A systematic study of the effect of the hard end-group composition on the microphase separation, thermal and mechanical properties of supramolecular polyurethanes

Article abstract of DOI:10.1016/j.polymer.2016.07.029

This paper reports a systematic study on a series of supramolecular polyurethanes that possess microphase separated morphologies which afford elastic materials at room temperature. Combinations of urea and/or urethane linkers in addition to a phenyl space

Full text of DOI:10.1016/j.polymer.2016.07.029

Polymer xxx (2016) 1e11
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Polymer
journal homepage: www.elsevier.com/locate/polymer
A systematic study of the effect of the hard end-group composition on
the microphase separation, thermal and mechanical properties of
supramolecular polyurethanes
Daniel Hermida Merino a, Antonio Feula a, Kelly Melia a, Andrew T. Slark b,
Ioannis Giannakopoulos c, Clive R. Siviour c, C. Paul Buckley c, Barnaby W. Greenland d,
Dan Liu e, Yu Gan a, Peter J. Harris a, Ann M. Chippindale a, Ian W. Hamley a,
,
*
Wayne Hayes a
a Department of Chemistry, University of Reading, Whiteknights, Reading, RG6 6AD, UK
b Henkel UK Limited, 957 Buckingham Avenue, Slough, SL1 4NL, UK
c Department of Engineering Science, Oxford University, Parks Road, Oxford, OX1 3PJ, UK
d The Reading School of Pharmacy, University of Reading, Whiteknights, Reading, RG6 6AD, UK
e Department of Physics, University of Surrey, Guildford, Surrey, GU2 7XH, UK
a r t i c l e i n f o
a b s t r a c t
Article history:
This paper reports a systematic study on a series of supramolecular polyurethanes that possess micro-
phase separated morphologies which afford elastic materials at room temperature. Combinations of urea
and/or urethane linkers in addition to a phenyl spacer have been used to study the effect of the rigidity of
the hard end group segments as well as the hydrogen bonding capability of the urethane-urea linker
units. Small angle X-ray scattering (SAXS) experiments have revealed characteristic microphase sepa-
rated morphologies. Wide angle X-ray scattering (WAXS) was used to probe the lateral packing of the
urethane and/or urea within the hard segments. Differential scanning calorimetry (DSC) analysis
confirmed that unsymmetrical soft/hard segment phases have been achieved by varying the urethane/
urea content. Transmission electron microscopy (TEM) and atomic force microscopy (AFM) determined
that a 1-D fibrillar structure was obtained when the hard segment featured ureas whereas a 3-D
structure was achieved when a combination of urea and urethane groups was used, giving rise to
enhanced elongation properties. Finally, we present mechanical testing data in which oscillatory
rheology at a range of frequencies and temperatures has revealed the effect of the connectivity of the
hard segments on the relaxation times of the supramolecular chains. Tensile tests showed that end
groups with ureas or a combination of a urea and urethane yielded elastic materials with strengths of ca.
5 MPa at room temperature.
Received 11 May 2016
Received in revised form
7 July 2016
Accepted 10 July 2016
Available online xxx
Keywords:
Supramolecular polymer
Self-assembly
Polyurethane
Phase separation
Rheology
© 2016 Published by Elsevier Ltd.
1. Introduction
materials [6e8]. In addition, the mechanical properties of PUs are
intimately related to their chemical composition [9,10]. The struc-
Polyurethane and polyurea (PU) based materials [1,2] have
found widespread applications in the industrial scale production of
synthetic rubbers, adhesives, protective coatings, foams, fibers and
elastomers as well as in biomaterials [3,4] and semi-permeable
membranes [5]. The extensive use of PUs has been facilitated by
their ease of synthesis from relatively inexpensive starting
ture/property relationships that exist between the monomeric
components and bulk material provide wide scope for tailoring and
optimisation of the physical properties of the PUs to target specific
applications [11e15].
One of the most intriguing properties of PUs is their elasto-
meric behaviour [16,17], which arises from the segmented nature
of the chemical composition within each polymer chain [18,19]. In
the solid state, PUs are observed to phase-separate into soft do-
mains (containing flexible polymeric components) and hard do-
mains which contain a high proportion of urea and/or urethane
* Corresponding author.
E-mail address: w.c.hayes@reading.ac.uk (W. Hayes).
http://dx.doi.org/10.1016/j.polymer.2016.07.029
0032-3861/© 2016 Published by Elsevier Ltd.
Please cite this article in press as: D.H. Merino, et al., A systematic study of the effect of the hard end-group composition on the microphase
separation, thermal and mechanical properties of supramolecular polyurethanes, Polymer (2016), http://dx.doi.org/10.1016/
j.polymer.2016.07.029

2
D.H. Merino et al. / Polymer xxx (2016) 1e11
linkages [9,20,21]. The urea and urethane components are closely
associated through hydrogen bonding interactions [2,8,22,23],
thereby providing physical crosslinking between the soft seg-
ments. Thus these PUs behave similarly to conventional, covalently
crosslinked rubbers, except that disruption of the hydrogen
bonding by elevating the temperature, permits melt-processing
[24].
2. Results and discussion
2.1. Design
The aim of this study was to gain a deeper understanding of the
relationship between the structure of the end-groups of linear
supramolecular PU based materials and resulting morphological
and mechanical properties in the solid state. It was predicted that
an increase in the rigidity and the number of possible hydrogen
bonding interactions within the polymer end-group would
generate materials with advantageous rheological properties under
ambient conditions, whilst still permitting facile processing at
elevated temperatures.
€
Yilgor and co-workers have investigated extensively the
structure-property relationship in segmented PUs including the
local packing of the urea moieties, which results in a high degree of
microphase separation [16,25]. In addition, phase separation may
be promoted by using apolar soft segments with low surface en-
ergies such as polydimethyl siloxane (PDMS) [5,26e28]. During
extensive studies by Bouteiller and co-workers [29], PDMS with bis-
urea motifs were synthesised, the resulting materials exhibited a
3D network of hard domains, as a consequence of the partial
crystallisation of the hard segments. In a related study, Sijbesma
and co-workers have investigated [30] the influence of increasing
the length of the PDMS soft segments. They found that it was
possible to obtain fibrillar structures by increasing the in-
compatibility of the soft and hard segments [31]. Such fibrillar
structures were not observed in PUs that contained PEB [32] or PCL
[33]. An interesting spherical morphology was observed when
longer soft segment chains were used as a result of a transition of
cylindrical to micellar ordering along with the enlargement of the
volume fraction of the external segment, as predicted by Flory-
Huggins [34]. Detailed studies conducted by Thurn-Albrecht et al.
have established [35] that for a series of telechelic polyisobutylenes
equipped with complementary hydrogen-bonding groups, in-
teractions between micellar aggregates leads to network formation
and solid-like properties at lower temperatures induced by gelation
without any specific ordering. Recent studies have not just been
restricted to low molecular weight polymers with functionalised
chain ends, for example Creton et al. have shown that the visco-
elastic characteristics of a series poly(n-butyl)acrylates whose
central units feature pendant hydrogen bonding motifs can be
tailored by the loading of the self-complementary recognition
groups [36,37]. Meijer and co-workers have introduced mono-
disperse hard segments within PUs, and these hard segments were
found to undergo crystallisation, and thus achieve a higher degree
of microphase separation [38]. It was found that the polymers,
which featured two urea groups, displayed a useful balance be-
tween high toughness and low melting temperature, the latter of
which is key to aiding processability. In agreement with recent
studies by Rowan and co-workers [39e41], we have found [42e44]
that tough materials could be formed by addition of hydrogen
bonding end-groups with low binding constants (between 1 and
45 Mꢀ1) [45] to relatively low molecular weight PUs (between
10,000 and 20,000 g/mol). These materials possess interesting
temperature responsive properties whereby the degree of phase
separation within the materials, and the temperature at which the
polymers started to weaken, correlated to the binding constant of
the end-group [2,7,22,46,47].
2.2. Synthesis
Prior to commencing polymer synthesis, two novel polymer
end-groups (1 and 2) were synthesised via a two-step route by
initial addition of morpholine derivatives 3 and 4 to 4-nitrophenyl
isocyanate to produce urethane 5 and urea 6 followed by reduction
of the nitro groups (Scheme 1, see also the Supporting Information
(SI), Figs. S1e8).
With the end groups in hand, the targeted supramolecular
polymers were generated through a two-step protocol [8,42,43].
Briefly, this involves the synthesis of pre-polymer 7 by the addition
of MDI to hydroxyl-terminated poly(ethylene-co-butylene) (Krasol)
in the bulk state at 120 ꢁC (NCO:OH ratio of 2:1). Subsequent
addition of the endgroups (either 1, 2, 3 or 4) to 7 results in the
formation of supramolecular PUs 8, 9, 10 and 11, respectively
(Table 1). The supramolecular PUs were isolated in good yields
(>84%) after several precipitations into methanol.
All four of the supramolecular polymers exhibited comparable
molecular weights and polydispersity as measured by GPS analysis
(Mn 6600 g/mol 10% and Ð ¼ 1.8 0.1). This increase in molecular
weight compared to the starting Krasol material (Mn ¼ 1730 g/mol,
Ð ¼ 1.2) demonstrates that the conditions used to produce the
prepolymer 7 result in repeatable degrees of chain extension cen-
tred around 3 repeating units of Krasol/MDI.
The resulting polymers were structurally related, possessing
identical morpholine end-groups but varying numbers of urea
groups (see comparison of 8 and 9 or 10 and 11), with polymers 10
and 11 featuring two additional urea groups when compared to
polymers 8 and 9 (see Table 1). Maintaining the composition of the
hydrophobic block whilst changing the end group enabled the
study of the changes in phase separation on the mechanical prop-
erties to be investigated in isolation from changes incurred through
altering the soft segments of the PUs.
2.3. Solid state structures of end group mimics
In order to provide an insight into how end groups of this type
may aid the assembly of polyurethanes in the bulk, the end group
mimics 12 and 13 (Fig. 1) were generated (see Figs. S17-20) and
solid state structures obtained from single-crystal X-ray diffraction
analysis (see Fig. 1, Table S1 and Figs S21e22).
The dominant intermolecular interaction evident for urethane
12 was found to be a strong hydrogen bond (2.142 Å and 139.73ꢁ)
between the morpholine oxygen and urethane hydrogen (see
Fig. 2a) whereas bifurcated hydrogen bonds [48] between urea
moieties dominated in the solid state structure of urea 13 as shown
(see Fig. 2b).
Herein we describe the systematic design and synthesis of four
novel, low molecular weight (Mn z 6600 g/mol, Ð z 1.8) PUs
which vary in the nature of their end-groups. The effect that sys-
tematically changing chemical structure and relative % composition
of the hard end group has on the morphology of the new materials
has been studied using multiple techniques. In addition we have
investigated the mechanical properties of these novel materials by
oscillatory rheometry and tensile testing. These extensive studies
provide new insights into the structure property relationships that
underpin this important class of materials.
2.4. Morphology
Variable temperature SAXS has revealed that each member of
Please cite this article in press as: D.H. Merino, et al., A systematic study of the effect of the hard end-group composition on the microphase
separation, thermal and mechanical properties of supramolecular polyurethanes, Polymer (2016), http://dx.doi.org/10.1016/
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D.H. Merino et al. / Polymer xxx (2016) 1e11
3
Scheme 1. Synthesis of novel end-groups 1 and 2.
Table 1
Synthesis and selected properties of polymers 8e11.
Polymer
R
Yield %
88
Hard contenta
17.4
%
End-groupb
3.1
%
8
9
97
95
17.4
20.3
3.8
7.0
10
11
84
20.3
6.9
a
Calculated as the %wt/wt of structural components in the polymer that include MDI units as part of the hard segments.
Calculated as the %wt/wt of structural components in the polymer that comprise the terminal MDI units and R groups.
b
the series of supramolecular polymers (8e11) exhibited phase
Time resolved SAXS analysis was conducted on each polymer
(8e11), acquiring a scattering pattern every 30 s as the samples
were heated from ꢀ70 ꢁC to 70 ꢁC and then returned to ꢀ70 ꢁC
(heating and cooling rate 5 ꢁC/min) (see Fig. S25). Each sample
presented a single Bragg reflection in the SAXS patterns, typical of
the phase-separated morphology of PUs. Qualitative analysis of the
SAXS patterns for the polymers 8 to 11 sequentially revealed that on
increasing temperature, the intensity increases slightly up to 0 ꢁC
for samples 8 and 9 and then decreases markedly. In contrast, for
samples 10 and 11 the intensity increases continuously up to at
least 50 ꢁC. Comparing the SAXS data for 8 and 9 which contain the
more weakly associating end groups with 10 and 11 (which contain
an increased number of strongly hydrogen bonding urea residues)
shows a clear increase in the phase contrast as the hydrogen
bonding potential of the end-groups increases.
separation (see Figs. S23 and 24) [45,49,50]. The Bragg spacing
associated with the phase separation was found to be 5.18 nm,
5.71 nm, 7.12 nm and 7.19 nm for 8e11, respectively. It can be seen
that a small increase in the hard group content from 17 to 20% (from
8e9 to 10e11, respectively) resulted in a dramatic increase in
d spacing (at least 1.5 nm). For each pair of polymers with the same
hard-group content (8e9 and 10e11), the substitution of a more
strongly hydrogen bonding urea group in place of a urethane
resulted in a subtle increase in spacing.
For the polymers with the lower hard block content (i.e. 8 and
9), the intensity of the primary peak decreased with increasing
temperature, above 0 ꢁC (see Fig. S23 and Fig. S24), suggesting an
Fig. 1. End group mimics 12 and 13.
Please cite this article in press as: D.H. Merino, et al., A systematic study of the effect of the hard end-group composition on the microphase
separation, thermal and mechanical properties of supramolecular polyurethanes, Polymer (2016), http://dx.doi.org/10.1016/
j.polymer.2016.07.029

4
D.H. Merino et al. / Polymer xxx (2016) 1e11
Fig. 2. Single-crystal X-ray diffraction analysis of a) urethane 12 and b) urea 13.
increase in miscibility between the end-groups and the mid blocks
as the temperature rises. Conversely, for polymers 10 and 11, the
domain spacing increased slightly with temperature. As observed
previously [29], this may be accounted for by stronger hydrogen
bonding interactions within the larger end groups, maintaining the
integrity of the hard blocks over the temperature range studied.
Thus, the hard domains remain intact whilst the soft domains un-
dergo thermal expansion, increasing the domain spacing at higher
temperatures. It is also not unreasonable to expect the hydrogen
bonding distances between the hard blocks to expand as the
temperature of the environment increases. This trend is consistent
with the observation that the temperature at which each polymer
exhibits the maximum in its domain spacing increases in the order
8e11, (from 0 ꢁC for 8 to > 70 ꢁC for 11) indicating that disruption of
the hard segments occurs at higher temperatures. The temperature
dependence of this molecular-scale packing is observed readily in a
plot of the intensity of the Bragg peak as a function of temperature.
It can be seen that for all polymers (8e11) the d spacing increases
with temperature, consistent with the thermally induced dissoci-
ation of the end-groups (Fig. 3). The WAXS data in Fig. 3 shows a
clear increase in the d spacing with temperature for all samples (the
full WAXS profiles are shown in Fig. S26). This distance reflects
4.92 Å
40000
35000
30000
25000
20000
15000
10000
5000
9
4.59 Å
4.22 Å
11
I
8
3.47 Å
10
4.98 Å
0
10
20
30
40
2θ / degrees
Fig. 4. WAXS scattering patterns of 8e11 at 25 ꢁC.
side-to-side packing of molecules.
Comparison of the WAXS scattering signals for each of the
polymers at 25 ꢁC revealed more information about the packing
within the hard domains (Fig. 4). Polymers 8 and 9, with the lower
hard block content give a broad scattering signal with maxima at
4.98 nm and 4.92 nm, respectively, consistent with an amorphous
system. In contrast, in the case of polymers 10 and 11, three sharp
signals are evident in their WAXS patterns at ca. 4.6, 4.2 and 3.5 Å.
These well-resolved reflections demonstrate the presence of re-
gions with a certain degree of ordering, which may be attributed to
5.10
5.05
8
5.00
4.95
9
4.90
4.85
4.80
4.75
4.70
4.65
4.60
4.55
4.50
4.45
4.40
4.35
4.30
4.25
10
11
urea and urethane lateral spacing (4.6 and 4.2 Å) and
(3.5 Å) [51] as a consequence of the aromatic spacer unit within
these end-groups.
p p stacking
ꢀ
Analysis of the DSC thermograms for each of the polymers
(8e11) supports the X-ray scattering analysis and indicates the
amorphous nature of 8 and 9 (ca. 17% hard block composition)
whereas broad peaks can be seen for polymers 10 and 11 (ca. 20%
hard block composition) suggesting that they are partly amorphous
(Fig. 5).
-80
-60
-40
-20
0
20
40
60
80
T / °C
Increasing hard segment content produced two notable effects,
firstly, a modest decrease in the Tg of the soft segment (from ca.
Fig. 3. Summary of the temperature-dependent WAXS
packing spacing of end-capped polymers 8e11.
d
spacing revealing local
Please cite this article in press as: D.H. Merino, et al., A systematic study of the effect of the hard end-group composition on the microphase
separation, thermal and mechanical properties of supramolecular polyurethanes, Polymer (2016), http://dx.doi.org/10.1016/
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D.H. Merino et al. / Polymer xxx (2016) 1e11
5
packing of the urea groups in contrast with less ordered association
of the urethane end groups. AFM analysis of 11 revealed assemblies
of fibres (67.5 nm thick on average) which were also densely
packed. The fibres formed by 11 appear to be ribbon-like [18]. The
fibrillar structure obtained suggests that the hard hydrogen bonded
domains are separated by the soft apolar segments, indeed this data
correlated with the variable temperature SAXS analysis which
revealed a broad peak centred at 7.5 nm. Higher order peaks are not
present, which shows the absence of
a long-range ordered
morphology such as hexagonal-packed cylinders. In contrast to
polymer 11, 10 does not form well-defined (fibrillar) aggregates but
exhibits an irregular phase-separated morphology with rounded
domains (Fig. 6) [51]. The difference in morphology is remarkable
and is ascribed to differences in hydrogen-bonding interactions
between the end groups (as shown by the WAXS data in Fig. 4)
which has a substantial effect on the phase separation morphology.
2.5. Mechanical characterisation
Fig. 5. DSC thermograms for the polymer series 8e11. The data shown was obtained
during the second heating ramp at a heating rate of 10 ꢁC/min.
Mechanical characterisation was performed using rheometry,
dynamic mechanical analysis (DMA) and tensile tests. The rheo-
logical behaviour of 8 and 9 was compared to the domain spacing
evolution with temperature in Fig. 7 (Fig. S29). Temperature
dependent changes in microstructure were observed to be consis-
tent with the mechanical response. In particular, the onset of phase
coalescence coincides with a transition in the rheological behaviour
of the materials. This result suggests that in the case of 8 and 9 the
onset of phase coalescence coincides with a transition in the
rheological behaviour of the materials. This result suggests that the
observed decrease in storage modulus versus temperature is
related to phase mixing and to the dissociation of secondary bonds
between the end-groups of the polymers [8]. Isothermal frequency
sweeps were also performed and time-temperature superposition
used to produce master curves presented in the SI (Fig. S29), along
with the shift factor as a function of temperature.
A ‘melting temperature’ was defined as the temperature above
which the storage modulus falls below 0.1 MPa. This temperature
increased from 38 ꢁC for 8e53 ꢁC for 9, as a direct result of the
stronger hydrogen bonding between the urea groups of compound
9. In contrast, examination of the plots of the storage modulus
versus temperature of 10 and 11 did not reveal any transition in
their rheological behaviour over this temperature range (Fig. 8).
The materials remained solid within the entire temperature range
studied (0e120 ꢁC) with storage moduli in the order of 1e10 MPa
(Table 2). It is noted that even after a further increase of the tem-
perature to 200 ꢁC, a rheological transition was not observed for
these materials.
ꢀ30 ꢁC for 8 and 9 to ca. ꢀ35 ꢁC for 10 and 11). The broad
appearance of this transition, which can extend to approximately
0
ꢁC may be as a consequence of the relatively disperse nature
(Ð z 1.8) of these samples. Secondly, the appearance of distinct
transitions for the hard segments of polymers which contained the
more extended end-groups (10 and 11). These higher temperature
thermal transitions (maxima ca. 190e200 ꢁC) are thought to
correspond to dissociation of the non-covalent bonds between the
polymer end-groups. In addition, these findings indicate a well-
defined microphase separation in these two polymers (10 and 11)
[52]. In contrast, the thermograms of polymers 8 and 9 show a
continuous decrease in heat flow at higher temperatures. These
transitions may be attributed to the formation of two composi-
tionally distinct hard block domains comprising primarily either
urea or urethane moieties [29].
The homogeneous morphology of the starting homopolymer
shown by AFM (Fig. S27) confirms the absence of microphase
segregation. The urethane 8 also exhibits no phase separation and
the AFM appears very similar to that of the un-endcapped mid-
block polymer (Fig. S27). In contrast, the interchange of the ure-
thane linkage for
a urea in 9 produced a phase-separated
morphology revealed by both AFM and TEM analysis (see
Fig. S28). The AFM micrograph of 9 (see Fig. S28a) reveals the
different domains produced by the immiscibility of block compo-
nents when compared with the characteristic monophase of the
unendcapped polymer. In addition, TEM analysis confirmed the
existence of microphase separated structure (see Fig. S28b). The
well-defined morphology observed in Fig. S28 corroborates the
capacity of ordering attributed to bis-urethane/urea system. The
pre-organised arrangement adopted by the rigidity of the end-
groups promotes the ordered stacking of the urethane/urea units.
The morphology featured in 10 and 11 consists of the hard segment
(in bright) surrounded by the soft segment matrix (dark) [30]. The
phase separation of the hard/soft segments is pronounced for
polymers 10 and 11 (Fig. 6) which is in agreement with the DSC data
obtained (Fig. 5). Interestingly, the striking differences between the
morphologies of 8 and 9 (see Figs. S27 and S28) and those of 10
(Fig. 6a and b) and 11 (Fig. 6c and d) were generated by a minimal
increase (ca. 3 wt%) of the hard content.
DMA was employed to further probe the mechanical response of
three of the materials (8, 9, 11) at low temperatures. As shown in
Fig. 9, at temperatures below about 0 ꢁC, the compounds present
the standard viscoelastic behaviour of high molecular weight
polymers. A glassy region is followed by the glass transition, with a
Tg (taken at the maximum of tan
all three materials. This finding is consistent with DSC results that
showed the Tg of 11 to be only slightly lower than that of 8 and 9. Of
d
) of around ꢀ10 ꢁC common for
note is the lower max tan
d value of 11 when compared to 8 and 9.
This higher degree of elasticity in 10 could be associated with a
more developed elastic-network for this compound or with the
presence of a crystalline phase. At temperatures above 0 ꢁC, two
distinctly different behaviours were observed. 11 showed a wide
rubbery plateau, with storage modulus in the order of around
100 MPa and no sign of melting up to 80 ꢁC. On the other hand, the
storage modulus of 8 and 9 never reached a plateau. Instead,
beyond the glass transition, E0 continued to decrease with
increasing temperature, albeit at a slower rate, up to an apparent
It has been shown that PUs aggregate in the bulk in hard do-
mains with dimensions ca. 300 nm, referred to as ‘urea balls’ (as
observed in this study with polymer 10, Fig. 6a) [51]. The fibrillar
structure featured by 11 is consistent with the strength of the local
Please cite this article in press as: D.H. Merino, et al., A systematic study of the effect of the hard end-group composition on the microphase
separation, thermal and mechanical properties of supramolecular polyurethanes, Polymer (2016), http://dx.doi.org/10.1016/
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6
D.H. Merino et al. / Polymer xxx (2016) 1e11
Fig. 6. a) AFM and b) TEM image of morpholine terminated urethane polybutadiene derivative 10; c) AFM and d) TEM image of morpholine terminated urea polybutadiene
derivative 11.
107
107
5.72
5.30
5.25
106
105
104
103
102
101
100
G’’
106
105
104
103
102
101
100
G’’
5.70
5.68
5.66
5.64
5.62
5.20
5.15
5.10
5.05
5.00
4.95
4.90
4.85
G’
G’
d/nm
d/nm
0
20
40
60
80
100 120
0
20
40
60
80
100 120
T / °C
T / °C
b)
a)
Fig. 7. G0, G00 and domain spacing of morpholine terminated polybutadiene derivative against temperature and for materials: a) 8 and b) 9 Data obtained in oscillatory shear at a
frequency of 10 Hz and cooling rate of 3 ꢁC per minute.
melting point at about 20e30 ꢁC. These findings confirm the results
from parallel plate oscillatory shear rheometry, and demonstrate
the effect of increasing the strength of non-covalent interactions
between the PU chains. From a practical point of view, it is desired
that the compounds have a broad rubbery plateau followed by a
sharp melting transition, to ensure a stable mechanical response,
while maintaining processability.
The mechanical properties of the materials were explored
further by tensile testing at room temperature and a constant true
strain rate of 0.04 sꢀ1. At least five samples were tested for each
material. Typical stress versus strain traces of 9, 10 and 11 are
shown in Fig. 10, note that data for 8 are not shown because it
produced a low quality film that could not be tested reliably to
failure. The failure strain and ultimate tensile strength (UTS) of the
materials are shown in Table 3. As evidenced by the rheometer data,
the addition of the aromatic rings to the end-groups of the com-
pounds results in a significant increase in stiffness. Here, the tensile
modulus (measured as the slope of the true stress versus true strain
plots between 0.2% and 0.8% strain) increases from 2.5 MPa in the
case of 9 to around 85 MPa for 10 and 55 MPa for 11. At the same
time the true strain at the point of failure decreases dramatically
(from 1.3 for 9 to 0.1 for 10) while the UTS increases from
approximately 0.2 MPa for 9 to around 5 MPa for 10 and 11. These
results can be interpreted by considering the strength of intermo-
lecular non-covalent bonds. In the case of compound 9 the disso-
ciation of the relatively weak hydrogen bonds at low loads is
followed by chain sliding up to the point of rupture. However, it is
proposed that the stronger non-covalent bonds in the case of 10
and 11 act as effective cross-link points that remain intact during
loading which would give rise to a different failure mechanism
Please cite this article in press as: D.H. Merino, et al., A systematic study of the effect of the hard end-group composition on the microphase
separation, thermal and mechanical properties of supramolecular polyurethanes, Polymer (2016), http://dx.doi.org/10.1016/
j.polymer.2016.07.029

D.H. Merino et al. / Polymer xxx (2016) 1e11
7
Table 3
102
101
Tensile properties for polymers 9, 10 and 11.
9
10
11
Modulus [MPa]
Failure true strain
UTS [MPa]
2.5 0.2
1.33 0.20
0.2 0.0
86.7 3.6
0.10 0.04
5.0 0.7
55.1 2.3
0.20 0.06
5.4 0.4
100
10-1
10-2
10-3
10-4
10-5
10-6
where the chains are first stretched and subsequently the network
is broken.
11
10
9
3. Conclusion
8
We have studied the morphology of a series of novel supra-
molecular polyurethanes that differ by hard end group content (ca.
17 or 20%) and the chemical composition (urea or urethane) of the
end group. Studies of the solid state structures of small molecule
analogues of the urea and urethane end groups were conducted.
These showed that the dominant intermolecular interaction for
urethane was a hydrogen bond between the morpholine oxygen
and urethane hydrogen, whereas the urea small molecule exhibited
a stronger bifurcated hydrogen bond. The morphology of the novel
polymers containing these end groups was studied in detail using
multiple techniques. Small angle X-ray scattering (SAXS) experi-
ments have revealed that the polymer 8 containing the weakly
hydrogen bonding urethane end groups and the lower hard end
group content did not form phase separated materials. However,
0
20 40 60 80 100 120 140
Temperature / °C
Fig. 8. Shear storage modulus versus temperature for molecules tested in oscillatory
shear at a frequency of 5 Hz and a cooling rate of 2 ꢁC per minute.
Table 2
Rheology data for polymers 8, 9, 10, 11, E0 was calculated as 3 ꢂ G0, assuming
incompressibility.
8
9
10
11
Melting point [ꢁC]
E0 (20 ꢁC) [MPa]
38
2.5
53
7.2
e
e
32
55
104
103
102
101
100
1
11
9
0.8
8
0.6
0.4
0.2
0
11
9
8
-80 -60 -40 -20
0
20 40 60 80
-80 -60 -40 -20
0
20 40 60 80
Temperature / °C
Temperature / °C
Fig. 9. Tensile storage modulus (left) and loss tangent (tan
d, right) versus temperature for polymers 8, 9 and 11.
increasing the hydrogen bonding capability of the end group by
replacing a urethane with a urea linkage resulted in characteristic
microphase phase separated materials. WAXS analysis of the ma-
terials showed that order within the hard segments increased with
increasing hydrogen bonding strength of the end groups and
increasing hard end group content. This was supported by DSC
analysis which revealed varying the urethane/urea content resulted
in unsymmetrical soft/hard segment phases. TEM and AFM
confirmed the lack of phase separation in the low hard end group
content urethane containing polymer (8). However, TEM and AFM
determined that a 1-D fibrillar structure was obtained when the
hard segment featured a urea (9) whereas the high hard end group
content materials (10 and 11) exhibited a highly ordered 3-D
morphology. The morphology of the materials had a direct effect
on the mechanical properties of the polymers. Urethane 8 could not
be cast into coherent films suitable for testing. However, urea 9
6
5
4
3
2
1
0
11
10
9
0
0.2 0.4 0.6 0.8
1
1.2
True Strain
Fig. 10. True stress versus true strain tensile plots for 9, 10 and 11.
Please cite this article in press as: D.H. Merino, et al., A systematic study of the effect of the hard end-group composition on the microphase
separation, thermal and mechanical properties of supramolecular polyurethanes, Polymer (2016), http://dx.doi.org/10.1016/
j.polymer.2016.07.029

8
D.H. Merino et al. / Polymer xxx (2016) 1e11
with its phase separated morphology was a low tensile modulus
(2.5 MPa), ductile material (failure true strain ¼ 1.33). The two
polymers with the higher hard end group content and extensive 3D
interconnected morphology (10 and 11) were an order of magni-
tude stiffer (E > 55 MPa) but also an order of magnitude less ductile
(failure true strain <0.2) than 9.
carbon grids (S162-3 Formvar/Carbon 300 Mesh Cu) consisted of
placing a droplet of the sample onto the grid and allowing the grid
to dry in air. The sample was then stained with a droplet of uranyl
acetate (1% in water) and left to dry in air prior to analysis.
SAXS/WAXS experiments were conducted at station BM26B of
the European Synchrotron Radiation Facility (ESRF) in Grenoble
(France). The samples were introduced in pans DSC 600 heating
stage from Linkam that were previously pierced in order to produce
X-ray windows. The windows are made of mica and the capsules
were sealed avoiding sample loss. DSC runs are performed to study
the thermal dependency of the interaction between the self-
assembled polyureas and polyurethanes as well as the physical
properties associated to the bisurethane systems [53]. The distance
to the SAXS detector was ca. 5 m using a wavelength of 0.145 nm. A-
2D multiwire detector with a resolution of 512 ꢂ 515 pixels and a
4. Experimental
4.1. Materials
The reagents used were commercially available from either
Acros Chimica, Aldrich Chemical Company or Alfa Caesar and were
used without further purification. The dry solvents involved in the
project were bought and used as supplied or distilled with appro-
priate dehydration agents i.e. tetrahydrofuran (THF) which was
distilled from benzophenone and sodium. All other reagents such
as polybutadiene Krasol LBH2000; Mn (1730) and Mw (2020);
Mw=Mn (1.17); concentration of hydroxyl groups (1.021meq.g 21);
and isomer content (60 wt% of 1,2-;15 wt% of 1,4-cis, and 25 wt% of
1,4-trans), and 4,4’-methylenebis(phenylisocyanate) were supplied
by Henkel Adhesive Technologies, Slough, UK.
pixel size of 260 m
m2. Standard corrections for sample absorption
and background subtraction have been performed. WAXS images
were recorded using a CCD-based X-ray digital camera (Photonic
science) with a pixel size of 44 m
m2.
Thin films of the polyurethanes were produced for mechanical
testing via a solution casting procedure. All of the polyurethanes
were dissolved in THF and the solution was poured into flat PTFE
molds. This was subsequently placed in a vacuum oven at 70 ꢁC
with a pressure of approximately 0.6e0.8 bar for a duration of 24 h.
Thin-layer chromatography (TLC) was performed on aluminium
sheets coated with Merck silicagel 60 F254. Column chromatog-
raphy was performed using either SI60 sorbent silica (40e63
mm)
Polyurethane films of uniform thickness between 200 and 500 mm
supplied from VWR international or Brockmann 1, standard grade,
neutral activated aluminium oxide (ca. 150 mesh) supplied from
Aldrich Chemical Company.
were obtained at the end of this procedure without residual sol-
vent. Rheological analysis was performed on circular samples
(25 mm diameter) that were obtained from the moulded film using
a steel punch cutter. For tensile testing, rectangular samples were
cut with a razor blade and paper end-tabs were bonded to the
samples with a commercial cyanoacrylate adhesive Loctite™. This
sample assembly was found to reduce slippage inside the tensile
grips of the tensometer. The gauge section between the grips was of
25 mm length, 5 mm width and approximately 0.5 mm thickness.
4.2. Instrumentation
1H Nuclear magnetic resonance spectroscopy (1H NMR) was
conducted on either a Bruker DPX250 (250 MHz), a Bruker AMX400
(400 MHz) or a Bruker DPX400 (400 MHz) spectrometers using the
deuterated solvent as the lock. In the same manner, 13C Nuclear
magnetic resonance spectroscopy (13C NMR) was performed on a
Bruker DPX250 (62.5 MHz), a Bruker AMX400 (100 MHz) or a
Bruker DPX400 (100 MHz) spectrometers. Infrared spectroscopy
was performed using a Perkin Elmer 1420-X spectrometer with the
samples analysed as either neat films or in solution between two
potassium bromide or sodium chloride disks.
Bruker MicroToF LCMS has been utilised in order to obtain Mass
spectrometry analysis with ionisation via electrospray and samples
injected by direct insertion via a syringe pump. Electrothermal
digital melting point instrument or DSC on a TA Instruments DSC
Q2000 differential scanning calorimeter were used to find out the
melting point of the SPUs. The average molecular weight of the
polymers generated was determined via Gel Permeation Chroma-
tography (GPC) Polymer Laboratories PL-GPC 220 high temperature
chromatograph fitted with PL Mixed Gel columns at 40 ꢁC. The
samples were dissolved in analytical grade chloroform (5 mg/mL),
and analysed against PL Easy-Cal Polystyrene as the calibrants.
AFM experiments have been carried out at Surrey University in
collaboration with Professor Joe Keddie and Dan Liu using a NT-
MDT instrument (Russia) in tapping mode, the tip spring con-
stant was ca. 9.8 N/m and the frequency was 180 KHz. The samples
were directly drop casted from a solution of 1% (w/v) on a mica
sheet placed on the AFM holder. The mica sheet was cleaved with
adhesive tape before the sample was placed on the mica surface.
The samples were also spin-coated in a silicon wafer from the same
solution 1% (w/v) in THF.
4.3. Mechanical testing
Rheometry was performed on two machines, both using parallel
plate oscillatory shear. Data in Fig. 7 and Fig. S32 were obtained
using a TA instruments AR2000 rheometer. Temperature sweeps
were performed at a single frequency of 10 Hz with a strain
amplitude of 0.1%. Initially, the samples were placed on the
rheometer, heated to 100 ꢁC and held to 1 h. Then, the temperature
was decreased to ꢀ50 ꢁC at a rate of 3 ꢁC per minute whilst data
were collected. For the isothermal frequency sweeps, the samples
were again placed in the rheometer and heated to 70 ꢁC for
approximately 5 min. The temperature was reduced to 20 ꢁC, and
again held for 5 min. The frequency sweeps were the performed
between 0.1 and 100 rad/s at a strain amplitude of 0.1% for each
selected temperature, each time the temperature was held for
5 min before data were gathered. Three cycles of tests were
measured, and the results show that the repeating heating does not
change the properties significantly (SI).
Experiments reported in Fig. 8 were performed with an Anton
Paar Physica MCR 301 rheometer. For the temperature sweep at a
single frequency of 5 Hz and a strain amplitude of 0.1%., the samples
were placed on the rheometer, heated to 70 ꢁC and held at this
temperature for approximately 5 min. The samples were then
subjected to a cooling temperature ramp of 2 ꢁC/min to 0 ꢁC before
being heated up to 120 ꢁC and held at this temperature for 1 h. After
this step they were cooled again from 120 ꢁC to 0 ꢁC at 2 ꢁC/min.
TEM analysis was performed in the Centre for Advance Micro-
scopy at the University of Reading (CFAM) using a CM20 Trans-
mission Electron Microscope with an electron beam voltage of
80 kV with assistance of Dr Peter Harris. Sample preparation on
During all of these steps the dynamic shear moduli (G0, G00 and tan
d)
of the samples were recorded. This method was applied to ensure
that chemical cross-linking had not occurred during the exposure
of the samples to high temperatures. Cross-linking was not
Please cite this article in press as: D.H. Merino, et al., A systematic study of the effect of the hard end-group composition on the microphase
separation, thermal and mechanical properties of supramolecular polyurethanes, Polymer (2016), http://dx.doi.org/10.1016/
j.polymer.2016.07.029

D.H. Merino et al. / Polymer xxx (2016) 1e11
9
observed, i.e. an increase in G0 during the isothermal step at 120 ꢁC
was not observed.
IR (thin film) ymax/cmꢀ1 3335, 2946, 2817, 1675, 1598, 1552, 1502,
1329, 1300, 1255, 1224, 1110, 850, 751; 1H NMR (400 MHz, CDCl3)
Dynamical mechanical analysis (DMA) was performed with a TA
Q800 DMA, in tensile mode. The samples were subjected to a
heating ramp from ꢀ80 ꢁC up to 80 ꢁC at a rate of 2 ꢁC per minute.
Dynamic tension was applied at a frequency of 5 Hz and a strain
amplitude of 0.1%. To avoid buckling, a ratio of static to dynamic
force of 120% was maintained throughout each test.
d
2.52 (4H, t, J ¼ 4.4), 2.57 (2H, t, J ¼ 5.6), 3.40 (2H, q, J ¼ 5.6), 3.74
(4H, t, J ¼ 4.4), 5.46 (1H, t, J ¼ 4.4), 7.56 (2H, AA’XX0 system, J ¼ 9.2),
8.17 (2H, AA’XX0 system, J ¼ 9.2); 13C NMR (100 MHz, CDCl3)
d 36.9,
53.5, 58.0, 66.8, 117.8, 125.2, 142.2, 145.6, 154.6; MS (CI) calc. for
C13H19N4O4: 295.1406, found 295.1415.
The tensile testing was carried out on an Instron universal
testing machine (model 5982) with a 100 N load cell. A constant
true strain rate of 0.04 sꢀ1 was used, for the purposes of controlling
the experiment (but not subsequent data analysis) this was calcu-
lated from the cross-head displacement. Strains for data analysis
were calculated from digital image correlation. A fine speckle
pattern was sprayed onto the specimen surface and imaged during
the experiment using a digital camera. The images were analysed
using commercial image analysis software (DaVis V8, LaVision) to
produce specimen strains.
4.7. Synthesis of 2-morpholinoethyl 4-aminophenylcarbamate 1
The nitro derivative previously produced was dissolved in
ethanol (25 mL) then Pd/C (5% w/w, 250 mg) was added. The re-
action was sealed with rubber septum and H2 from a balloon
bubbled through the rapidly stirring suspension for 10 min. The
reaction was left before stirring under an H2 atmosphere for 1 h.
The bubbling and stirring cycle was repeated twice more. The re-
action was continued to completion then filtered through a pad of
Celite® which was washed thoroughly with ethanol and the solvent
removed to leave a brown oil which crystallised on standing,
4.4. X-ray structure determination
affording a yellow solid in yield of 86% (0.79 g). IR (thin film) ymax
cmꢀ1 3338, 2958, 2858, 1710, 1630, 1601, 1516, 1430, 1303, 1227,
1113, 1067, 830, 516; 1H NMR (400 MHz, CDCl3)
2.50 (4H, br), 2.63
/
Crystals of 12 and 13 were mounted under Paratone-N oil and
flash cooled to 150 K under nitrogen in an Oxford Cryosystems
Cryostream. Single-crystal X-ray intensity data were collected using
ꢀ
d
(2H, t, J ¼ 5.6), 3.51 (2H, br) 3.70 (4H, t, J ¼ 4.4), 4.25 (2H, t, J ¼ 5.6),
an Agilent Gemini
S Ultra diffractometer (Cu Ka radiation
6.60 (2H, d, J ¼ 8.8), 7.03 (1H, br), 7.11 (2H, d, J ¼ 7.2); 13C NMR
€
ꢀ
€
(e ¼ 1.54,180 Å) or Mo Ka radiation (e ¼ 0.71,073 Å). The data were
reduced within the CrysAlisPro software [54]. The structures was
solved using the program SIR92 [55] and all non-hydrogen atoms
located. Least-squares refinements on F were carried out using the
CRYSTALS suite of programs [56]. The non-hydrogen atoms were
refined anisotropically. All the hydrogen atoms were located in
difference Fourier maps, then those attached to C were placed
geometrically with a CꢀH distance of 0.95 Å and a Uiso of 1.2 times
the value of Ueq of the parent C atom. The fractional coordinates of
the H atoms attached to the N atoms were refined freely whilst
those of the hydrogen atoms attached to C were then refined with
riding constraints. All crystals of compound 13 were found to be
racemic twins.
(100 MHz, CDCl3) d 53.7, 57.5, 61.4, 66.8, 115.5, 116.7, 121.0, 129.1,
138.6, 142.8, 154.0; MS (ESI) calc. for C13H20N3O3: 266.1499, found
266.1502.
4.8. Synthesis of 1-(4-aminophenyl)-3-(2-morpholinoethyl)urea 2
The urea 2 was produce according to the procedure used to
synthesise
1 to afford 2 as a pink solid (0.74 g, 83%).
m.p. ¼170e171 ꢁC; IR (thin film) ymax/cmꢀ1 3340, 2956, 2868, 2823,
1646, 1602, 1559, 1514, 1285, 1225, 1183, 1113, 684, 516; 1H NMR
(400 MHz, d6-DMSO) 2.33e2.36 (6H, m), 3.15 (2H, q, J ¼ 5.6), 3.57
(4H, t, J ¼ 4.8), 4.68 (2H, br), 5.83 (1H, t, J ¼ 5.2) 6.45 (2H, AA’XX0
system, J ¼ 8.6), 6.98 (2H, AA’XX0 system, J ¼ 8.6), 8.04 (1H, s); 13C
4.5. Synthesis of 2-morpholinoethyl 4-nitrophenylcarbamate 5
NMR (100 MHz, d6-DMSO) d 36.0, 53.2, 58.0, 66.1, 114.1, 120.2, 129.6,
143.3, 155.6. MS (ESI) calc. for C13H21N4O2: 265.1559, found
265.1659.
To
a stirred solution of nitrophenylisocyanate (4.00 g,
24.4 mmol) in THF 75 mL under nitrogen was added N-(2-
hydroxylethyl)-morpholine (3.52 mL, 26.8 mmol) dropwise from
a syringe. After 18 h at room temperature the solvent was removed
under reduced pressure to leave a sticky yellow solid. Trituration in
hexane (50 mL) afforded the target urethane as a yellow powder
(6.20 g, 86%). Recrystallization from EtOAc (50 mL) and methanol
(17 mL) of a portion from chloroform produced orange crystals
suitable for X-ray crystallographic analysis. m.p. ¼ 108e109 ꢁC; IR
(thin film) ymax/cmꢀ1 3265, 3087, 2860, 2822, 1738, 1613, 1596,
1555, 1504, 1454, 1414, 1333, 1304, 1223, 1177, 1112, 1067, 941, 854,
4.9. Synthesis of 4-((4’-carbamic acid 2-[bismorpholine]-ethyl
ester) benzyl)-phenyl-amino-carbonyl terminated poly(butadiene)
diol 8.
Poly(butadiene) diol end-capped with 4,4’-methylenebi-
s(phenylisocyanate) (8.06 g, 3.23 mmol) was dissolved in dry THF,
N-(2-hydroxylethyl)-morpholine (0.84 g, 6.45 mmol) was added
and the solution was left stirring at 50 ꢁC for 6 h. The solvent was
removed in vacuo and 4-((4’-carbamic acid 2-[bismorpholine]ethyl
ester) benzyl)-phenyl-amino-carbonyl terminated poly(butadiene)
diol 8 diol was purified as a light orange viscous liquid in 88% yield
(8.06 g) via a multiple slow precipitation in methanol. IR (CDCl3,
KBr) nmax/cmꢀ1; 3324, 3072, 2963, 2915, 2843, 1734, 1703, 1638,
751, 690, 531; 1H NMR (250 MHz, CDCl3)
d
2.53 (4H, t, J ¼ 4.1), 2.69
(2H, t, J ¼ 5.6), 3.74 (4H, t, J ¼ 4.6), 4.34 (2H, t, J ¼ 5.6), 7.22 (1H, s),
7.55 (2H, AA’XX0 system, J ¼ 9.1), 8.20 (2H, AA’XX0 system, J ¼ 9.1);
13C NMR (62.5 MHz, CDCl3)
d 53.8, 57.3, 62.3, 66.8,117.7,125.2,143.1,
143.9.0, 152.7; (NCO2; MS (CI) calc. for C13H17N3O5: 295.1168, found
295.1169.
1598, 1533, 1436; 1H NMR (250 MHz, CDCl3)
d 1.25e2.02 (18Hn, br),
2.50e2.53 (8H, t, J ¼ 5.0), 2.64e2.69 (4H, t, J ¼ 5.0), 3.71e3.74 (8H, t,
J ¼ 5.0), 3.88 (4H, s), 4.26e4.30 (4H, t, J ¼ 5.0), 4.96e5.84 (7Hn, br),
6.55e6.88 (4H, br), 7.08e7.11 (8H, AA’XX0 system), 7.26e7.29 (8H,
4.6. Synthesis of 1-(2-morpholinoethyl)-3-(4-nitrophenyl)urea 6.
Synthesis of nitro urea 6 was achieved by the addition of 4-(2-
aminoethyl)morpholine (3.52 mL, 26.8 mmol) to 4-nitrophenyl
isocyanate (4.00 g, 24.4 mmol) following the procedure detailed
for 5 to give 6 as an orange powder (6.05 g, 84%). m.p. 191e194 ꢁC;
AA’XX0 system); 13C NMR (62.5 MHz, CDCl3)
d 25.6, 27.4, 30.2, 32.8,
38.2, 39.7, 43.8, 53.8, 57.5, 66.9, 68.0, 112.9e114.9, 118.9, 127.7, 130.1,
131.7, 135.9e136.4, 142.7e143.6, 153.4e153.7; GPC (THF) Mw
13,079, Mn7244, Ð 1.8.
Please cite this article in press as: D.H. Merino, et al., A systematic study of the effect of the hard end-group composition on the microphase
separation, thermal and mechanical properties of supramolecular polyurethanes, Polymer (2016), http://dx.doi.org/10.1016/
j.polymer.2016.07.029

10
D.H. Merino et al. / Polymer xxx (2016) 1e11
4.10. Synthesis of 4-((40-2-[bismorpholine]-ethyl) ureidobenzyl)-
phenyl-amino-carbonyl terminated poly(butadiene) 9
powder, 0.96 g, 91%. Mp 104 ꢁC; IR (ATR) nmax/cmꢀ1 1222, 1313,
1445, 1500, 1567, 1720, 2816, 2866, 2950, 3140, 3257; 1H NMR
(400 MHz; CD3Cl3)
d
2.53 (4H, t, J ¼ 4.4), 2.67 (2H, t, J ¼ 5.6), 3.73
This product has been produced via the direct addition approach
used to synthesise polymer 8. 4-((40-2-[Bismorpholine]-ethyl)
ureidobenzyl)-phenyl-amino-carbonyl terminated poly(butadiene)
diol 9 was produced as an elastomeric solid in a 97% yield (3.18 g)
from poly(butadiene) diol end-capped with 4,4’-methylenebi-
s(phenylisocyanate) (2.88 g, 1.15 mmol) with N-(2-aminoethyl)-
morpholin-4-(2-aminoethyl)morpholine (0.30 g, 2.31 mmol). IR
(CDCl3, KBr) nmax/cmꢀ1 3337, 3072, 2963, 2915, 2844, 1734, 1639,
(4H, t, J ¼ 4.8), 4.30 (2H, t, J ¼ 5.6), 6.79 (1H, br), 7.06 (1H, t, J ¼ 7.6),
7.28e7.38 (4H, m); 13C NMR (100 MHz; CDCl3)
d 53.8, 57.5, 61.8,
66.8, 118.6, 123.51, 129.1, 137.8, 153.4.
4.14. Synthesis of 1-(2-morpholinoethyl)-3-phenylurea 13
Phenyl isocyanate (3.8 mmol, 0.46 g) was added to 2-
morpholinoethan-1-amine (3.8 mmol, 0.50 g) in dry THF (10 mL)
and the mixture was stirred for 16 h at room temperature. The
solvent was then removed in vacuo to afford compound 13 as a
white powder, 0.89 g, 94%. Mp 140 ꢁC; IR (ATR) nmax/cmꢀ1 1243,
1309, 1442, 1550, 1594, 1653, 2811, 2857, 2940, 2969, 3310; 1H NMR
1598, 1544, 1414; 1H NMR (250 MHz, CDCl3)
d 1.25e2.02 (18Hn, br)
2.41e2.50 (12H, t, J ¼ 5.0), 3.29e3.36 (4H, br), 3.61e3.64 (8H, t,
J ¼ 5.0), 3.88 (4H, s), 4.96e5.77 (7Hn, br), 6.66 (4H, br), 7.04e7.11
(8H, AA’XX0 system), 7.26e7.29 (8H, AA’XX0 system); 13C NMR
(62.5 MHz, CDCl3) d 25.0, 27.4, 30.2, 32.7, 37.4e38.4, 41.0, 43.5, 53.3,
(400 MHz; CD3Cl3)
d
2.34 (4H, br), 2.38 (2H, t, J ¼ 6.0), 3.28 (2H, q,
57.8, 66.8, 113.8e115.0, 121.2e121.5, 128.0, 129.4, 131.8,
136.2e136.8, 142.7e143.1, 153.7-156-2; GPC (THF) Mw 10,229,
Mn6010, Ð 1.7.
J ¼ 5.6), 3.59 (4H, br), 6.15 (1H, s), 6.96 (1H, t, J ¼ 7.6), 7.20 (2H, t,
J ¼ 7.6), 7.31 (2H, d, J ¼ 7.6); 13C NMR (100 MHz; CDCl3)
d 36.8, 53.4,
57.9, 66.8, 119.8, 122.7, 128.9, 139.3, 156.8.
4.11. Synthesis of bis(4-(4-(3-(4-(3-(2-morpholinoethyl)ureido)
phenyl)ureido) benzyl)phenylcarbamate) terminated
poly(butadiene) 10
Acknowledgements
The authors would like to thank EPSRC (EP/J010715/1 and EP/
J011436/1) for post-doctoral fellowships for AF and IG; EP/
G026203/1 and EP/D07434711 for BWG), and Henkel UK Ltd for
PhD studentships for DHM and KM. Spectroscopic and thermal data
were acquired in the Chemical Analysis Facility at the University of
Reading. AFM and TEM images were obtained using the Electron
Microscopy Laboratory (EMLab) at the University of Reading and
the microscopy facilities at the University of Surrey.
This product has been produced via the direct addition approach
used to synthesise polymer 8. 4-((40-2-[Bismorpholine]-ethyl)ure-
idobenzyl)-phenyl-amino-carbonyl terminated poly(butadiene)
diol 10 was produced as a pink elastomeric solid in a 95% yield
(7.20 g) from poly(butadiene) diol end-capped with 4,4’-methyl-
enebis(phenylisocyanate) (6.25 g, 2.5 mmol) with (1.32 g,
5.0 mmol) 2-morpholinoethyl 4-aminophenylcarbamate 1. IR
(CDCl3, KBr) nmax/cmꢀ1 3309, 3072, 2969, 2915, 2843, 1728, 1711,
Appendix A. Supplementary data
1693, 1595, 1519, 1412; 1H NMR (250 MHz, CDCl3)
d 1.25e2.02
(18Hn, br) 2.51 (8H, br), 2.65 (4H, br), 3.69 (8H, br), 3.80e3.87 (4H,
m), 4.26e4.34 (4H, br), 4.96e5.79 (7Hn, br), 6.55 (4H, br), 6.85e6.87
Supplementary data related to this article can be found at http://
dx.doi.org/10.1016/j.polymer.2016.07.029.
(4H, br), 7.04e7.47 (24H, br); 13C NMR (62.5 MHz, CDCl3)
d 25.0,
27.4, 28.6, 32.0, 38.2, 41.7, 43.6, 53.8, 57.5, 61.7, 66.8, 113.8e115.1,
119.1e121.5, 128.3, 130.2, 131.8, 135.8e136.2, 142.7e143.4, 153.9;
GPC (THF) Mw 12,538, Mn7249, Ð 1.7.
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