-l 'IEt-650 2.4 2.531 0.6u u.59 TETIOOO 2.2 5.3r o.49 rJ.49

TET2UX) 1.8 12.3t 0.6 14.0t 0.ó I 0.48 0.60 'lB'12900 I.5 15.6 0.8 0.47 0.46 .TPCLI25O t.4 7.6 0.4u 0.66 TPCLzOOO 1.2 lI.2t 0.9 10.0r 06 0.55 0.77 'rPCL3000 t.l r8.2(0.7) I3.8r 0.5 0.54 0.57

lET(ó50)2 2.L 4.8( I 0.48 0.ór rB'r'0uuj)2 t.l IZ.lt 0.5 0.56 0.66 TET(2000)2 1.4 1 21.7t 0.9 0.47 o.49

S orption and desorption curves are shown in Figures 6.4a-f . Sorption coefficients D5 show identical trends throughout polymer series to that observed previously i.e. increases in soft segment MW for all homopolymers were accompanied by significant increases in the sorption coefficients.

Desorption coefficients D¿ also showed increases in the same manner as Ds through a polymer series. The differences between Ds and 156

t.2

1.0 .{ oc' t€.-å: tr*.ff'æ o ***þ rEd 0.8 a + 8 O @ {+ OE a o ¿ o EE ra 0.6 ao à ++ d"ttdF o + s +* 650 0.4 OO + + 1000 oo df o 2000 o+ o o.z E++- . 2900 #fl 0.0 0 1000 2000 3000 4000 út2 2l

Figure 6.4a Reduced sorption curves for TET urethane acrylates containing PTMO soft segments of molecula¡ weights given above. Samples equilibrated at98Vo r.h. .

t.2

1.0 o t + % qÞ'F s + o 0.8 ofa_ E""Ë 8 + à o*s 0.6 tr ¿ o+tr tr od' 0.4 + tr 1250 ¡E T + 2000 0.2 i . 3000

0.0 0 1000 2000 3000 4000 lt2 2l

Figure 6.4b Reduced sorption curves for TPCL urethane acrylates containing PCL soft segments of molecular weights given above. Samples equilibrated at9\Vo r.h. . 157

t.2

1.0 cr qt .. rl?'lÉf "t % o glD 0.8 8 t**d À + g 0.6 a ¿ a+tr E + 0.4 o +tr tr (650)2 o *q + (1m0)2 0.2 tr o Ë (2000)2

0.0 0 1000 2000 3000 4000 lt2 2l

Figure 6.4c Reduced sorption curves for TET(PTMO)2 urethane acrylates containing PTMO soft segments of molecula¡ weights given above. Samples equilibrated at987o r.h.

t.2

1.0 qü. .tt tg, tt 8' a ry 0.8 oo.¡* o + 8 o + ¿l=< o d FT 0.6 E + ¿ t + a E 650 0.4 +* trtr + r9 EI 1000 I EE o 2000 0.2 l"' otr . 2900 0.0 0 1000 2000 3000 4000

¡L 12 2l

Figure 6.4d Reduced desorption curves for TET urethane acrylates containing PTMO soft segments of molecula¡ weights given above. Samples initially equilibrated at987o r.h. 158

1.2 qit 1.0 .|lr å¡ Sa o ot a + F + 0.8 8 + |

Figure 6.4e Reduced desorption curves for TPCL urethane acrylates containing PCL soft segments of molecular weights given above. Samples initially equilibrated at 98Vo r.h. .

t.2

1.0 S+ ¡O +rr rw 5* .r* '+ .ßF 0.8 o+ 8 a+ !r a ¿ + !r 0.6 a ¿ + o o + 0.4 E (650)2 a +E + (1000)2 Or E 0.2 'd o (2000)2 EI å 0.0 0 1000 2000 3000 4{n0 ¡112 2l

Figure 6.4f Reduced desorption curves for TET(PTMO)2 urethane acrylates coñtaining PTMO soft segments of molecular weights given above. Samples initially equilibrated at98Vo r.h. . 159

D6 for each polymer decreased fo¡ the homopolymers sorbed at 98Vo r.h. with the sole exception being the TPCL3000 sample where D¿ (13.8 x 108 cm2 sec-l) decreased relative to D, (18.2 x 108 cm2 sec-1). Correlation coefficients obtained in determining ns and nd were between 0.97 and 0.998 - improved fits were obtained relative to the samples conditioned at lower humidity. ns and nd showed wider variations at this relative humidity than at 79Vor.h. with the largest nd being that for the TPCL2000 polymer. In general, however, values of n¿ exceeded those of n, throughout each polymer series indicating greater departures from Fickian behaviour for the desorption runs. Increases in water contents over the EWC values for samples conditioned at 79%or.h. were pronounced in the TET polymer series with samples at 98Vor.h. taking up about 207o more water. TPCL polymers showed an increase in EV/C of close to I)Vo, while the TET(PTMO)2 samples showed negligible further water uptake compared to the 79Vor.h.

samples.

6.3. 1 (dlSorption in water

S orption and desorption res ults for homopolymers in w ater are given in Table 6.4, with reduced sorption and desorption curves shown in Figures 6.5a-f.

All sorption/desorption plots showed less scatter for the samples equilibrated in water compared to humid environments - since EWCs were highest for samples in water reduced errors involved in sample weighing may account for this. 160

t.2

1.0 o dFo ooQ ++û + + G ?or q¡r/ åE o o +* trtr 0.8 o + E I t6) + H o o +* ¿ o + #*"" 0.6 o àF o& + E ¡o .+ 6tr E 650 o+ otr o 0.4 o + + 1000 oFd o 2000 0.2 . 2goo

0.0 0 1000 2000 3000 4000 lt2 2l

Figure 6.5a Reduced sorption curves for TET urethane acrylates containing PTMO soft segments of molecula¡ weights given above. Samples equilibrated in water .

1.2

{9o to"o 1.0 rE b + eþE'}sre ot o q+ ++ + + o 0.8 + o 8 EI F o ""tt'" à + o 0.6 à o+ O +tt 0.4 + E¡ 1250 ,tr ofE + 2000 o++¡ 0.2 æ É o 3000 rt 0.0 0 1000 2000 3000 {t2 2l

Figure 6.5b Reduced sorption curves for TPCL urethane acrylates containing PCL soft segments of molecular weights given above. Samples equilibrated in water. 161

r.2

1.0 a #+-tU o û a 0 o dF Ed 0.8 o #te 8 o tr ¿= o.¡ F 0.6 o aE à +tr o [t 0.4 t.-t E (650)2 o+o + (1000)2 qrtr 0.2 qt o + (2w)2

0.0 0 1000 2000 3000 4000 5m0 út2 2T

Figure 6.5c Reduced sorption curves for TET(PTMO)2 urethane acrylates containing PTMO soft segments of molecular weights given above. Samples equilibrated in water.

t.2

g. r.0 oa sÞn 8, gf a o O a a oo + + E + 0.8 ¡O + 8 oo ¿F oo + E ¿r. 0.6 o a o + 650 + o+ tr 0.4 + 1000 a + o + o 2000 o 0.2 o+ o . 2900 É+" 0.0 0 1000 2000 3000

¡112 2l

Figure 6.5d Reduced desorption curves for TET urethane acrylates containing PTMO soft segments of molecular weights given above. Samples initially equilibrated in water. L62

t.2

1.0 El 4D Ebr EDÞ óq ôq6 EI o 0.8 + g 8 t F¿ +Eo +OE rl 0.6 E ¿ e8 + 0.4 .f E 1250 +9 + d 2000 0.2 t o 3000

0.0 0 1000 2000 3000 ¿1000 út2 2l

Figure 6.5e Reduced desorption curves for TPCL urethane acrylates containing PCL soft segments of molecular weights given above. Samples initially equilibrated in water .

t.2

1.0 sFdö Eð o@o ::?ffË o+o+ "oF 0.8 o+ + 8 + tr ¿ o + E o + ÀF 0.6 df E tr (650)2 0.4 # + (1000)2 tr 0.2 +b o (2000)2 E¡

0.0 0 1000 2000 3000 4000 ¡r12 2t

Figure 6.5f Reduced desorption curves for TET(PTMO)2 urethane acrylates containing PTMO soft segments of molecular weights given above. Samples initially equilibrated in water. 163 TABLE ó.4 Sorption resuls for TET, TPCL and TET(PTMO)2 polymers equilibrated in water. The standard error for D5 and D¿ given by the DATAFT progrlam is shown in brackets. Polymer Ds (xl D¿ (xl Dg nd (7o) cm2sec-l cm2sec-1 ïBI'650 3.I 2.66t 0.48 0.45 TETlU}O 2.8 0.48 0.5u 'tEI-zUX, 2.U I Ó.7, u.3 17.6(0.3 0.48 0.52 TET29OO 2.O 32.0(Z.t) tJ.41 0.45

TPCL125O 1.8 7.51 0.47 0.45 IPCLz000 l.ó r0.1 .0.7 r3.9(0.3 0.50 0.53 TPCL3OOO 1.5 I r 1.7(0.3 0.41 0.5ó

TET(650)2 3.0 ) 0.40 0.46 'r E r'(tu)0)2 2.5 0.6 l4.l(u.3 o.47 0.52

TET(20m)2 2.1 23.Ot I I 26.O(tJ.4 0.46 0.48

Sorption coefficients Ds increased markedly for TET2900 and

TPCL3000 samples while any increases for the other polymers were small. Desorption coefficients D¿ were generally close to those for Ds, with the exception of TET2900, where D¿ (32.0 x 108 cm2 sec-l) was greater than Ds Q4-4 x 108 cm2 sec-l), and TPCL3000 where D¿ (1I.7 x 108 cm2 sec-l) was approximately half that of D5 (22.0 x 108 cm2 seò-t). Values for both n5 and ns were appreciably closer to 0.5 for the polymers equilibrated in water than for any of the samples equilibrated in humid environments - indicating that both sorption and desorption mechanisms are close to Fickian. Correlation coefficients exceeded 0.99 for all samples except the TET(650)2 polymer.

The absolute EWC decrease for desorption generally tended to be Iess than the EWC increase on sorption, with the discrepancy between water taken up and rwater lost from the polymer decreasing in sorbing environments with higher relative humidity and in water (Table 6.5). 164

TABLE 6.5 EWC decrease on desorption expressed as a fraction of EWC increase on sorption for homopolymers at three different relative humidities and in water.

en 33Vor.h. 79Vor.h 98Vor.h. water TET65O 0.7 0.7 U.E 0.E 'IEl'100u 0.8 0.8 0.8 0.9 TET2OOO o.l 0.ó 0.8 t.l 1'b'r2900 0.8 0.8 0.9 1.1 .IPCLI25LI 0.7 0.8 0.9 0.9 TPCL2OOO 1.0 o.7 0.9 0.9 TPCL3OOO 0.8 0.8 0.8 0.9

TET(650 t2 0.6 0.8 t.0 1.0 )2 0.9 0.1 1.0 1.0 t2 o.l 0.8 1.0 1.0

6.3.2 Variation of Ds and D¿ with water uptake Sorption and desorption coefficients for the TET homologous series at 33%or.h. \ilere calculated using Equations 6.2 and 6.3 and plotted Mr against -'-- (Figures 6.6a and b).

Sorption coefficients for TET650 and TETl000 both showed little change through the sorption run. TET2000 and 2900 showed large decreases in Ds uft". ff = 0.6, with the final sorption coefficients for these polymers being between a third and a quarter of their maximum values. Desorption coefficients for the TET series all showed initial increases up ,o # = 0.5, thereafter values for TET650 and 1000 remained relatively constant. D¿ for TET2000 and 2900 polymers decreased abruptly urt". ff = 0.9. 165

4.W,-1 E 650 a + 1000 o o 3.00e-7 ¡O a 2000 a o 2900 o o ooogoo a v, o o a z.we-'t o o o

o 1.00e-7 ++ + + ++ o a .r.* +++ +++q.üi oO Etr trEE I o E El EEEtro Et o E a 0.00e+0 0.0 0.2 0.4 0.6 0.8 1.0 Mt/M"" Figure 6.6 (a) Variation in sorption coefficient, D., with fractional weight uptake for TET ureth ane acrylates equi librat ed at 337or.h..

5.00e-7 Er 650 o a o + 1000 a 4.00e-7 o o a o 2000 o o 3.00e-7 . 2900 a o o ! o a a o o 2.ffie-7 a o o + + +++ * .1' 1.00e-7 o + + ++ E Eft EE + o s g tr s O tr 0.00e+0 0.0 0.2 0.4 0.6 0.8 1.0 Mt/M-

Figure 6.6 (b) Variation in desorption coefhcient, D¿, with f¡actional weight loss for TET urethane acrylates i niti ally equili brated at 331or.h.. 166

6.3.3 Copolymer sorption in water. 6.3.3 (a) TET650 copolymers Sorption results for copolymers of TET650 containing the indicated weight percentages of methyl acrylate, TEGDA and HDDA are shown in

TABLE 6.6 Sorption results for TET650 copolymers equilibrated in water. M, H and T represent methyl acrylate, HDDA and TEGDA respectively. The standard error for D5 given by the DATAFT is shown in brackets. Polymer EWC CalculatedbWC Ds (xl0u) cmzsec-r fl5

TET65O 3.1 ) 0.48

'lBl'650M2t) 2.9 2.8Ur U.U3 0.50 TET65OM5O 2.7 3.40(0.05) 0.49

'l.b'l'ó50I20 3.5 4.2 l.EEt 0.02 0.47 'lEr.ó50f50 4.7 6.0 1.70r 0.01 0.49

T100 8.8 7 0.05 I 0.50

'rh,r'ó50H20 2.5 2.8 2-02t 0.02 0.47

TET65OH5O 2 2.4 1 0.02 ) o.49 HIOO 1.6 2.24t 0.08 0.41

Table 6.6. Reduced sorption curves for this copolymer series are shown in Figures 6.7a-c.

The calculated EWC was obtained by assuming that the EWCs of the constituent polymers in the copolymer were additive. Sorption coefficients and EWC were not calculated for poly(methyl acrylate) homopolymer - when photopolymerised pMA was removed from casting

sheets poor dimensional stability and extremely uneven surfaces rendered the samples unsuitable for gravimetric sorption experiments. The

mechanical integrity of copolymer samples containing up to 50wtVo methyl acrylate was satisfactory, however, enabling sorption runs to be performed. t67

M. Plots of log -m- against log t (seconds) gâve ns values with correlation coefficients exceeding 0.99 in all cases. With the sole exception of the HDDA homopolymer, all values of ns were close ,ol 0., indicating predominantly Fickian sorption in the copolymers.

Dr increased slightly on copolymerisation with methyl acrylate from 2.66 x 108 cm2 sec-l in TET650 homopolymer to 3.40 x 108 cm2 sec-l in the 50/50 copolymer. EWC decreased slightly on copolymerisation with methyl acrylate.

Copolymers of TET650 containing 20 and 5ÙwtVo TEGDA showed a decrease in D5 from that for both TET650 and TEGDA homopolymers. Calculated EWCs were 0.7 and 1.37o higher than the experimentally determined EWCs for the 20 and 50wtVo TEGDA samples respectively.

TET650 co HDDA copolymers also exhibited lower values for Ds when compared to the constituent homopolymers, with the diffusion coefficient for the TET650co507oHDDA sample being only 60Vo of that for the TET650 homopolymer. Calculated EWCs, while being 0.3-0.47o higher

than the experimental EWCs, showed less deviation than for the

equivalent wtToTEGDA copolymer sample. 168

1.2

1.0 .,o os d#'o Eû 3ao 0.8 a 8 à Àl

0.4 -

s 20 wtVo 0.2 ,o*- . 50 wt%o

0.0 "Ê 0 1000 2000 3000 4000 ¡u2 2l

Figure 6.7a Reduced sorption curves for copolymers of TET650 containing methyl acrylate in the proportions given above. Samples equilibrated in water .

t.2

1.0 OO E.Po o$o 0.8 8 ¿FT 0.6 $ ¿ gs

0.4 g E 20 wtVo 0.2 "t" ' 50 wt%o I ¡ " 0.0 0 1000 2000 3000 4000

¡112 2l

Figure 6.7b Reduced sorption curves for copolymers of TET650 containing TEGDA in the proportions given above. Samples equilibrated in water . 169

1.2

1.0 a a 0.8 É¡. 8 lr:¡ o'oo_OE , ¿ Ea 0.6 EE À a

0.4 t E 20 wt%o sB 0.2 EB ' 50 wt%o I 3 0.0 0 r000 2000 3000 4000 {t2 2l

Figure 6.7c Reduced sorption curyes for copolymers of TET650 containing HDDA in the proportions given above. Samples equilibrated in water .

t.2

1.0 @D Ð ss4 t¡ ;st o'oeo 0.8 ^do o 8 Sooo àF à 0.6 o oog o EI E 0.4 5 o ZjwtVo

0.2 I 50wf7o Ê '

0.0 0 1000 2000 3000

¡t12 2l

Figure 6.8a Reduced sorption curves for copolymers of TET1000 containing methyl acrylate in the proportions given above. Samples equilibrated in water. t70

6. 3.3 (b) TETl000 copolymers Sorption results for TET1000 copolymers are shown in Table 6.7, with the corresponding reduced sorption curves being shown in Figures 6.8a-c.

TABLE 6.7 Sorption results for TET1000 copolymers equilibrated in water. M, H and T represent methyl acrylate, HDDA and TEGDA respectively. The standa¡d error for Ds given by the DATAFT progmm is shown in brackes.

Polymer EWC Calculated D5 (x108) cm2sec-l fì5 (7o) EWC

TETIOOO 2.8 I 0.48 -lEI-t0u)M20 2.3 0.45 IETlUÐM50 3.1 5.3(0.4) 0.41

TET1OOOTlO 2.8 3.4 6.3( 0.1 0.45 'rï.r'r000120 3.1 4.0 4.7e(0.0e) 0.47 TETlOOOT3O 3.6 4.6 3.81( 0.48 TETIOOOT4O 4,7 5.2 3. r8(0.04 0.49 TETlOOOT5O 4.9 5.8 2.65t U.U3 U.5U T100 8.8 2.211 0.0s) 0.s0

TETlOOOHIO 2.4 2.7 7 01 0.47 TETlOOOH2O 2.3 2.6 5.7t Lì. I 0.46 'r'þ't'1000H30 t 2.2 2.4 4.8t 0l 0.44 TETlOOOH4O 2.1 2.3 4.5( 0.1 0.45

TETlOOOH5O 2.0 2.2 3.7 t 0.2 0.43 H100 1.6 2 0.08) 0.41

Dt values decreased on copolymerisation with methyl acrylate to a

value of 5.3 x 108 cm2 sec-l for the 50/50 copolymer. EWC in the 20wtVo methyl acrylate copolymer initially decreased from the TET1000 homopolymer value, then increased to above 3.|Vo in the 50/50 copolymer. t7t

t.2

1.0 t- ÐÂ qÞa ç?a 4 w" Â 0.8 8 lF( ¿ E¡ 70 wtVo 0.6 ¿F EI + + 2O wtVo o .# 0.4 a o Â 30 wt%o

. 4O wtVo 0.2 o Â s 50 wt%o 0.0 0 1000 2000 3000 4000 5000 ¡u2 2t

!18ure 6.8b Reduced sorption curves for copolymers of TET1000 containing TEGDA in the proporrions given above. Samples èquilibrated in warer.

t.2

+ 1.0 Ð qb doa tr I A Â 0.8 8 fl L à E l0 wtVo 0.6 ¿ + I Â 20 wtVo 0.4 o ufll 30 wtTo I t 40 wtVo 0.2 # ^ 50 wtTo 0.0 0 1000 2000 3000 4000 s000 ¡112 2l

Fig.ure 6.8c Reduced sorption curves for copolymers of TETI000 containing HDDA in the proportions given above. Samples equilibráted in water. l't 2

TETl000/TEGDA copolymers showed an initial small increase in D. in the

TET1000 co I}VoTEGDA samples, thereafter Dr rapidly decreased to a value, in the 50/50 copolymer, very close to that for the TEGDA homopolymer. Values of the coefficient n, showed more variation in all three TET1000 copolymer series than for TET650 copolymers. Deviations from the Fickian value of nr= 0.5 were most prominent for the TET|000 co 50Vo methyl acrylate copolymer (ns = 0.41) and for the TET1000 copolymer samples with greater than llVo HDDA, where ns ranged from 0.43 to 0.46. The TET1000 co TEGDA series showed no increase in EWC for the l\Vo TEGDA copolymer sample, thereafter EWC increased to 4.9Vo in the 50/50 copolymer sample. Calculated EWCs varied from 0.6 to l.l7o higher than the experimentally observed values, with the discrepancy being lowest for the lO7oTEGDA sample. Experimental deviations from calculated EWC values were considerably smaller in the TETl000Æ{DDA copolymer series (0.2-0. 3Vo EWC).

6.3.3 (c) TET2000 copolymers Sorption results for TET2000 copolymers containing the indicated weight percentages methyl acrylate, TEGDA and HDDA are shown in Table 6.8, with reduced sorption curves shown in Figures 6.9a-c.

S orption coefficients decreased on copolymeris ation for all three copolymer series, with the greatest decrease occurring in the TET2000 co

507o TEGDA sample where Ds = 4-16 x 108 cm2 sec-l' Ns coefficients were very close to 0.5 for all copolymers, with correlation coefficients

exceeding 0.99 in all cases. 173

1.2 t3.$or¡or 1.0 lrolltl¡ a.# 0.8 8 d Èr troo à tra lF{ 0.6 EO ¿ Ea Eo 0.4 a E a o 20 wt%o E 0.2 I . 50 wtVo

0.0 0 r000 2000 3000 4000 út2 2l

Figure 6.9a Reduced sorption curves for copolymers of TET2000 containing methyl acrylate in the proportions given above. Samples equilibrated in water.

1.2

1.0 G it!Â ?a ça E

0.8 E 8 ¿g s* o I0 wtVo E.o+ 0.6 E+ à Oa + 20 wt%o Etfo¡ A trtooa 0.4 o 30 wt%o ' 40 wtVo 0.2 Â 50 wt%o

0.0 0 1000 2000 3000 4000 5000 ¡r12 2L

Figure 6.9b Reduced sorption curves for copolymers of TET2000 containing TEGDA in the proportions given above. Samples equilibrated in water . tl4

t.2

1.0 É)ô d)t ç-o? + Iws@+ + 0.8 + 8 5À +o oÂa Gr lO wtVo àF 0.6 q,. oô + 20 wt%o 9o 0.4 o 3O wtVo "3^ dÂ 40 wt%o 0.2 ' Bi Â 50 wt%o

0.0 0 1000 2000 3000 4000 5000 lt2 2l

Figure 6.9c Reduced sorption curves for copolymers of TET2000 containing HDDA in the proportions given above. Samples equilibrated in water .

t.2

oElañ 1.0 ,a"o E-' o Bú' +qËttJ', d 0.8 tr 8 E F< o ¿ E ao 0.6 tr àt=r Eo o EO 0.4 Oo tr tr¡ 20 wt%o 0.7 o ' 50 wt%o

0.0 0 1000 2000 3000 4000

¡112 2l

Figure 6.10a Reduced sorption curves for copolymers of TET2900 containing methyl acrylate in the proportions given above. Samples equilibrated in water. r75

TABLE 6.8 Sorption results for TET2000 copolymers equilibrated in water. M, H and T represent methyl acrylate, HDDA and TEGDA respectively. The standard error for D5 given by ttre DATAFT is shown in brackes. Polymer EWC (7o) Calculated Ds (x10u) cmzsec-r fl5 EWC TET2OOO 2.0 16.11 0.3 0.48 .IEIzUUOMzU 2.2 13.2t 0.4 0.47 l E l2U)0M50 2.3 9.8r 0.48

1-ET2000'I'10 2.3 2.7 l3.Ur tJ.7 0.48 TET2OOOT2O 3.0 3.4 9-710.2 0.47 TET2üNT3O 3.6 4.0 7.2t 0.48 'lB'I20U)'140 4.r 4.7 5.0ór 0.u9 o.49 TET2OOOT5O 4.8 5.4 4.1 6r 0.06 o.49 T100 8.8 2.27t 0.05 0.50

TET2OOOH1O 1.9 2.0 0.3 0.4ó 'IB't2000Ft20 T.7 L2.7 t 0.7 0.50 TET2OOOH3O 1.8 1.9 0.47 1b'12000H40 t.8 1.8 '1.9t tJ.z 0.47 TET20OOH5O l.E 1.8 1.Ot 0.47 H100 1.6 2.24t (.).08 0.4I

EWCs calculated from simple additivity of copolymer constituent EV/Cs were generally closer to the experimental EWCs than for the TET650 and 1000 copolymers. Calculated EWCs for HDDA copolymers were only 0.1-O.27o high for the 10-30 wtVo HDDA samples and did not differ from the experimental values for the 40 and 50wtVo HDDA copolymers. Calculated EWCs for TET2000/TEGDA copolymers were again higher than the experimental values (0.4-0.67oEWC).

6.3.3 (d) TET2900 copolymers Sorption results for TET2900 copolymers equilibrated in water are given in Table 6.9. Reduced sorption curves for TET2900 copolymers are plotted in Figures 6. l0a-c. t76

t.2

1.0 # tt e. ffB-* o tr 0.8 8 E4 o ¿ + *o a ts 0.6 s *o¡ tr ¿ E+oo l0 wtVo

+ 20 wt%o 0.4 "*o' "-oo;O o 30 wt%o o.z E + f o 40 wtVo o

0.0 0 1000 2000 3000 4m0 5000 lt2 2T

Figure 6.10b Reduced sorption curves for copolymers of TET2900 containing TEGDA in the proportions given above. Samples equilibrated in water.

t.2

qt o 1.0 É# E tdb E t foo ot o o a oo 0.8 O 8 tr à!r a àF 0.6 tr o o 0.4 a

tr o tr 20 wt%o 0.2 o o ' 50 wt%o

0.0 0 1000 2000 3m0 4000 l/2 2l

Figure 6.10c Reduced sorption curves for copolymers of TET2900 containing HDDA in the proportions given above. Samples equilibrated in water . 117

TABLE 6.9 Sorption results for TET2900 copolymers equilibrated in water. M, H and T represent methyl acrylate, FÐDA and TEGDA respectively. The standard error for D5 given by ttre DATAFT is shown in brackes. Polymer EWC (7o) Calculated Ds (xlOE) cm2sec-l rlg EWC TE12900 2.O 0.41 .

1ET2900M20 z.I 18. 1t 0.9 o.42 TETz9OOM5O 2.1 I l.4r 0.2 0.41 'rB'r29tx.t'l'10 z.z 2.7 19.4t 0.8 0.45 'rE12900T20 2.8 3.4 Il.2t 0.5 o.46 TETz9OOT3O 3.4 4.U E.ót tJ.2 U.5U TET29OOT4O 4.2 4.7 6.5r 0.I 0.48

Tt00 8.8 2.27 ) 0.50

TET29OOH2O t.0 I.9 16.7 t o.l 0.42 'rE1.2900H50 1.6 1.8 6.6r 0.2 0.50 H100 1.6 Z.Z4t 0.0E 0.41

Sorption coefficients for the TET2900 copolymers showed similar trends to those seen in the TET2000 copolymers i. e. a continuous decrease as the wtVo methyl acrylate, TEGDA or HDDA in the copolymer increased-

Ns coefficients calculated from log-log plots versus t "t # showed more variation than for TET2000 copolymers, ranging from 0.42 up to 0.50. Calculated EWCs again slightly overestimated the water uptake for the TEGDA and HDDA copolymer, although in common with the TET 2000 copolymers, the extent of this overestimation was substantially less than for the TET650 and 1000 copolymers.

6.4 Discussion

S trict compliance of the homopolymers with the three crltena outlined fo¡ Fickian diffusion was not observed The reduced desorPtion 178 curve invariably lay above the reduced sorption curve, particularly for polymers containing PTMO soft segments. Only one sample, TPCL3000, met the condition that D, exceed D¿ in all environments. This sample also showed an obvious tendency to decrease in weight after maximum weight increase had been initially achieved. The sorption of water above the final uptake value has been observed in a number of other systems Ll'12 and has been analysed by Vrentas et. al. 13.

This decrease in water uptake (termed "overshoot") has been attributed 14-16 to crystallisation in the polymer caused by the solvent, with denser crystalline regions rejecting solvent already sorbed. Although this overshoot phenomenon has been observed in non- crystalline systems 17. DSC measurements carried out on TPCL3000 samples (Chapter Four) indicate the presence of crystallites in this polymer :- formation of crystalline regions during sorption may account for the observed overshoot

The values of D¿ for TPCL3000 also showed a tendency to decrease relative to Dr as the relative humidity of the wet environment increased. This may be a reflection of the experimental techniques used :- sorption runs commenced within an hour of sample removal from a 50"C vacuum oven, while desorption took place after samples had been sorbed at 25oC for a number of days. Crystallinity probably developed during the sorption run but may have existed prior to the commencement of desorption runs, with crystalline regions slowing the removal of water from the polymer by increasing network tortuosity 8. Compliance of the TPCL3000 polymer with the Fickian criteria outlined is probably fortuitou s. Values of n. and n¿ were both close to 0.5 for samples sorbed in water - ñs was somewhat lower in humid atmospheres, while values of n¿ tt9 showed large variations, generally exceeding the value of n=0.5 characteristic of Fickian diffusion.

other workers l8 found opposite trends in ns and nd for water sorption in copolymers of pHEMA with oligo(ethylene glycol) dimethacrylates : high values of n, were often associated with low values of n¿. A number of copolymers had desorption coefficients which were lower than Dr - the opposite to the general trend for the urethane acrylate polymers.

Desorption coefficients for water in crosslinked PMMA networks 19, in common with urethane acrylates, exceeded those for sorption. Turner 2 considered that swelling in PMMA resulted from processes with long relaxation times, and continued after the intial solvent front had moved through the polymer. Differences in the relative magnitudes of D5 and D¿ may result from different rates of network swelling by imbibed water, with network relaxation due ro swelling in hydrophilic networks occurring faster than the Fickian diffusion component. In polymers with a low proportion of polar, water solubilising groups any network swelling would be slowed as the water first has to diffuse to hydrophilic sites before swelling can occur.

Diffusion coefficients calculated from curve fitting (Equation 6. 1) decreased approximately linearly with crosslink density. Figure 6. 11 shows the values for D, plotted as a function of molecular weight between crosslink's, Mç, for TET, TPCL and TET(PTMO)2 homopolymers at 987o humidity.

Similar trends in D have been noted in orher polymers at low to moderate degrees of crosslinking 20'21,1.

Linear extrapolation of the data in Figure 6. l l to zero molecular weight between crosslinks would yield a negative diffusion coefficient;in common with other networks 22 D, shows a non-linear decrease at high 180

20 o (t) a

fl r0 o +

0 0 1000 2000 3000 4000 5000 Mc Figure 6.11 Variation in sorption coefficient, Dr, calculated from Equation 6.l,with molecular weight between crosslinks for TET, TET(ruMO)Z and TPCL urethane acrylate homopolymers sorbed at 98Vo r.h.. M. was taken as the total molecular weight of soft segment between TDIÆ{EA end groups in the prepolymer. 181 crosslink densities. The slopes of the D. versus Mc plots are similar for rhe rwo different types of soft segments, suggesting that at this relative humidity the chemical composition of the polymers does not influence the rate of water sorption.

Ds obtained from Equation 6. 1 in the TET copolymer series containing TEGDA and HDDA generally decreased as the proportion of low molecular weight crosslinker in the copolymer increased. The variation of D, with methyl acrylate content was not as consistent as that seen for TEGDA and HDDA. TET650 copolymers showed a slight increase in Ds with increasing methyl acrylate content, TET1000 copolymers showed a slight decrease, and TET2000 and 29OO copolymers

both showed significant decreases in D, as the copolymer methyl acrylate content increased. This may reflect the network diffusion coefficient approaching that for methyl acrylate homopolymer.

A number of reduced sorption curves showed two-stage sorption behaviour where the initial steeply sloped region was followed by a second nearly linear region of reduced slope, which levelled off to the final equilibrium. This two-stage behaviour was most apparent for the TET2000, TET2900 and TET(650)2 samples at 33Vor.h. and resulted in larger standard errors in D, calculated from the DATAFT program. Two

stage behaviour was absent in desorptlon runs.

Turner 2 observed pronounced two-stage sorption behaviour for the sorption of water in PMMA. This non-Fickian sorption behaviour was accounted for by initial accommodation of the water in microvoids, followed by a slower redistribution of water accompanied by network

swelling. Sorption of gases and vapours at low activities in epoxy based

polymers has been described by a s uperposition of Henry's Law and t82 Langmuir isotherms 23, where the Henry's law term is generally attributed to molecular solution of penetrants in the glassy matrix, and the Langmuir sorption mode occurs as a result of the diffusion of gas molecules into pre-existing gaps or voids in the polymer 24. These systems differ considerably from the urethane acrylate polymers considered here in that PMMA and epoxy polymers are glassy at room temperature. The diffusion of organic vapours into polymers well above the glass transition generally obeys Fick's Law 6, so clearly an alternative mechanism to void filling must be invoked to account for the observed two-stage sorption behaviour.

A considerable number of studies have focused on the diffusion of gases in polyurethane elastomers. Mc Bride et. a1.25 investigated gas diffusion as a function of temperature in a series of polyurethane block copolymers. Discontinuities were observed in Arrhenius plots for 02 and COz diffusion through polymer membranes, with increases in gas diffusivity above a certain temperature being correlated with the onset of the glass transition temperature in hard segment domains. Incremental vapour sorption in a phase segregated polyurethane elastomer 26 with a strongly swelling solvent (ortho dichlorobenzene, ODCB) also showed pronounced two-stage sorption behaviour. The non- Fickian sorption curves resembled experimental results obtained by Berens ând Hopfenberg 21 for vapour diffusion in glassy poly(vinyl chloride) and polystyrene microspheres. Two stage sorption behaviour

was treated as the sum of independent Fickian and relaxation processes :

M(t)=Me'(t)+Mp(t) 183

where Mr.(t) is the Fickian contribution and Mn(t) is the relaxation controlled contribution, described by a first order rate equation :

Mp(t) = Mn(-) [ 1- exp ( - kn t) ] where kp is the relaxation rate constant.

S chneider 26 concluded that the observed two-stage sorption behaviour resulted from the swelling of glassy hard segment microdomains at high solvent concentrations. Important differences exist between the results obtained by

Schneider et. al., and those for the sorption of water in the urethane acrylate networks considered here. The reverse trend in two-stage sorption behaviour was found for the urethane acrylates, with two-stage sorption being absent in all samples equilibrated in water and at high relative humidity. Two-stage sorption behaviour was also not observed for all polymers, with reduced sorption curves for TPCL 1250 and TPCL 2000 at33Zor.h. being linear up to s-= 0.8. (The TPCL3000 polymer is omitted from discussion here due the complicating effects of polymer crys tallinity). Apparent inconsistencies were also found in values for EWC in the homologous series and in the copolymers. While the EWC decreased with increasing soft segment molecular weight for both PTMO and PCL soft segment polymers, the extent of this decrease was generally more marked for the samples containing PTMO units. EWCs in urethane acrylate copolymers were consistently less than those calculated from simple additivity of component EWC values, with deviations from calculated EWCs being greatest in TET 650 and TET 1000 copolymers. Petrik et.al. also found that simple additivity of contributions from hydrophilic poly(ethylene oxide) (PEO) and hydrophobic poly(propylene oxide) (PPO) was not observed for water vapour sorption in oligomeric 184 triblock copolymers of PEO and PPO.28 Calculated water uptake values derived from the values for PEO and PPO monomers overestimated the water uptake in the copolymers, particularly in those with high PPO contents and at low water vapour activities. Petrik et. al. were able to describe the observed sorption behaviour in terms of the Zimm-Lundberg theory 29. This treatmenr defines a cluster integral, Grr/Vt, which gives a measure of the clustering trend of the solvent molecules, and a mean cluster size f1 (the mean number of solvent molecules per cluster). As the water concentration increased the clustering tendency decreased, and the size of the clusters increased. Hydrophobic propylene oxide units promoted water molecule clustering, while decreased clustering in PEO rich copolymers was due to the strong interaction of water with the ethylene oxide units. Petrik also considered 30 wate¡ sorption in segmented polyurethanes with MDI hard segments with low molecular weight PEO, PPO and triblock PEO/PPO as the soft segments. Water clustering was promoted both by the MDI and the propylene oxide units. Overall water uptake in the polymers decreased as the weight fraction of MDI units increased, the inverse situation to that found for weight fraction TDI in the PTMO and

PCL networks considered here. A combination of the hydrophobic PTMO units and the relatively hydrophilic TDI hard segments may promote water clustering in a similar manner to that seen in the triblock copolymers examined by Petrik. The difference in hydrophilic/hydrophobic character between TDI and PCL soft segments is probably less than thar for TDI/PTMO, leading to decreased water clustering tendency in the polymer.

The decreases in Ds (calculated from Equations 6.3 and 6.4) observed in the TET2000 and 2900 polymers sorbed ar 33Vor.h at values 18_5 .Mt of M- exceeding 0.6 may be due to clusrer formation Barrie and Platt

3l investigated warer vapour diffusion into PMMA samples already containing water. As the initial concentration of water in the PMMA increased, D decreased - this was attributed to decrqasing solvent mobility as the w ater molecu les increasingly participated in cluster formation 32' The values of D¿ (calculated from Equation 6.3) for the TET 2000 and 2900 urethane acrylates desorbed from 33Vor.h. showed sudden decreases uft". ff= 0.9, at a considerably later stage than the observed decrease in Dr.

As shown in Table 6.5, desorption generally did not remove all the water initially taken into the polymer in the sorption run. The proportion of water removed from the polymer by desorption did increase, however, as the relative humidity of the sorbing environment increased and total water uptake increased. This behaviour suggests that a constant amount of water is retained in the network on desorption, with the relative proportion of this more strongly bound water decreasing as the overall water content increases. The observed variation in D¿ with uptake may be accounted for by an abrupt slowing of desorption as the last non-bound water is removed from the polymer.

Water aggregation at specific sites in polymers is well documented, e.g. in epoxy resins 33-36 with specific interactions being observed with secondary amine and pendant hydroxyl groups 37. Typical linear polyurethanes are extensively hydrogen bonded 38 with the hydrogen bond donor being the NH group of the urerhane linkage. The hydrogen bond acceptor may be either in the hard segment (the carbonyl of the urethane group) or the soft segment ( an ester carbonyl or ether oxygen) 39'40 Interaction of water with urethane NH 186 groups has also been confirmed by dynamic mechanical measurements on wet Polyurethanes 4l '42.

The water uptake of the PCL soft segment networks was less than that of the PTMO soft segment networks despite the fact that the PTMO soft segment is more hydrophobic than the PCL soft segment.. Specific hard/soft segment interactions may account for this apparently anomalous water uptake behaviour. Schneider and Paik-Sung 43 attribured improved phase segregation in 2,6 TDI/PTMO polyurethanes compared to 2,6TDI/PBA (polybutylene adipate) urethanes to hydrogen bonding between urethane NH groups and ester groups being stronger than urethane-ether interactions. Kolarik 44 proposed that solvent sorption in PHEMA networks involved disruption of polar polymer-polymer interactions and subsequent replacement with polymer-solvent interactions. If urethane-ester hydrogen bonding was appreciably stronger than urethane-ether bonding, replacement of the polymer-polymer bonds by polymer-solvent bonds could occur to a greater extent in networks containing ether based soft links. If it is accepted that the urethane NH groups are the most likely specific sites for water absorption this would explain why the urethane acrylate homopolymers containing ester soft segments take up less water than analogous networks containing PTMO soft segments.

Urethane-ester and urethane-ether hydrogen bonding may also explain the deviations from calculated and experimental EWCs observed in the TET urethane acrylate copolymers. The discrepancies observed were generally greater in the TET650 and TET1000 copolymers, i.e. in those polymers with a greater concentration of urethane NH groups capable of hydrogen bonding with carbonyl and ether oxygens in the TEGDA and HDDA comonomers. 187

Non-Fickian water sorption behaviour observed under certain conditions for these polymers can be ascribed to water clustering at the urethane NH group, with observed deviations from calculated EWCs in urethane acrylate copolymers possibly being due ro changès in the accessibility of the NH groups caused by polar interactions with comonomer carbonyl and ether moieties. 6.5 Summary

Diffusion coefficients for the urethane acrylate polymers calculated from Equation 6. I showed a linear increase with crosslink density. Changes in relative humidity affected the water sorption kinetics with sorption coefficients calculated according to Frisch's relation, Ds, being near the Fickian value of 0.5 for samples at high relative humidities and in water.

Non-Fickian or anomalous sorption was observed for some samples in low relative humidity sorbing environments. Two-stage sorption curves were attributed to water clustering near urethane NH groups, with this slow relaxation sorption component being most apparent in networks with lower crosslink density where Fickian diffusion was rapid. 188 CHAPTER SEVBN

FTIR/NIR MBASURBMENTS

7.1 Introduction ATR-FTIR and NIR spectroscopy were used to determine the percent double bond conversion achieved in homopolymer samples, since the properties of crosslinked polymer networks can vary significantly, depending on the extent of reaction of methacrylate or acrylate end groups l. For standard transmission IR spectra sample preparation usually 2,3 consists of casting thin polymer films from volatile solvent mixtures. Such films, when used in transmission IR must necessarily be no greater than 100-200pm thick, limiting their use for other (e.g. mechanical testing). The Attenuated Total Reflectance (ATR) technique, in which the polymer surface is brought into contact with a crystal through which the infrared radiation travels, permits analysis of otherwise intractable opaque solid polymers.

Copolymer samples were examined using ATR in order to determine whether the decreases in Tg for TEGDA and HDDA observed in DMTA runs on copolymers resulted from any specific intermolecular interactions. FTIR has proven useful in the study of polymer blend miscibility where improved signal to noise ratio, higher energy throughput and rapid spectral scanning compared to conventional dispersion IR enables relatively small changes which occur as a result of interactions between the blend components to be detected 2.4,5 .

The sorption of water into the homopolymer networks was 189 examined using NIR techniques

7.2 Results and Discussion 7.2.L ATR Results The ATR technique involves bringing the polymer into "onru", with a crystal (referred to as rhe Internal Reflection Element (IRE) through which the laser beam travels. At the polymer/crystal interface total internal reflection occurs due to the difference in refractive index between the sample and the crystal, with the quality of the ¡esultant

spectrum being determined by the angle of incidence of the IR beam and the crystal/sample interface. If the angle of incidence of the beam is below the critical angle, 0., spectral distortion may result. The critical angle is a function of the sample refractive index n2 and the IRE refractive index nt (2.37 for the KRS-5 crystal used here):

n2 0s = 5i¡-11 nl

Sample refractive indices were close to 1.4, giving a critical angle of 36o. The angle of incidence used was 45o, enabling distortion free spectra to be obtained. Adequate physical contact between the sample and the IRE element is essential for adequate spectral resolution 6 ' generally the rubbery nature of the urethane acrylate polymers promoted good "wetting" of the crystal by the sample - additional clamping pressure was only required for copolymer samples with a high proportion of glassy copolymer. The ATR specrrum of TET 650 polymer is shown in Figure 7.1.

2 1n

Hydrogen bonding in polyurethanes is reflected in the absorption positions for the NH stretch between 3600 and 3200 cm-l 190 and the carbonyl stretching region (1750-1600cm-l). Qualitatively, the NH stretch centred at 3310cm-l in the TET650 polymer indicates nearly complete NH hydrogen bonding. This peak is, however, not suitable for quantitative data concerning the fraction of "free" non-hydrogen bonded and hydrogen bonded groups 7. The infrared carbonyl band envelope is composed of separate absorptions attributable to free and hydrogen bonded carbonyl groups with previous workers observing the bonded carbonyl peak at 1700cm-l and that due to free C=O at 1732cm' 11.

The presence of the acrylate groups in the urethane acrylate polymers hampers examination of carbonyl absorptions. To facilitate understanding of this region, 1 molar equivalent of PTMO650 diol was endcapped with TDI units, and then subsequently crosslinked with glycerol. The ATR spectrum of the resultant polymer is shown in Figure 7.2, from which, based on the assumption that the extinction coefficients of free and bonded carbonyl groups are about the same 8,9, it is apparent that there are slightly more bonded C:O groups than free C=O groups.

Bonded C=O groups have been regarded as being due to hard segments residing in the interior of hard domains, while free C=O groups are ascribed to hard segments present in a mixed soft phase or at the interface between hard and soft segments 10. 191

É960 tr U) 40 F.Ì: èa

4000 3500 3000 2500 2000 1 500 1 000 500 Wavenumbe. c--t

Figure 7.1 ATR spectrum of TET650 polymer sheet (45o angle of incidence)

1 708 1.5

1.0 oc) F -o ot< tr4 .t> .o

0.5

0.0

1800 1750 1700 1650 1600 1550 Vy'avenumb", .--t Figure 7.2 ATR spectrum of the carbonyl region of a glycerol cross-linked TET650 analog. t92

Hydrogen bonding can, however, also occur within a single urethane group with the NH group hydrogen bonding to the urethane

alkoxy oxygen ll. Lack of correlation between DSC endotherms and

hydrogen bond dissociation shown by infrared spectroscopy has led some workers to conclude that hydrogen bonding is not indicative of differences in polyurethane domain morphology 12.

While the C=O absorption of the model polymer suggests that hydrogen bonding occurs to a similar extent in the crosslinked state as in linear polyurethanes, it is uncertain how closely this system approximates a crosslinked acrylate network.

7.2. l. (b) Conversion The quantitative determination of percent conversion, by monitoring the intensity of the C=C stretch at 1636cm-1 was not attempted, chiefly

due to the presence of interfering aromatic C=C stretches at 1620 and 1600cm-1. Qualitative comparison of the double bond conversion on

+ either side of a 2mm thick TET650 polymer sheet was attempted in order to determine whether light screening by the photoinitiator itself l3 had any effect on the uniformity of sample cure. Guthrie et. al. l4 observed decreasing percent conversion beyond

initiator concentrations of l7o with samples 2mm thick or greater, due to this light screening effect. ATR specrra of both sides of a TET650

polymer sheet (Figure 7.3) after curing showed rhar the double bond

conversion on the side furthest away from the u.v. source was the same as that of the side nearest the lamp, confirming that any light screening due to the Irgacure 651 photoinitiator is insignificant at this sample thickness and photoinitiator concentration. 193

0.6

0.5

0.4

0.3

0.2

0.1

lß 1680 1660 1640 t620 1600 1580 1560 Vy'avenumb". .--t Figure 7.3 C=C stretching region (a)TET650 prepolymer (b) upper and (c) lower surface of TET650 polymer sheet. 194

7.2. 1 (c)Copolymers Figure 7.4a shows the ATR scale expanded spectra for pure TET2000 and HDDA polymers, together with the spectra of TET2000/HDDA copolymers with from 10-50 wtVo HDDA in the 1100- l25O cm-1 region.The band at 1154cm-l -uy be tentatively assigned to the C-O streteh in the HDDA acrylate ester group 15. As the concentration of HDDA increases, the band shifts from 117Ocm-l in the lÙwfVo HDDA copolymer to 1159cm-l in the 50:50 copolymer. The position of this peak in TET1000 and TET2900/HDDA copolymers was within 1 or 2cm-l of the values obtained for TET 2000 copolymers at the same HDDA wtVo. TET65O/HDDA copolymers did not give adequately resolved spectra, presumably due to the increasingly glassy nature of the network resulting in poorer contact with the crystal. The TET2000/TEGDA copolymer series with TEGDA compositions between 10 and 50wt%o, together with the spectra of TET2000 and TEGDA homopolymers, are shown in Figure 7.4b. The

TEGDA C-O absorption peak shifted from 1 161cm-l in the homopolymer to near 1170cm-l in the I}wtVo TEGDA sample, and occurred at 1164cm-l in the 50:50 copolymer sample. The TEGDA peak position in TET1000 and TETZ9}}ITEGDA copolymers was close to that fóund in the TET2000/TEGDA copolymers. The frequency shifts observed for the HDDA and TEGDA C-O peaks in the copolymers indicate some interaction with the urethane acrylate networks. The larger magnitude of the shift for HDDA (AV-u* = 15cm-l) at l0wt%o in copolymers compared to TEGDA (Avrax = 9cm- t¡ may indicate, at least semi-quantitatively, a greatcr dcgree of interaction of the HDDA with the urethane acrylate. 19s

t.4

1.2

1.0 o o É 0.8 cd -ot< ano -o 0.6

0.4

0.2

0.0 1154 ll80 ll70 1160 ll50 ll40 ll30 Wavenumber c--t Figure 7.4 a ATR spectra of (a) TET2000 homopolymer, (g) HDDA homopolymer and I0-50wt%o HDDA (b-Ð in TET2000ÆIDDA copolymers. Absorbance indicated is for spectrum (a) only.

t.2

1.0

0.8

e) 0.6

0.4 (b)

0.2

1 161 0.0 1200 ll80 1160 ll40 rt20 Wavenumber a--t Figure 7.4 b ATR spectra of (a) TET2000 homopolymer, (g) TEGDA homopolymer and 10-50wtVoTEGDA (b-Ð in TET2OO0||EGDA copolymers. Absorbance indicated is for spectrum (a) only. 196

Frequency shifts for peaks in the TET2000 were not observed, hence no conclusions regarding any specific interactions (e.g. hydrogen bonding) leading to compatibility could be made.

7 .2.2 N IR Res ul ts

7. 2. 2 (a) Conversion Double bond conversion was monitored by taking NIR spectra of the prepolymer in the casting assembly after u.v. irradiation for varying intervals. In calculating percent conversion, the area of the peak at 6161cm-l due to the CH2=Qg stretching overtone (Figure 7.5) was ratioed to the combined C-H band envelope between 6060 and 5600cm-1, with the initial ratio in the prepolymer being set to ÙVo convarsion. The intensity of the C-H band varied slightly ( t1.5Vo) in the series of spectra taken during the photopolymerisation ; band ratioing eliminated any inconsistencies in absorbance due to slightly differing path lengths for different spectral runs. The plot of percent conversion against time for TET650 prepolymer containing 3Vo wlw Irgacure 651 is shown in Figure 7.6. After a short induction period of about l5 seconds, presumably reflecting oxygen inhibition of the propagation reaction, the reaction rate increased rapidly up to about 807o conversion, then decreased. The increase.in polymerisation rate after about 35Vo conversion occurs as a result of the bimolecular termination of macroradicals becoming diffusion controlled 16. Since the initiation and propagation rates are unaffected at low conversion, decreasing radical termination results in an attendant increase in radical concentration with an increase in overall polymerisation rate. The decrease in polymerisation rate after 80Vo conversion can be attributed to the onset of r97

0.14

0.t2

0.r0 (l) o F 0.08 -oLr o.t) -o 0.06

0.04

0.02

0.00 6220 62W 6180 6160 6140 6120 6100 Wavenumber cm-l Figure 7.5 NIR spectra of TET650 prepolymer before (a) and after u.v.exposure for vanous times.

100 o E

80 o v) q) 60 E I 40 SR

0 0 100 zffi 300 time (sec)

Figure 7.6 Percent double bond conversion versus time for TET650 polymerisation. 198 diffusion control for the initiation and propagation reactions, either due to depletion of initiator or monomer, or to the propagation step largely involving macroradicals whose ends move by the "reaction diffusion" process L7. The most significant result here is that conversion approaching 95 Vo was attained after 5 minutes of irradiation. Conversion proceeded to l00%o ( i.e. the peak at 6161cm-l became indistinguishable from the baseline) for all other homopolymers. 7.2.2 (b) Water uptake

The very thin samples required for transmission IR spectra and changes in sample refractive index on sorption are two factors complicating the use of infrared spectroscopy as a means of probing the nature of the interaction of sorbed water with polymer networks. Workers have recently used NIR spectroscopy to investigate the interaction of water with epoxy resins 18. The NIR spectrum of dry TET650 polymer, together with those for the same sample after 1,3, 6,8 and 25 hours water sorption at 25oC are shown in Figure 7.7. The water peak af 5201cm-l increased in intensity with longer sorption time. The position of this peak was found to be the same for all the TET and TET(PTMO)2 polymers. Similar increases in water peak intensity with sorption time were also observed for the TPCL polymer series. Figure 5.8 shows that the position of the water peak in TPCL polymers differed, however, being found near 523lcm-"1 for all TPCL homopolymers, close to that found for epoxy resins (5223cm-l) 18. The change in the water peak position in the NIR in polymers with different compositions may result from differences in the extent of water-polymer interactions. r99

0.7

0.6 q) () F 0.5 st< 9 Ð 0.4

(b) 0.3

5400 s300 szm 5100 5000 Wavenumber cm

Figure 7.7 NIR spectra of TET650 polymer after (a) 0 (b) 1 (c) 3 (d) 6 (e) 8 and (Ð 25 hours water sorption at25C.

0.9

520r 0.8 523t 0.7 o F 0.6 -ok o

0.4

0.3

0.2 5500 5400 5300 5200 5100 5000 4900 4800 Wavenumber c--t Figure 7.8 NIR spectra of (a) TET650 and (b) TPCL1250 polymers after 25 hours water sorption at25C. 200

7 .3 S ummary

The ATR spectra obtained for a glycerol crosslinked TET650 analog suggested that crosslinked u¡ethane acrylates, in common with

linear polyurethanes, may be extensively hydrogen bonded. .

Photopolymerisation proceeded to 100 percenr double bond conversion in all samples, with the exception of TET650 homopolymer,

where ultimate conversion was found to be near 95To.

Shifts in peak positions for HDDA and rEGDA copolymers indicated some interactions with the urethane acrylate networks. Specific intermolecular interactions accounting for these peak shifts were not identified.

The position of rhe water peak for sorbed samples in the NIR changed depending on the urethane acrylate soft segment, possibly reflecting varying degrees of interaction between water and the polymer. 201 CHAPTBR BIGHT

CON CLUS IONS

Dynamic mechanical analysis (Section 3.2.1) of urethane-acrylâtes containing a single soft segment chain between crosslinks showed large decreases in the polymer glass transition temperature as the molecular weight of the soft segment was increased up to 3000 for both PTMO and PCL soft segments. Decreasing crosslink density for higher molecular weight PTMO soft segments in samples with two soft segment chains interconnected by a single TDI group between acrylate crosslinks, led to an increased tendency for the PTMO chains to crystallise. TET(PTMO)2 polymers showed considerable decreases in Tg for the polymers containing PTMO 650 and 1000 soft links. The Tg of TET(2000)2 was identical to that for the TET2000 and TET2900 polymers. The cx,¿ peak for the TET(2900)2 was shifted about l0'C upwards in temperature from that for the TET(2000)Z polymer; this, together with a substantial decrease in the tan ô peak height for the amorphous glass transition and the development of a crç peak at 10'C confirmed the presence of extensive crystallinity in the sample. A Tg approximating that of the anticipated hard segment glass transition in the urethane acrylates was obtained from averaging the Tg values for linear and highly crosslinked TDIÆIEA model polymers, both of which were glassy and brittle at room temperature. Application of the Fox equation (Equation 3.2) to the urethane acrylate networks, considering the polymers to be made up of a hard and soft phase, considerably underestimated the actual polymer Tgs.

A different approach involved considering the observed Tgs as reflecting constraints on soft segment chain motion by crosslink points 202 in the network, with the relatively bulky TDIÆIEA moiety considered to be part of the immobile crosslink. DeBenedetto's equation gave reasonable agreement between calculated and experimental glas s transition temperatures. The most significant departures from the crosslinking equation were for TET650 and TET(2900)2 polymers.

Extensive crystallinity in the TET(2900)2 polymer, not considered in the crosslinking equation may account for the lack of agreement between experimental and actual Tgs for this sample. Reasons for the overestimation of the TET650 polymer Tg are not obvious. Conversions for TET 650 measured by NIR (Section 7.2.2) were around 95Vo. One possible explanation for the lower than expected Tg for the TET650 polymer could be plasticisation by a small amount of residual monomer - since DSC measurements established the TET650 prepolymer Tg as being near -74oC presumably only a minor amount of residual monomer could account for a decrease in polymer Tg from the calculated value.

The improved prediction of polymer Tg by a model based solely on crosslink density, together with the fact that no higher temperature tan ô peaks were observed, suggests that hard domains are not present in these polymers. Considerable decreases in Tg for TET(650)2 and TET(1000)2 compared to TET650 and 1000 confirmed that the dynamic mechanical response of the urethane acrylate homopolymers is largely governed by the number of acrylate crosslinks, with the considerable decrease in Tg from TET1000 to TET2000 being adequately accounted for by decreased crosslink density, rather than improved phase segregation between hard and soft phases. Aggregation of hard segments into discrete domains does not occur, perhaps due to the acrylate crosslink point near the TDI group inhibiting cooperative hard segment orlentatlon. 203

Dynamic mechanical properties of TET 650 and TET 1000 copolymers with methyl acrylate, TEGDA and HDDA showed predictable changes based on changes in network crosslink density (Section 3.2.2). TET 650 and 1000 methyl acrylate copolymers showed decreased storage modulus with increasing methyl acrylate content. TET 650 and 1000 copolymers with TEGDA and HDDA showed the opposite trend, with the increased crosslink density being manifested in an increase in polymer storage modulus as the weight Vo of reactive diluent increased. Tan ô peak heights for the single Tg peak observed decreased for

TETl000/TEGDA copolymers from 10 to 4}wt%o TEGDA, then increased for the SOwt%o TEGDA copolymer. The tan ô peak height in the TETI000/50wtVoHDDA copolymer was slightly less than that of HDDA homopolymer, suggesting that a similar trend in tan ô peak height could apply to the TET1000/HDDA copolymers. The increase in tan ô peak height was accompanied by a decrease in the peak width, indicating a decrease in the number of environments for the motional units responsible for the glass t¡ansition. Minima in tan ô peak heights for intermediate copolymer compositions probably correspond to maximum polymer heterogeneity with a large spread in the polymer motions giving rise to the glass transition. Copolymers of TET2000 and 2900 contrasted with those for the

TET650 ând 1000 in that two loss peaks were usually found in the tan ô

-temperature plots, with the peak near -5OoC remaining unchanged in position throughout the range of copolymer compositions. Some TET 2900 copolymers showed increases in storage modulus through the DMTA scans after -40"C which were characteristic of recrystallisation and subsequent melting of PTMO crystallites. While this may have been expected for methyl acrylate copolymers, 204 recrystallisation of PTMO units in copolymers containing HDDA and TEGDA, where crosslink density was increased even at the lowest levels of comonomer incorporation, was not. A possible rationale for this may be that, while the total number of crosslinks was undoubtedly increased on addition of TEGDA or HDDA, the number of crosslinks with two bulky TDI groups held in close proximity to each other by the linking of adjacent acrylate groups was decreased. The position of the higher temperature tan ô peak due to the glassy comonomer varied in a non-linear fashion with copolymer composition. Predictions of the composition of the polyurethane and "plastic" phases by applying the Fox equation to the low and high temperature peak temperatures yielded contradictory results. Phase compositions inferred from peak positions indicated that TEGDA showed increased miscibility with the polyurethane phase in TET2900 copolymers, compared to

TET2000 copolymers. The hydrophobicity of PTMO chains would be expected to increase with increased PTMO molecular weight - greater miscibility of relatively hydrophilic TEGDA with TET2900 is altogether unlikely. Tan õ peak heights for the soft segment phase in TPCL and

TET(PTMO)2 copolymers were retained to a much greater extent in the urethane acrylates with higher molecular weight soft segments. Tan õ

(Tg) dropped to 0.23 from the homopolymer value of 0.43 in the TET( 1000)2 20wtVoTEGDA copolymer, with a significantly smaller proportionate decrease found for the TET(2000)2 20wtVo TEGDA copolymer : from 0.53 in the TET(2000)2 homopolymer to 0.35.

The extent of plasticisation of the higher temperature peak appears to be related, not to the miscibility of the copolymer with the urethane soft phase, but to the volume fraction of soft phase available to plasticise the reactive diluent transition. Transmission Electron Microscopy (TEM) 205 could possibly be used to correlate polymer morphology with copolym". dynamic mechanical response. DMTA runs on saturated urethane acrylate polymers (Section 3.2.3) showed three distinct types of dynamic mechanical response compared to the dry polymer. For the TET650, TPCL |ZSO and TET(650)2 polymers decreases in Tg consistent with water plasticisation were seen, with the extent of Tg depression being in agreement with that calculated from the Fox equation, where the Tg of water was taken to be 134K. A further group of polymers : TET1000, TPCL 2000, TPCL 3000 and TET(1000)2 showed little change in Tg, with slight increases in tan ô (Tg) for the saturated polymer. A final group : TET 2000,.TET2900 and TET(2000)2 showed increases in Tg, together with large decreases in peak height, which were attributed to ice formation in the polymers. The observed changes in polymer dynamic mechanical properties when equilibrated in water may be due to water residing in different environments within the network.

Water in pHEMA networks has been classified I as being freezing or non-freezing - recent work 2 further subdivides the non-freezing water into water strongly hydrogen bonded to the polymer, together with a secondary water hydration shell around this hydrogen bonded water. The bulk water, capable of freezing, is considered to be accommodated in polymer free volume.

While the water uptake in the urethane acrylates is low (1.5 - 3.lVo EV/C) compared to pHEMA networks (37-39Vo) similar distributions of water may occur, with water directly bonded to the urethane NH group and associated with carbonyl groups being in the non -freez\ng water category.

Freezing water probably forms more extens ively in networks where the hydrophobicity of the polymer chains promotes water 206 clustering: - TPCL 3000 containing PCL soft segments is unaffected by water sorption while the amorphous glass transition of saturated TET 2900 increased by nearly 20oC, with a concomirant decrease in tan ô(Tg).

NIR spectral results (Section 7.2.2) showed distinct differences in the water peak position for saturated polymers between PCL and PTMO soft segments, with the water peak in TPCL polymers being closer to that found in wet epoxy resins, where extensive water - polymer hydrogen bonding is known to occur.

DSC experiments (Chapter Four) failed to detect the presence of a hard phase, either directly through the observation of a high temperature glass transition, or indirectly from the observation of any increases in the soft segment Tg caused by phase mixing produced by high temperature annealing followed by rapid quenching to below the soft segment Tg. Minor peaks attributable to hard segment crystallite melting were seen in DSC thermograms for samples with very low crosslink den s ity. Non-Fickian water sorption kinetics, attributable to water clustering, was most pronounced in the TET2000 and 2900 polymers at

337o r-h.. It seems likely that similar clustering would also occur in the TET650 and 1000 networks, but two - stage sorption was probably not observed for these samples as a result of the greater crosslink density reducing the rate of the diffusion component of sorption relative to the relaxation component resulting from water clustering. Two-stage behaviour was absent in all desorption runs at ambient temperature in dry atmospheres, however, all the sorbed water taken up at various humidities was generally not removed. The discrepancy between water taken up and water lost on drying may be due to residual water directly hydrogen bonded to polar groups in the polymer. 201

Water uptake for all PCL based networks was less than that for polymers with PTMO soft segments. Copolymer water contents were generally less than those calculated from simple additivity of component EWCs, with greater deviations from calculated values being found in networks containing a higher weight fraction of urethane gròups. Vy'ater uptake in the urethane acrylate polymers and copolymers may be governed by the accessibility of the NH groups - if orher groups in the polymer can hydrogen bond to these groups overall water uprake may be reduced.

PEMAS 13C Ntr¡R time constants for PTMO carbons in the TET650 -2900 homopolymers showed increases in Tlp(C) and T5¡ as the PTMO chain lengthened, indicating increases in both mid-kHz and low frequency components of motion. Time constants for the PTMO carbons showed little change on copolymerisation with methyl acrylate and TEGDA - this correlated well with the observed constancy in the position of the PTMO Tg peak in the copolymers. Biexponential decay behaviour was only found for the TET2900 polymer 3 months after casting and may be attributable to the development of limited crystallinity in the polymer. The NMR data alone cannot be used to reject the presence of a hard phase, however, if biexponential decay had been found for the PTMO carbons in TET 650 and 1000 phase mixing could have been inferred.

S tructure/property relations in the urethane acrylates have been previously described in terms similar to those applied ro linear polyurethanes i. e. in terms of phase segregation and hard domain organisation. Taken together, the DMTA, DSC and NMR data do not support the existence of a hard phase in the urethane acrylate polymers

209

The differences in properties for TET and TPCL urethane acrylates may be important in their selection for various uses; e.g. in optic fibres where urethane acrylates are commonly used to provide the inner (buffering) layer in two-layer coatings. The rubbery polymer serves to dissipate stresses in the glass fibre which may lead to signal att€nuation. In moist environments where it is important to maintain the buffering properties of the polymer over a wide temperature range, clearly ice formation at subambient temperatures would be undesirable. Under such conditions, TPCL urethane acrylates would be the coating of choice, despite the greater hydrophobicity of PTMO soft segments. 2to REFERBNCES

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