>-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
Statement
Acknowledgements
Abbreviations
Chapter One Introduction
Chapter Two Experimental Techniques2.1 Sample sources
2.2 Sample preparation
2.3 Sample synthesis
2.4 Sample polymerisation
2.5 Dynamic Mechanical Thermal Analysis
2.6 Nuclear Magnetic Resonance
2.7 Dffferential Scanning Calorimetry
2.8 Sorption/Desorption Measurements
2.9 FTIR/NIR Spectroscopy
2.10 Thermogravimetric Analysis
Chapter Three Dynamic Mechanical Measurements
3.1 Introduction
3.2 Results and discussion
3.2.L Homologous series
(a) TET homopolymers
(b) TPCL homopolymers
vl
vt tt
tx
x
1
7
7
7
Il2
l2
l4
20
22
23
23
24
24
25
25
25
34
ll
(c) TET(Fr|MO)2 homopolymers
(d) Hard segment model compounds
3.2.2 Copolymers
(a) TET copolymers
(b) TPCL copolymers
(c) TET(F fMO)2 copolymers
3.2.3 Water in Urethane Acrylate Networks
3.3 Summary
Chapter Four DSC Results
4.1 Introduction
4.2 Results and discussion
4.2.I Homopolymer Thermograms
(a) TET homopolymers
(b) TPCL homopolymers
(c) TET(Fr[MO)2 homopolymers
(d) Hard segment model polymers
4.2.2 Annealing experiments
4.3 Summary
Chapter Five NMR
5.1 Introduction
5.2 Results and discussion
5.2.1 l3C liquid phase NMR
5.2.2 SOIid StAtC PEMAS 13C NMR
(a)TET homopolymers
(b) TET2g0O/methyl acrylate copolymers
39
43
53
53
7L
73
79
93
t20
120
120
120
t20
L20
r30
95
95
96
96
96
L02
104
108
tr7
118
ut
(c) TET29O0/TEGDA copolymers
(d ) Vãialion in TEI2ffi ad 290 NMR pææræs wih tirre
5.3 Summary
Chapter Six So¡ption and Diffusion6.1 Introduction
6.2 Diffusion kinetics
6.3 Results
6.3.1 Homologous series
(a) Sorption at 33Vor.h.
(b) Sorptio n at 7 9Vor.h.
(c) Sorption at 987or.h.
(d) Sorption in water.
6.3.2 Variation of diffusion coefficienb with water uphke
6.3.2 Copolymer sorption in water
(a) TET 650 copolymers
(b) TET 1000 copolymers
(c) TET 2000 copolymers
(d) TET2900 copolymers
6.4 Discussion
6.5 Summary
Chapter Seven FTIRNIR Measu¡ements7.1 Introduction
7.2 Results and discussion
7.2.1 ATR results
(a) Hydrogen bonding
133
136
140
142
142
142
145
t45
145
1s1
155
159
L64
L66
L66
L70
172
t75
177
187
188
188
189
189
189
IV
7.3 Summary
Chapter Eight
References
(b) Conversron
(c) Copolymers
7.2.2 NIR results
(a) Conversion
(b) Water uptake
Conclusions
L92
t94
196
t96
198
200
20t
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
BD
acp
AH*
ÂH.
DMTA
DSC
ED
EWC
FTIR
HDDA
HEA
HEMA
MA
MDI
NIR
NMR
PCL
PEMAS
PTMO
TDI
TEGDA
Tob
TET
ABBREVIATIONS
buta¡re diol
heat capacity change at Tg
DSC peak melting enthalpy
DSC peak crystallisation enthalpy
Dynamic Mechanical Thermal Analyser
differential scanning calorimetry
ethylene diamine
equilibrium water content
Fourie¡ Transform Infra¡ed S pectroscopy
1,6 hexanediol diacrylate
2-hydroxy ethyl acrylate
2-hydroxy ethyl methacrylate
methyl acrylate
4,4'-diphenylmethane diisocyanate
Near Infra¡ed Specroscopy
nuclea¡ magnetic resonance
poly(caprolactone) diol
proton enhanced magic-angle spinning
poly(tetramethylene oxide) diol
toluene diisocyanate (2,4 isomer unless otherwise noted)
tetra(ethylene glycol) diacrylate
glass transition temperature
urethane acrylate with one F fMO chain between two terminal TDVggA
unlts
X
TET(PTMO)2
TPCL
urethane acrylate with two mMO chains connected by a single TDI group,
between two terminal TDVFIEA units
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
I1HO
Figure 1.1 Typical linear polyurethane strucrure.
à)sof t seçm¡n15/\,/\
H
c((cH2)
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 urethaneacrylate copolymer networks 39 attributed modulus changes oncopolymerisation to improved phase separation brought about bypreferential 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
l
2.3 Sample synthesis(i) TET and TPCL prepolymers
The synthesis scheme for these prepolymers is shown in Figure2.L. HEA was added dropwise to an equimolar amount of TDI withstirring 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 wtVostannous 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
a further 0. 15 wtTo stannous octoate. The reaction was heated for afurther 2 hours at 70oC with reacrion completion being verified by IR
spectroscopy from the absence of the -NCO stretch at 2273 cm-1 in theprepolymer.
liiil Hsrd seørnent rnrìdel r.rì mnnrrnrlc
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 moleof 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
(i)=C=O
=C=O
+
o
(ii) 2r
oil-o-cH2cH2-o-c-cH=cH230"c
+
on
NH-C-O-
IilHO-CH2CH2-O-c-cH=cHz
HO-((CHt4O),¡-H pTMO diol
oHO-
il(c-(cHtso)n -H PCL diol
0. I 57o stannous octoate
70'c
olt
cH2cH2-o-C-CH=CH2
or
Hrcolt
or
NH- C - O -((CHt4O)n -TDr-HEAoil
(c-(cHr5o),
Figure 2.1 Reaction scheme for TET and TPCL urethane acrylate prepolymersynthesis. N refers to the numhr of repeat units in the soft link.
10
(i) 2 Ho-((cH2)4o)n-H +
o
Ho-((cH2)4o1"-[ *
N=C=O
N=C=O
0.15% stannous octoate700c
o-((cHr4o)n-H
oilc-
II
(ii) u+
2l
o
HEA-rDI-O -(cHzlqo)- [ -¡¡',r70"c
0.157o stannous.octoate
NH
ö-oI
o(cHr4o),TDr-HEA
Figure. 2.2 Reaction scheme for TET(PTMO)2 urethane acrylate prepolymersynthesis.
11
CH3
NH-
NH_
{{
9?c-o.cH2cH2-o-c-cH{H2
ooll il
NH-C-O -CH2CH2-O -C-CH=CH2
9?c-o-cH2cH2-o-c-cH{H2
I
III
Figure 2.3 Ha¡d segment model compound prepolymers : (I) 1: I mole rario TDIÆ{EAand (IIÐ 1:2 mole ratio TDIÆ{EA.
T2
Synthesis of the l:2 mole ratio TDI/FIEA hard segment modelcompound involved addition of a further 1 molar equivalent of dried HEA
containing 0.15 wtVo stannous octoate catalyst to the 1:1 mole ratiocompound. 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 thephotoinitiator. 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 VibratorD isplacementTransducer
:,
TemperaiureE nclosure
Sample
tii
Liqu idN itrogen
Drive Shaft Clamps
(a)
(b)
Figure 2.a @) Schematic diagram of the Polymer t-abs DMTA head used in thebending 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.
The DMTA Analyser solves the relevant equations of motion 3'4 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
CDC13 triplet at't1.0ppm. The spectrometer operated at 20.1MHz for l3C
observation. Standard techniques were used for measurement [ 13C 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
followed by a 90o phase shift and continuous application of a strong lH
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
components of motion, can then be calculated from the variation in l3C
17
90"Til2
lH
decouple
13c
acqulre recycle
Figure 2.5 (a) Pulse progranìme for obtaining spin-lock, cross polarised NMR timeconstants, T5¡.
900
1H
13c
hold
Figure 2.5 (b) Pulse programme for obtaining relaxation time constants for I3C in the
rotating frame, Tlp(C).
lockÞpln
tcp
900tú2
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 Tslr1p (H)
(2.2)rl p (C)
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 beingrl p (C)
insignificant and Equation 2. 1 reduces to
I=+rexprffi;l - expclfril (Tsl r lp(H)I I ) (2.3)
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
2É.
g1-
é
oJe1-
o
l.lcltrng tndothcrm
t
Tm
Tcmpcrarurc (qc)
Figure 2.6 (a) The peak onset method used for determining melting and crysrallisariontemperatures from DSC thermograms.
-,,
ce
Lcpr ?
T
)-ro
Figure 2.6 (b) The midpoint method used for determining glass transitlontemperatures from DSC thermograms. T, is taken as being f f þCp) i.e. thetempemture at which tl¡e heat capacity change is half of that of the total change from theglassy to the rubbery state.
20
alternation and received phase cycling (CYCLOPS ) to remove quadimages. 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 astemperature 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
-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 tocalculate Tg for samples with small ACo.
Oneet
cpDo ItoT9
TI
f2'c'c'cJ/g-C'c
-72_ sJ31
-20. o00
-52.5?1
0. 353
-58.8t9
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
the V/C determined when the
particular wet environment.
Equilibrium water content (EWC) was
sample had reached equilibrium in a
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 MBCHANICALMEASUREMBNTS
3.1 IntroductionThe 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 Discussion3.2.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
ct
0.)!Ect
0.3
0.2
0.5
0.4
0.1
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 softsegments of indicated molecular weighr
650
o 1000
2000
tr 2900oo
I!
I¡t
rh6+a!
o %I
o
oa
a
t ho oo
%b
qbo
ho. % oo
O
!to
ocflro.'tlô
aoo
%gq
!o
O
oo
oO
o
o
os
t¡ E!
oO
oo
a
d
o
I
o
+
a
f
aa
aa
aa
ot
a
oEI
s
E
a
EI.o o'ì\*o-
27
rrèoo
9.5
8.5
7.s
6.5
5-5
-r25 -100 -15 -50 -25 0 25 50 75 100 r25
Temperature (oC)
s 650
o 1000
o 2000
+ 2900
*t++
g
trIIeeo+
oo
orD
++
+
Iofo
o+ Ea+o
EIOo
oaI o
oI oar (,
ogOr o
oI
ú
*1
\
\
aa
\
Figure 3.1^b t og E' - temperature plots for TET u¡ethane acrylates with pTMO softsegments of indicated molecular weighr
289
8
=t¡lèoo
't
4
6
a
5
.r25 -100 -75 -50 -25 0 25 50 15 100 r25
Temperanue (oC)
4,
E 650
o 1000
! 2000
+ 2900
E+o o
oooo
+!+F
E++
u
o(,o
oo
rD
oo
E
EI+I|lI
1troo
ootrOq
ao¡E
oootrona
Eo
#oEl++ os-q'
Figure 3.l^c t og E" - temperature plots for TET urethane acrylates with pTMO softsegments 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 3.1Dynamic 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.)
* 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.
0.Ll-65131*3076.7 |-502900
0.236588*3026.85-502000
0.37-M823297.38+131000
0.48-1554.4627.95+356s0
wtrTg(E")wp('c)Peak heightLogE'(N/m2)
(25.C)Tg('c)PTMO MWt.
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 Bpeak 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 oxyethylenerepeat 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 tanô(Tg) from TET650 to TET2000 may reflect restrictions on chain motion
which occur as a prelude to the development of more extensivecrystallinity 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 considerabletemperature difference between the Tg peak in the tan õ and the log E" -
temperature plots, with the maximum in the log E" temperature plotsbeing 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 lossmodulus and tan ô peak positions in a series of copolymers of maleic
anhydride crosslinked polystyrene (p(ScoMA)) with linearpolyvinylmethylether (PVME). Samples containing small amounts ofeither 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 lossmaximum 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
GI
q)-é
ñ
0.3
0.2
0.5
0.4
0.1
0.0
-100 -75 -50 -25 0 25 50 15 100
Temperature ('C)
tr 1250
o 2000
o 3000
Elo
oE¡
oE
aooo o
o
o
ao
(to
(t
rt
do
a
o
ao
o
o
o
ta t"
ooog
E
oo
trE
atr
oO
E¡a
tr
o
troo o
oo
O
oæ
ooo o
oao
oo
oE-ooo
oo
Oa
&'O¡
"4otro
tr
E
trE
oo
oo
o
oo
o
Figure 3.2 a Tan ô - temperature plots for TPCL urethane acrylates with PCL softsegments of indicaæd molecula¡ weighr
36
9
l0
8
6
tièoo
7
-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 softsegments of indicated molecular weight.
o 1250
. 2000
o 3000
ael oo¡oo o -ereo ot
oEI
O¡
o
o
Oi,oo
osotr
OOaoa
0ooo
ooo
oo
oooo
37
9
8
7
f¡ìõoo
6
5
4
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 softsegments of indicatirC molecular weight.
o 1250
o 2000
o 3000
ooooa
oo
E¡oE
EIEI
%o(lo
EoE
E
t gEtr
o
sooo
EE
o
atrE
o
ûo
ooa
a
tE.. tr
ooo
oo
Etr
tr
EE
tr
ooo Qcoo
oo
Oo
38
TABLE 3.2Dynamic 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
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.
-373000 -453E0.38707.6300ó.EE2I-262000 -4054o.4045
rz50 1.2027+8 -27550.3711
'l'g("c)PCLM.V/t
LogE'(N/mz¡çZS"C) Tg(E")'Wp("C)Peak height
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 3.3Dynamic 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 glasstransition peak.
-570.1 l36.3s9-41t2¿9VU-63320.53436.446-542¿t )vI-46480.42156.831-'t2t21000I-32630.49876.ó68-6
Tg(8")wtp("c)Peakheight
LogE'(N/m\25"ClsCC)FTTMOlv1Wt
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
clq,
cÉ
0.6
0.5
0.4
0.3
0.1
0.0
o.2
-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 withFrIMO soft segmens of indicated molecular weight.
EI (6s0)2
+ (1000)2
a (2000)2
o (2e00)2E+
o+
**+ s
EI
+a
a
t?+
++
a
+
+
a
+o +
so
a
oa
GI
++ +++
¡**++
Eb
tr
o
tat +
+
a
+a
ooo
o
+
****EtsEI
otr
oPFtEttr
EIE
.ofr.ao
E
EIE
Et
oE
str
tr
E
EgE
tro
EI
++ç+*+
4r
9.0
8.0
7.0
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 withFTMO soft segmens of indicated molecular weight.
l¡lÞoo
E (6s0)2
o (1000)2
o (2000)2
 (29O0)2
aoeoooræo'
Etr
É
Â.-D-_.
o l-lto
.U'
ooo
o
a
OEI
O¡
a
o
aaaa
ao
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ooa
a
a
ao
OOo
ooa
E 'âEtoEI
o
o
a
42
:Hè0o
9
8
7
6
5
4
-100 -15 -50 -25 0 25 50 15 100
Temperanne ('C)
(6s0)2
+ (1000)2
a (2000)2
 (2e00)2
** 8E+tr
t+k
rÐc
aa
a
E
+
+
+aoa
oGI
+ EGI
oI
++
++
+++
IEgt 3.3 c Log E" --.temperature plots for TET(PTMO)2 urethane acrylares withPTMO 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
-50 -25 0 25 50 75 100 r25 150 r75 2NTemperature ('C)
Figure 3.4 Tan ô - temperature plots for linear (1:l mole ratio) and crosslinked (1:2mole ratio) TDI/FIEA hard segment model polymers.
-30 -5 20 45 10 95 t20 r45 170 195 220Temperature (oC)
Figure 3.5 Tan ô - temperature plot for linear (1:1 mole ratio) TDVHEA hard segmentmodel polymer scanned at a slower rate(1K/min) than in Figure 3.4.(3lVmin).
cÉ
q)õGI
0.5
0.4
0.3
0:2
0.1
0.0
cË
q)!
cl
0.6
0.5
0.4
0.3
0.2
0.1
0.0
o 1:1o l:2
acb.
" ¡o\'sdÐ ho*l@
oO
E
tr
Oa
troo
Ea
otr
oEO
trtrE¡EEtrEoE
45
9.5
8.5
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 hardsegment model polymer.
320
300
280
260
240
220
200
180
0.1 0.2 0.3
wh
0.4 0.5
Figure 3.7 Variation in experimental and calculated Tgs with hard segment weightfrattion for TET and TET(PTMO)2 polymers. Tgs were calculated using the Foxequation.
l¡lèoo
Tgexp
Tgcalc
E¡
O
v(u
GIl-(uÀEq)
ooo
o
oE¡
Eo
E
tr
o
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 :
ITo-õ
wlTgl
w2Te2+
(3.2)
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 thepolymer 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¡¡lwhere 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
t
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 totalnumber of crosslinks and junctions, X., is given by :
fXc =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
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
Xc = 0'a79 Z
ã4=
Ts
600.5 1t254.36
2oo[1
= 0.479
fi=2 (i-2)/ 2
+
for i=4 (tetrafunctional crosslinker)
Xc = O.479
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.
350
300
5l
o Tgexp+ Tgcalc
v(¡)¡i
cɡr(uq
q)
250
2000.0 0.1 0.2 0.3 0.4 0.5
Crosslink density (Xc)
Figure 3.8 Variation in experimental and calculated Tgs versus network crosslinkdeñsity for TET and TET(PTMO)2 polymers. Tgs were calculated using DeBenedetto'sequation.
+o
c
I
a
o
+
o
+
to
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 3.4Dynamic mechanical results for TET650 ursthane acrylate copolymers. The Tg (glasstransition temperature) was determined from the tan &temperature plot. Tg@") refers tothe highest temperature peak in the log E"-temperature plot. Peak height refers to the tanõ glass transition peak.
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
+5784.ó50 I0.2258.7700.3078.512+5058.020 -z
HDDA+44ðu.55U rJ.3248.0ó5 +9
8.353+4350.820 00.356TEGDA
1.195+2793.550 +7t.234+2918.4zt) +I3o.6617.801
methylacrylate
'lg(8")Peakheight
refc)I|gdoIVodiluent
WtTodiluent
54
1.5
1.0
0.0-100 -75 -50 -25 0 25
Temperature (oC)50 75 100
Figure 3.9 Tan ô - temperature plots for TET 650 copolymers containing the indicatedwt%o methyl acrylate.
0.3
0.2
0.1
0.0
cq
{)õct
0.5
0.4
c€
q)õC!
100 -15 -50 -25 0 25Temperature (oC)
50 75 100
E
o
20
s0
rD
EO
Ett
o
a
o
o
trEo
E¡ gtr
o
qoE
O TEGDA. HDDA
s o s tr ".".r.råTf.t l".".tåt t o o o o'r
E o6 oEEoO
oEo
oooO
o
GI
E
o oooo to
Figure 3.10 a Tan ô - temperature plots for TEGDA and HDDA homopolymers.
55
frlè0o
9
8
l0
7
Figure 3.10 b Log E' - temperature plots for TEGDA and HDDA homopolymers
0.3
0.1
0.0
0.4
lm -15 -50 -25 0 25Temperature (oC)
-25 0 2sTemperanre (oC)
50 75 100
0.2
ctl
o)'tt
GI
-100 -15 -50 50 75 r00
S TEGDAO HDDA
Egg EggE
oooooooaoo
EEtr ry
E
cooote
tr¡t
E
O
20
50
a o
o
O
cf
tro
Eo
EIo
oOg
ao
O
Figure 3.11Tan õ - temperature plots for TET 650 copolymers containing the indicatedwtTo 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 3.5Dynamic 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 ôglass transition peak.
+8l6E I8.376+528850-I17888.102+538340-6.19848.218+437530
-18.230t7.938+356420-zr.27081.664+2344l0
HDDA+9.26388.295+4I8450-242t597.80ó+391840-17.258t8.027+367030-2126627.859+285720-33.29587.602+253710
TEGDA+41.00886.878+ZtJ9550+l.79426.933+IÓ9340-2.62747.160+178930-2.48357.214+21E2ztJ-17.39227.278+l'l6810
methylacrylate
Tg(8")Peakheight
l-ogE'(N/m¿)l¿5"c)'IsCC)lv1ol7odiluent
WtTodiluent
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
50 50 75 100
Figure 3.13 Tan ô - temperature plots for TET1000 copolymers containing theindicated wt% methyl acrylate.
GI
q)õcl
0.4
0.3
0.2
0.1
0.0
-100 -75 -50 -25 0 25 50 75 100Temperanue (oC)
Figure 3.12 Tan ô - temperature plots for TET 650 copolymers containing the indicatedWt% HDDA.
CI
q)õcÉ
t.2
1.0
0.8
0.6
0.4
0.2
0.015100 -25 0 25
Temperature ('C)
o20o50
oa
gE'tr
EoEOtr%
EotrOEoE
.10
.20o30.40.50
þta. .
Q+a
aa.
a
a
aa
.Oa+
o
Itrf
ra
ot
.oa
aaaa
.loo(¡
;.{
roa
o
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
0.3
0.1
0.0
lm -75 -50 50 75 100
Figure 3.14 Tan ô - temperature plots for TET 1000 copolymers containing theindicated wtTo TEGDA.
0.3
0.2
0.1
0.0
-25 0 25Temperaure (oC)
Figure 3.15 Tan ô - temperature plots for TET 1000 copolymers containing theindicated wtTo HDDA.
0.2
cll
q,)õcl
-25 0 25Temperature ("C)
ctq)õCE
-100 -75 -50 50 75 100
E¡ 10+20o30a40o50
El,¡'o
Ff*Jq3
o o o
+eAa+
aO+tr+EI
stros
tr10.20!30o40.50
a
õI'
ae
a Eo.otooa
o!OoE
ao
o
oa
¡EEoa
õa
Eoa
EoO
O+E
at""
!o
å.
.t t tf,.
EEo
EO
Ð
;%
o o.!
o o'E ro.
O Ia+E oio'Eo"tt
60
C!
q)õcg
0.8
0.6
0.4
0.2
0.0
Figure 3.16 Tan õ - temperature plots for TET2000 copolymers containing theindicated wt%o methyl acrylate.
0.3
0.2
0.1
0.0
GI
q)!
6
100 -15 -50
-100 -75 -50
-25 0 25Temperatue (oC)
-25 0 25Temperature CC)
50 75 100
50 15 100
Figure 3.17 Tan ô - temperature plots for TET 2000 copolymers containing theindicated wtTo TEGDA.
ol0.20E30o40.50
"".,.l".:.TfË,
"*iËii'.ãTi:'o' "
!9
¡!¡
o o'
o
an o
O!
to+ o
EE a
E
oo
Es o
"EtroEEooErEtraaaaoooooo
iïiiii::::
Er 10.20r30o40.50
E}¡
E Þaa
"iþ
o
!.o
OE
61
TABLE 3.6Dynamic mechanical results for TET2000 urethane acrylate copolymers. The Tg (glass
transition temperature) was determined f¡om the tan &temperature plot. Tg@") refers tothe highest temperature peak in the log E"-temperature plot- Peak height refers to the tan ôglass transition peak.
-59r3048. IÓÓ-50,+599Z50-59.12838.018-50,+538940-59r37 57.90r-50,+248330-59.16037.596-50,+107420-5919117.189-50,-I256l0
HDDA-61.066I,.20198.169-56,+459050-620948,.19777.925-51,+428540-63.rt70,.t6251.715-57,+307930-ól.15'39,.t727.428-52,+ZI68zrJ-58.L937,.1827.157-48,-lU49IO
TEGDA+l.73386.135+1ó9750-58.54096.688+149540-54.40166.74tJ+79330-59.30836.723I88ZtJ-58.25'376.6t:2-387710
:l'g(ts")Peak herghtLngE'(N/m¿)l¿J"v)'lg("c)MoLTodiluent
WtTodiluent
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 weightpercentages. 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
c!
q.)
E
c!
0.2
0.1
0.0-100 -75 -50 -25 0 25
Temperarure (oC)
50 75 100
Figure 3.18 Tan ô - temperature plots for TET 2000 copolymers containing theindicated wtTo HDDA.
0.6
0.2
0.0
-100 -15 -50 -25 o 25 50 75 100Temperature (oC)
Figure 3.19 a Tan ô - remperature plots for TET2900 copolymers containing theindicated wtTo methyl acrylate.
0.8
0.4
GI
q)!
ñ
o10.20E30o40.50
otrtrotretrtr"t""tr "t
E ^...tr ooot_I{.^
_... ._:.,.',,:.
"":::id, ir.ri,
t _ooo ooo ..*t o o ¡"::;
