A biomimetic modular polymer with tough and adaptive properties

Article abstract of DOI:10.1021/ja9009666

(Figure Presented) Natural materials employ many elegant strategies to achieve mechanical properties required for survival under varying environmental conditions. Thus these remarkable biopolymers and nanocomposites often not only have a combination of me

Full text of DOI:10.1021/ja9009666

Published on Web 06/09/2009
A Biomimetic Modular Polymer with Tough and Adaptive Properties
Aaron M. Kushner, John D. Vossler, Gregory A. Williams, and Zhibin Guan*
Department of Chemistry, UniVersity of California, 1102 Natural Sciences 2, IrVine, California 92697-2025
Received February 6, 2009; E-mail: zguan@uci.edu
Scheme 1. Synthesis of Biomimetic Linear Modular Polymer
Whereas man-made polymers can be prepared to meet particular
parameters one at a time, it remains a challenge to design synthetic
polymers with a combination of mechanical properties such as high
modulus, toughness, and resilience. A further challenge is to
introduce adaptive properties into polymers. In contrast, many smart
strategies have evolved in nature to achieve biopolymers possessing
excellent combinations of mechanical properties.1 To survive in
often variable environments, natural materials have also evolved
to be adaptive, maintaining functions across a range of stress or
strain, or changing properties in response to stimuli such as
temperature or moisture level.2 In recent years, the elucidation of
molecular mechanisms for natural materials has prompted many
biomimetic materials designs.3 Herein, we report a biomimetic
design of a modular polymer that has a combination of high
modulus, toughness, and resilience, while possessing adaptive
mechanical properties.
Our biomimetic concept is based on the modular domain design
observed in the skeletal muscle protein titin, which possesses a
remarkable combination of strength, toughness, and elasticity.4 The
ability of titin to absorb energy by the reversible rupture of
intramolecular secondary interactions, followed by refolding induced
recovery, makes it an intriguing model for the design of adaptive
materials. Following titin’s modular design, our group first syn-
thesized polymers incorporating the quadruple hydrogen bonding
2-ureido-4[1H]-pyrimidone (UPy) motif5 as the modular domain-
forming mimic of the Ig domains in titin.6 To overcome issues
such as structural heterogeneity and interchain cross-linking, we
further developed a cyclic modular polymer using a peptidomimetic
ꢀ-sheet dimer.7 Despite the well-defined single molecule unfolding
properties, the synthesis of the second generation polymers is
tedious, and the rupture forces of the H-bonded modules are much
lower than the UPy dimer. To simplify synthesis and improve
mechanical strength, the cyclic modular concept was applied to
the UPy core. In our previous report, a cyclic UPy dimer was
incorporated as a cross-linker into a 3-D network that then showed
significant enhancement of mechanical properties,8 which we
attribute to the reversible opening of the closed UPy dimer.
Well-defined hydrogen bonding motifs have been introduced to
a number of polymers at chain ends,5,9,10 in the main chain,6,11 or
at side chains,12-14 in the pursuit of new material designs. However,
most of the studies were focused on the dynamic properties in
solution or melt.9-14 In one study, UPy groups were introduced as
thermoreversible interaction sites to a chemically cross-linked
polymer network, resulting in shape memory properties.15 Never-
theless, this thermoset is not processable and its mechanical
properties are moderate in terms of strength and extensibility. Here
we synthesized a linear polymer composed of a tandem array of
biomimetic cyclic UPy modules, closely mimicking the titin
architecture, yielding a strong, tough, processable, and highly
adaptiVe material.
by the Grubbs Gen-2 catalyst, acyclic diene metathesis (ADMET)16
polymerization of monomer 3 (Scheme 1) afforded polymer 1 (Mn
) 18.0 kDa, PDI ) 1.7; see Supporting Information (SI) for
experimental details). The polymer is soluble in CHCl3 or CHCl3/
DMF (9/1 v/v), and homogeneous transparent films are easily
prepared from solution casting. Dynamic Light Scattering (DLS)
of polymer 1 in solution shows a single peak around 7-8 nm (see
Figure S9 in SI), indicating no aggregation present in the polymer
solution.
The system was designed so that the polymer backbone has no
hydrogen-bonding or other strong interaction sites that may
complicate the reversible opening and closing of the UPy modules.
ADMET was chosen because it results in a hydrocarbon backbone
with minimal secondary interactions, and the hydrocarbon spacer
is attached to the UPy module by an ester linkage rather than the
H-bonding carbamate linkage of our previous system.8 We decided
to make a homopolymer of 3 by incorporating a UPy module in
every repeat unit to minimize the possibility of phase segregation,
which would likely prevent the linear tandem UPy modules, trapped
in rigid-phase aggregates, from exhibiting titin-like sequential
unfolding and reversible energy-absorbing behavior. As a control,
monomer 3 was protected with the ortho-nitrobenzyl (NBn) group
to block H-bonding and then polymerized to yield polymer 2. The
control polymer is almost identical to the real sample except that
the protected UPy units cannot form dimers. As will be shown later,
without reversible hydrogen-bonding capability, the control polymer
2 is drastically different from the real sample (1), lacking toughness,
extensibility, or any adaptive properties.
Figure 1 compares the stress-strain curves for polymer 1 and
the control 2. While control is brittle and fractures at 7% strain,
polymer 1 undergoes large deformation with maximal strain >100%.
Polymer 1 is stiff with a relatively high Young’s modulus, ∼200
MPa. After yielding at ∼5% strain, it shows a large deformation,
with a relatively small increase in stress, resulting in the absorption
of a large amount of energy. We postulate that this high toughness
The cyclic UPy core was synthesized as described previously8
and further converted to a terminal diolefin monomer 3. Catalyzed
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Figure 2. Shape-memory cycle for polymer 1 (color code: red, temperature;
blue, stress; and black, strain).
quickly recovered to its original dimension upon further heating,
with both shape fixity and recovery approaching 100% (Figure 2).
The unusual stress-strain behavior (Figure 1) and the unique
shape-memory properties (Figure 2) indicate that the modular
polymer 1 exhibits a combination of mechanical properties found
in titin, including high modulus, toughness, large extensibility, and
intriguing adaptiVe behavior. In contrast, the H-bonding blocked
control polymer 2 is a simple brittle material, suggesting that the
reversible rupture of UPy dimer modules contributes to the
macroscopic properties. Our proposed mechanism for these ob-
served properties is illustrated in Figure 3. Upon stretching, the
modules gradually unfold, resulting in a large extension and
absorption of energy. When the polymer is cooled and stress is
removed, many of the unfolded UPy units are in close proximity
to other opened units on neighboring chains, with which they can
dimerize to form a temporary network, freezing the new polymer
shape. When temperature/time is applied to the system, the newly
formed interchain UPy dimers become dynamic, and the modules
can return to their original more stable cyclic self-dimerized state,
regaining their original dimensions and properties.
Figure 1. Stress-strain curves for sample 1 and control 2.
behavior is enabled at the molecular level by the continuous
unfolding of UPy dimer modules upon stretching. Whereas
traditional thermoplastic polymers can undergo large plastic
deformation after yielding, this is usually accomplished through
crazing or necking leading to permanent damage.17 In our case,
however, neither necking nor crazing was observed. The specimens
remain uniform and transparent throughout the tests.
In our further studies, we observed that this plastic deformation
is not permanent and could gradually recover with time or upon
heat treatment, suggesting an interesting “self-healing” property for
polymer 1. After each loading-unloading cycle, the sample did
not return to its original length, resulting in a temporary set. After
10 strain cycles, the material was set at ∼135% of its original length
(blue curve, Figure 1). With time this sample gradually regained
its shape and properties, recovering to 110% of its original length
overnight at rt (green curve, Figure 1). In addition to slow
spontaneous recovery, heating the material to ∼80 °C induced full
recovery of dimension, stiffness, strength, and toughness in ∼30 s,
returning these properties to a level comparable to that achieved
by slow healing at rt. It should be noted that the unusual
stress-strain behavior reflects the inherent mechanical properties
of the polymer and is highly reproducible.18 Because the sample
did not recover to its original length following the loading-unloading
cycle, the initial region of subsequent stretching showed zero-stress
until the sample reached the new set length. During unloading, the
sample tension reaches zero-stress before the strain reaches the new
temporary set point and then recovers fully over ∼60 s, so that
stress again increases when the new set strain is reached on the
next loading cycle. Similar stress-strain behavior has been observed
in biological materials such as silks,19 though, in that case, the
postloading strain at zero-stress represents a permanent set. Finally,
to test the sensitivity of the polymer to water, we performed static
tensile tests on a polymer 1 sample after incubating it in water for
18 h. Excitingly, very little change was observed in its tensile
mechanical properties after this treatment (see SI, Figure S17).
Presumably, due to the hydrophobic nature of the polymer, water
was not able to effectively permeate the material and, thus, had no
detrimental effect on the mechanical properties.
Figure 3. Proposed molecular mechanism.
The proposed molecular mechanism accounting for the rare
combination of mechanical properties of polymer 1 is supported
by the following factors. First, our single molecule force spectros-
copy (SMFS) for previous and current modular polymers shows
sawtooth force-extension curves corresponding to tandem unfold-
ing of biomimetic modules,6,7,20 indicating that, at molecular scale,
UPy modules can unfold sequentially under stress. Second,
computer modeling demonstrates that cross-dimerization of unfolded
UPy units from adjacent chains is totally feasible (see SI, Figure
S18-20). Third, the transition temperature of the shape-memory
cycle (Ttrans) of polymer 1 matches the UPy dimer unfolding
temperature, suggesting a direct connection between the UPy
module unfolding event and the observed macroscopic properties.
The Ttrans of the shape-memory cycle for polymer 1 centers at ∼50
°C (Figure 2), agreeing with a TR transition (∼52 °C) observed in
To further explore the adaptive properties of this material, we
investigated the temperature dependence of the self-healing process
and observed an interesting shape-memory behavior (Figure 2).
After heating to 80 °C, the sample was extended to 250% strain
and then cooled to 5 °C to freeze the shape. The stress was then
released, and then the temperature was ramped up gradually. After
reaching ∼27 °C, the sample specimen began to retract and then
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the tan δ peak of its DMA temperature scan (see SI, Figure S10),
which, most importantly, matches the UPy unfolding temperature.15
Fourthly, our negative control polymer 2 has shown that blocking
the hydrogen bonds of the UPy unit completely destroys the high
toughness and adaptive properties, confirming the importance of
reversible hydrogen bonding for the macroscopic properties. Lastly,
our X-ray, DSC, and AFM data have ruled out the possibility that
microphase separation or crystallinity is responsible for the observed
properties. SAXS data reveal a complete lack of phase separation
on the scale of 2-45 nm for both untreated and prestretched samples
(see SI, Figure S11-12). AFM imaging further confirms no
microphase separation in polymer 1 (see SI, Figure S15). The
WAXS (see SI, Figure S14) pattern reveals a broad peak at ∼4.6
Å, which molecular modeling indicates is a reasonable match with
the spacing between two adjacent UPy dimers. An additional peak
is observed in the WAXS spectrum at ∼2.2 nm, corresponding to
the approximate molecular size of the UPy dimer macrocycle.21
The DSC trace shows no crystallinity for polymer 1 (see SI, Figure
S16), confirming its amorphous nature in bulk.
In summary, we have demonstrated the first biomimetic modular
polymer that exhibits a rare combination of high modulus, high
toughness, and adaptiVe properties, including self-healing and shape
memory. Notably, this new polymer combines high toughness and
self-healing properties in one material, something that has proven
extremely difficult to achieve.24 Our ongoing research seeks to
elucidate the structure property relationships of various aspects of
the modular polymer architecture, such as the loop size and
intermodule chain length, as well as new UPy module macrocycle
morphology. These studies are expected to yield further novel
biomimetic polymeric materials with advanced properties.
Acknowledgment. We acknowledge the financial support from
the National Institute of Health (R01EB004936) and U.S. Depart-
ment of Energy (DE-FG02-04ER46162). We thank Professor Adam
Summers for the MTS instrument, Jane Bai for AFM imaging, and
M. Davinagarcia Jr. for SAXS/WAXS spectra. Z.G. acknowledges
a Camille Dreyfus Teacher-Scholar Award. A.K. acknowledges an
Eli Lilly fellowship.
For shape-memory polymers, the molecular mechanism for
“memorizing” the initial permanent shape is usually one of the
following: (1) permanent covalent cross-links; (2) physical cross-
links such as crystalline domains or segregated microphases; or
(3) chain entanglement for polymers with very high molecular
weight.22 For polymer 1, the first two possibilities are ruled out
because there are no covalent cross-links and the system is totally
amorphous. This leaves the third one, chain entanglement, as the
most probable mechanism for holding the permanent shape and
preventing permanent flowing at elevated temperature. Although
our polymer 1 molecular weight is not particularly high, the unusual
molecular topology should contribute to the formation of stable
entanglements. With bulky and rigid UPy modules dispersed
between flexible alkene spacers, the chains can easily be trapped
in entangled states, as was shown with the iptycene polymers by
Swager and co-workers.23 Though interchain UPy cross-dimeriza-
tion is feasible and, as discussed above, plays a key role in the
temporary fixation of polymer shape, it should be noted that this
temporary cross-linking cannot be responsible for holding the
permanent initial shape, as deformation at elevated temperature
would lead to a whole new H-bond network, replacing the initial
one, and the original shape would be completely erased.
Supporting Information Available: Synthesis and characterization
of monomers and polymers, stress-strain, DMA, X-ray, AFM, and
molecular modeling experiments. This material is available free of
charge via the Internet at http://pubs.acs.org.
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