U.S. patent application number 16/813969 was filed with the patent office on 2020-11-05 for blends of oligopeptide terminal polyisobutylene or polystyrene.
The applicant listed for this patent is ECOLE POLYTECHNIQUE FEDERALE DE LAUSANNE (EPFL). Invention is credited to Emmanuel Croisier, Holger Frauenrath, Su Liang, Veronique Michaud.
Application Number | 20200347217 16/813969 |
Document ID | / |
Family ID | 1000004959702 |
Filed Date | 2020-11-05 |
View All Diagrams
United States Patent
Application |
20200347217 |
Kind Code |
A1 |
Croisier; Emmanuel ; et
al. |
November 5, 2020 |
BLENDS OF OLIGOPEPTIDE TERMINAL POLYISOBUTYLENE OR POLYSTYRENE
Abstract
Various blends of polymers are disclosed, comprising
oligopeptide functionalised polymers such as polyisobutylene and
polystyrene. Mono-functionalised and di-functionalised polymers
(each containing 0 to 5 peptide units beyond its terminal amide
group) may be blended with each other and/or with
non-functionalised polymers to produce blended compositions. Such
compositions are of use, for example, in vibrations dampers.
Certain blends also exhibit self-healing properties.
Inventors: |
Croisier; Emmanuel;
(Lausanne, CH) ; Frauenrath; Holger; (Lausanne,
CH) ; Liang; Su; (Lausanne, CH) ; Michaud;
Veronique; (Lausanne, CH) |
|
Applicant: |
Name |
City |
State |
Country |
Type |
ECOLE POLYTECHNIQUE FEDERALE DE LAUSANNE (EPFL) |
Lausanne |
|
CH |
|
|
Family ID: |
1000004959702 |
Appl. No.: |
16/813969 |
Filed: |
March 10, 2020 |
Related U.S. Patent Documents
|
|
|
|
|
|
Application
Number |
Filing Date |
Patent Number |
|
|
14646950 |
May 22, 2015 |
10752765 |
|
|
PCT/EP2013/074793 |
Nov 26, 2013 |
|
|
|
16813969 |
|
|
|
|
Current U.S.
Class: |
1/1 |
Current CPC
Class: |
C09D 125/06 20130101;
C09D 123/36 20130101; C08L 2205/02 20130101; C08L 2205/025
20130101; C08H 1/00 20130101; C08L 23/36 20130101; C08L 25/06
20130101; C08L 89/00 20130101; C08L 23/22 20130101 |
International
Class: |
C08L 23/36 20060101
C08L023/36; C08L 23/22 20060101 C08L023/22; C08L 89/00 20060101
C08L089/00; C08H 1/00 20060101 C08H001/00; C08L 25/06 20060101
C08L025/06; C09D 123/36 20060101 C09D123/36; C09D 125/06 20060101
C09D125/06 |
Foreign Application Data
Date |
Code |
Application Number |
Nov 26, 2012 |
GB |
1221246.0 |
Claims
1. A polymer blend, comprising: a first oligopeptide-terminal
polymer component selected from the group consisting of: a
hydrophobic, flexible polymer having a glass transition temperature
below 20.degree. C. and only one monodisperse oligopeptide end
group, the monodisperse oligopeptide end group having 1 to 5 amino
acid repeating units; and a hydrophobic, flexible polymer having a
glass transition temperature below 20.degree. C. and two
monodisperse oligopeptide end groups; and at least one additional
polymer component selected from the group consisting of: a
hydrophobic, flexible polymer that is different from said first
oligopeptide-terminal polymer component and that has a glass
transition temperature below 20.degree. C. and only one
monodisperse oligopeptide end group, the monodisperse oligopeptide
end group having 1 to 5 amino acid repeating units; and a
hydrophobic, flexible polymer that is different from said first
oligopeptide-terminal polymer component and that has a glass
transition temperature below 20.degree. C. and two monodisperse
oligopeptide end groups.
2. The polymer blend according to claim 1, wherein at least one of
said first oligopeptide-terminal polymer component and said at
least one additional polymer component comprises repeating units
selected from the group consisting of isobutylene, butadiene,
siloxane, acrylate, and fluoropolymer units.
3. The polymer blend according to claim 1, wherein at least one of
said first oligopeptide-terminal polymer component and said at
least one additional polymer component comprises one or more of
isobutylene, isoprene or styrene units.
4. The polymer blend according to claim 1, wherein said first
oligopeptide-terminal polymer component and said at least one
additional polymer component include a hydrophobic, flexible
isobutylene polymer having a glass transition temperature below
20.degree. C. and only one monodisperse oligopeptide end group, the
monodisperse oligopeptide end group having 1 to 5 amino acid
repeating units, blended with a hydrophobic, flexible styrene
polymer having a glass transition temperature below 20.degree. C.
and only one monodisperse oligopeptide end group, the monodisperse
oligopeptide end group having 1 to 5 amino acid repeating
units.
5. The polymer blend according to claim 1, wherein an oligopeptide
moiety of said first oligopeptide-terminal polymer component
comprises L-alanine units.
6. The polymer blend according to claim 1, wherein a polymer
segment of at least one of said first oligopeptide-terminal polymer
component and said at least one additional polymer component is
selected from the group consisting of: polyisobutylene,
poly(isobutylene-co-isoprene), polyisoprene, polybutadiene,
polysiloxane, polyacrylate, poly(ethylene-co-butylene),
hydrogenated poly(isoprene), hydrogenated poly(butadiene), and a
fluoropolymer.
7. The polymer blend according to claim 6, wherein said
fluoropolymer is poly(tetrafluoroethylene-co-ethylene).
8. The polymer blend according to claim 1, wherein said first
oligopeptide-terminal polymer component and said at least one
additional polymer component include a hydrophobic, flexible
polyisobutylene polymer having one or two monodisperse oligopeptide
end groups blended with a hydrophobic, flexible polystyrene polymer
having one or two monodisperse oligopeptide end groups.
9. The polymer blend according to claim 1, in the form of a
shape-persistent thermoplastic elastomer.
10. The polymer blend according to claim 1, wherein each of said
oligopeptide end groups of said first oligopeptide-terminal polymer
component and said at least one additional polymer component is the
same.
11. The polymer blend according to claim 10, wherein each of said
oligopeptide end groups of said first oligopeptide-terminal polymer
component and said at least one additional polymer component has
the same 2 amino acid repeating units beyond its terminal amide
group.
12. The polymer blend according to claim 1, comprising
interpenetrating supramolecular polymer networks in which two or
more specific supramolecular interactions result in the formation
of two or more independent, interpenetrating supramolecular
networks with different transition temperatures, that is,
deaggregation temperatures.
13. A vibration damping material comprising the polymer blend
according to claim 12.
14. The vibration damping material according to claim 13, being a
composite material including one or more of the following: a
plasticizer; and a reinforcing filler comprising carbon fibre,
carbon black, or silica particles.
15. The vibration damping material according to claim 13, in a form
adapted to reduce vibration within a vehicle, the form being a pad
or other layer which can be interposed between members of the
vehicle subject to vibration.
16. A vehicle which includes the vibration damping material
according to claim 13.
17. The vehicle according to claim 16, which is a motor vehicle or
an aerospace vehicle.
18. A method of vibration damping which involves use of the polymer
blend according to claim 1 upon or within a structure or a
vehicle.
19. A polymer blend, comprising: at least one hydrophobic, flexible
polymer having a glass transition temperature below 20.degree. C.
and only one monodisperse oligopeptide end group, the monodisperse
oligopeptide end group having 1 to 5 amino acid repeating units;
and at least one hydrophobic, flexible polymer having a glass
transition temperature below 20.degree. C. and two monodisperse
oligopeptide end groups.
20. The polymer blend according to claim 19, wherein said at least
one hydrophobic, flexible polymer having only one monodisperse
oligopeptide end group and said at least one hydrophobic, flexible
polymer having two monodisperse oligopeptide end groups comprise
the same type of polymer segment.
Description
CROSS REFERENCE TO RELATED APPLICATIONS
[0001] This application is a continuation of co-pending U.S.
application Ser. No. 14/646,950 which is a 371 National Stage of
International Application No. PCT/EP2013/074793, filed Nov. 26,
2013, which claims the benefit under 35 USC 119 of United Kingdom
Application No. GB 1221246.0, filed Nov. 26, 2012.
FIELD OF THE INVENTION
[0002] This invention relates to polymer blends, in particular to
blends of polymers with oligopeptide end groups (named
oligopeptide-terminal polymers in the following), and in particular
to blends of oligopeptide-terminal polyisobutylenes, polyisoprenes,
polystyrenes and copolymers of the aforementioned polymers. It also
relates to the use of materials from such blends, for example as
reinforced, shape-persistent thermoplastic elastomers, as vibration
damping materials, as self-healing materials and in other
applications.
[0003] In typical thermoplastic materials, the polymer molecules
start to flow along each other upon heating, so that the material
can be macroscopically deformed. If this already happens at room
temperature, albeit very slowly, the terms "creep" or "cold flow"
are typically deployed. Materials exhibiting creep are not
form-stable; they may be useful in certain applications but
inadequate for others where shape-persistence is important.
[0004] In typical elastomers (rubbers), the polymer molecules are
linked together by chemical (covalent) bonds to form a
three-dimensional, covalently linked network of chains with loops
and holes. The chain segments cannot move anymore, and on trying to
deform the material macroscopically, it will retain its original
shape upon release of the deforming force (this is what makes the
material a rubber). As covalent bonds only break when the material
is destroyed at high temperatures, the covalent networks are
irreversibly formed and stable.
[0005] High molecular weight polyisobutylene is widely used
commercially as a thermoplastic material. It shows elastomeric
behaviour to a certain degree due to chain entanglement but, unless
covalently cross-linked ("butyl rubber") shows a large degree of
creep.
[0006] Interpenetrating polymer networks (IPNs) are elastomers that
contain two (or more) independent, mechanically interlocked
(interpenetrating) three-dimensional covalent networks consisting
of different types of polymers. Firstly, one type of covalent
network is formed from the first type of polymer chains by
covalently connecting (cross-linking) them by a chemical reaction.
Then the second type of polymer chains that fill the loops and
holes of the first network are cross-linked by a chemical reaction.
IPNs are useful as or in vibration damping materials. Without
wishing to be bound by theory, we postulate that by subjecting such
an IPN to a vibration of a certain frequency, only one of the two
independent networks will start to respond (assuming the two types
of polymers have different mechanical properties, e.g., stiffness)
so that at the molecular level, the segments of this network
apparently start to move along the segments of the other network,
which dissipates energy and thus damps the vibration.
[0007] Mechanical vibration damping is highly desirable for
structures, in order to maintain their stability, performance and
durability. Vibration damping is typically achieved by either
"active damping", using sensors and actuators such as piezoelectric
devices, or "passive damping" by means of materials that dissipate
vibrational energy. Due to their viscoelastic properties, polymer
materials are useful passive damping materials. Since the "loss
factor" tan .delta. (defined as the ratio of loss and storage
modulus, such as G''/G' in shear mode) as a measure of the
material's intrinsic damping properties attains its highest values
around the material's glass transition temperature, factors that
control the latter are an important aspect of research in the field
of damping materials. For instance, small molecule additives have
successfully been used as plasticizers, in order to broaden the
glass transition region. "Interpenetrating polymer networks"
(IPNs), on the other hand, often feature a superposition of the
damping properties of the constituting homopolymers and, thus,
offer opportunities for materials with excellent damping
performance over large temperature or frequency ranges. Finally,
the incorporation of inorganic nanostructures or fibres as
reinforcing fillers typically leads to a raising and broadening of
the loss factor maximum around the glass transition temperature of
the polymer matrix. This effect is particularly pronounced for
fibres with a low aspect ratio such as carbon whiskers or short
microfibers. Moreover, it was demonstrated, that the interface
properties between the matrix and the reinforcement strongly
influence the energy dissipation mechanisms, with intermediate
"interface friction" providing increased damping.
[0008] In contrast to true elastomers described above,
"supramolecular materials" are an example of so-called
"thermoplastic elastomers" (TPEs). In TPEs, the polymer chains are
cross-linked (as per elastomers), but the network points
(cross-links) are not covalent bonds (unlike elastomers). Instead,
the network points are "weak" non-covalent bonds (secondary bonds,
supramolecular interactions, such as electrostatic interactions,
hydrogen bonds, dipolar interactions, or Van der Waals
interactions) that create stable network points but are weak enough
to be reversibly broken/re-formed upon heating/cooling. This means
that the material is a three-dimensional network and behaves like
an elastomer at temperatures below the critical temperature where
the supramolecular interactions start to be disrupted. Above this
temperature, the network is broken, the polymer chains can flow,
and so the material becomes thermoplastic and can be
macroscopically deformed. This is advantageous because it enables
processing the material into any desired shape, and also from a
sustainability standpoint allowing a measure of re-use and
recycling.
[0009] In some cases, the reversible disruption/re-formation of the
networks can even dynamically happen at room temperature (or under
mild heating or mechanical treatment). In this case, the materials
may show a degree of "self-healing". This means that, by
macroscopically damaging as by cutting for example, the material,
the damage may heal just by itself, because molecules from both
sides of the crack apparently can "leave" their network and reform
a network with molecules from the other side.
[0010] In the prior art document WO 2011/045309 (BASF SE)
individual oligopeptide-terminal polymers as fall within definition
(a) herein and (b) herein (e.g. (a) monofunctional oligopeptide
terminal polymers, and (b) difunctional oligopeptide terminal
polymers) have been disclosed along with their preparation and some
postulated uses, but of an entirely different nature. However, in
that prior patent, there is no disclosure or suggestion of blends
of different oligopeptide-terminated polymers or of blends of
oligopeptide-terminated polymers with other polymers such as medium
or high molecular weight PIB. Still less, is there any suggestion
of the particular advantageous properties, discovered now,
attributable to such blends.
SUMMARY OF THE INVENTION
[0011] The invention provides, in one aspect, a polymer blend
comprising: at least two (a) monofunctional oligopeptide terminal
polymers, or at least two (b) difunctional oligopeptide terminal
polymers.
[0012] The invention also provides a polymer blend comprising: at
least one polymer (a) as defined above, together with at least one
polymer (b) as defined above.
[0013] The invention further provides, in an alternative aspect, a
polymer blend comprising: at least two (a) monofunctional
oligopeptide-terminated hydrophobic flexible polymers, or at least
two (b) difunctional oligopeptide-terminated hydrophobic flexible
polymers, or at least one of (a) together with at least one of (b)
as defined above.
[0014] In a further alternative aspect, the invention provides a
polymer blend comprising: at least two (a) monofunctional
oligopeptide-terminated hydrophobic polymers, or at least two (b)
difunctional oligopeptide-terminated hydrophobic polymers, or at
least one of (a) together with at least one of (b) as defined
above.
[0015] In some embodiments, the polymer blend may further comprise
at least one (c) non-functionalised polymer, such as an isobutylene
polymer. The non-functionalised polymer may have, for example, a
molecular weight exceeding 10,000.
[0016] The invention also extends to a polymer blend comprising: at
least one polymer (c) as defined above, and at least one polymer
(a) as defined above and/or at least one polymer (b) as defined
above.
[0017] In some embodiments, polymers (a), (b) and/or (c) comprise
repeating units selected from isobutylene, butadiene, siloxane,
acrylate, fluoropolymer, isoprene or styrene units. In some
embodiments, the polymer segment of polymer (a) and/or (b) and/or
(c), may be selected from the group consisting of:
poly(isobutylene-co-isoprene), polyisoprene, polybutadiene, a
.quadrature.polysiloxane (in particular poly(dimethylsiloxane)), a
polyacrylate (in particular poly(methyl acrylate) or poly(butyl
acrylate)), poly(ethylene-co-butylene), hydrogenated
poly(isoprene), hydrogenated poly(butadiene), or a fluoropolymer
(in particular poly(tetrafluoroethylene), and
poly(tetrafluoroethylene-co-ethylene).
[0018] Polymers (a) and (b) may, in some embodiments, comprise the
same type of polymer segment. Where present, polymer (c) may, in
some embodiments, comprise the same type of polymer segment as
polymer (a) and/or (b).
[0019] In some particular embodiments, the blend comprises
oligopeptide-terminated isobutylene polymer with
oligopeptide-terminated styrene polymer.
[0020] Polymers (a), (b) and/or (c) may comprise, for example,
flexible hydrophobic polymers.
[0021] In some embodiment, the oligopeptide-terminated polymer (a)
may have from 0 to 5 (for example 2 to 5) peptide units beyond its
terminal amide group and/or polymer (b) may have, at each end, 0 to
5 (for example 2 to 5) peptide units beyond its terminal amide
groups.
[0022] The oligopeptide moiety of polymer (a) and/or (b) may
comprise, for example, L-alanine units. In some embodiments, the
oligopeptide moiety consists only of such units.
[0023] The polymer segment of polymer (a) and/or (b) and/or, if
present, (c), may, in some embodiments, be a hydrophobic polymer
with a glass transition temperature below 20.degree. C.
[0024] Polymer blends according to the invention may, for example,
be in the form of shape-persistent thermoplastic elastomers.
[0025] In some embodiments, the polymer blend may comprise
interpenetrating supramolecular polymer networks in which two or
more specific supramolecular interactions result in the formation
of two or more independent, interpenetrating supramolecular
networks with different transition (deaggregation)
temperatures.
[0026] The invention also extends to vibration damping materials
comprising the polymer blends of the invention.
[0027] Some embodiments of vibration damping material comprise: (i)
monofunctional oligopeptide terminal polyisobutylene, having from 2
to 5 peptide units beyond its terminal amide group; and (ii)
monofunctional oligopeptide terminal polystyrene, having from 2 to
5 peptide units beyond its terminal amide group.
[0028] The vibration damping material may, in some embodiments,
additionally comprise: (iii) a high molecular weight
non-functionalised polymer.
[0029] Further particular embodiments of vibration damping material
according to the invention include, by way of example-- [0030]
A.--a material comprising: (i) monofunctional oligopeptide terminal
polyisobutylene, having from 2 to 5 peptide units beyond its
terminal amide group; and (ii) non-functionalised polyisobutylene;
wherein the monofunctional oligopeptide terminal polyisobutylene
may, for example, have 2 peptide units beyond its terminal amide
group; and [0031] B.--a material comprising: (i) monofunctional
oligopeptide terminal polyisobutylene, having from 0 to 5 peptide
units beyond its terminal amide group; (ii) difunctional
oligopeptide terminal polyisobutylene, having from 0 to 5 peptide
units beyond its terminal amide group; and (iii) non-functionalised
polyisobutylene; wherein each of oligopeptide terminal
polyisobutylene polymers (i) and (ii) has the same number of
peptide units beyond its terminal amide group; wherein each of
oligopeptide terminal polyisobutylene polymers (i) and (ii) may,
for example, have 2 peptide units beyond its terminal amide
group.
[0032] Vibration damping materials according to the invention may,
for example, be composite materials including one or more other
components, such as one or more of the following: a plasticizer or
a reinforcing filler (such as carbon fibre, carbon black, silica
particles).
[0033] Vibration damping materials of the invention may be used,
for example, in a form adapted to reduce vibration within a
vehicle, such as in the form of a pad or other layer which can be
interposed between members of such vehicle subject to vibration.
The invention also extends to vehicles which incorporate such
damping materials, for example motor vehicles and aerospace
vehicles.
[0034] In alternative aspects, the invention provides the use of
any of the polymer blends of the invention as a damping material
and to a method of vibration damping which involves the use of any
of the polymer blends of the invention upon or within a structure
such as a vehicle.
[0035] Moreover, the present invention also embraces polymer blends
of monofunctional and/or difunctional oligopeptide-terminated
polymers and/or the corresponding non-functional, higher molecular
weight polymers of flexible hydrophobic polymers other than
polyisobutylene, such as polyisoprene, polybutadiene,
polysiloxanes, polyacrylates, or fluoropolymers.
[0036] The present invention embraces, for example, blends of
monofunctional oligopeptide-terminated polyisobutylenes with
oligopeptide segments comprising 0-5 amino acid repeating units
(designated for convenience M0-M5) and/or their difunctional
analogues (similarly designated for convenience as D0-D5) and/or
oligopeptide-terminated polystyrene with oligopeptide segments
comprising 0-5 amino acid repeating units (designated for
convenience S0-S5) and/or oligopeptide-terminated polymers other
than polyisobutylene/polystyrene (such as polybutadiene,
polyisoprene, .quadrature.hermogravime, polyacrylates,
polymethacrylates, or copolymers of any of the aforementioned
polymers including polyisobutylene and polystyrene) with
oligopeptide segments comprising 0-5 amino acid repeating units
(designated for convenience P0-P5) and/or (optionally high
molecular weight) non-functionalised polyisobutylenes of different
grades (molecular weights) and/or flexible hydrophobic polymers
other than polyisobutylene, such as polybutadiene, polyisoprene,
polysiloxanes, polyacrylates, or fluoropolymers, or copolymers of
any of the aforementioned polymers including polyisobutylene and
polystyrene.
[0037] The above blend(s), if required, can be used as a precursor
blend to be admixed with other non-functionalised, higher molecular
weight hydrophobic flexible polymer such as polyisobutylene, to
form the preferred thermoplastic elastomer polymer blends of the
invention. Alternatively, just one polymer (a) or (b) as defined
herein may be admixed with the non-functionalised, higher molecular
weight hydrophobic flexible polymer, such as polyisobutylene.
[0038] These blends result in materials with novel properties
because the oligopeptides deployed in the present invention form
aggregates by hydrogen-bonding between the peptide groups; they are
chiral; and their (exact) length determines what kind of aggregates
form; and this length-dependence is not only selective but even
specific (self-sorting) so that different types of aggregates can
persist in mixtures. The present invention also embraces polymer
blends using oligopeptide terminal groups based on other amino
acids apart from those specifically exemplified herein.
Furthermore, whilst the present blends (materials) may be used on
their own, they may also be used as the matrix material for a
composite, that is, a reinforcing filler may be added into the
present blends (such as carbon fibres, or carbon black, or silica
particles). This may result in stronger, stiffer materials with the
same beneficial damping properties.
[0039] Further optional and preferred features are to be found
amongst the sub claims herein:
[0040] (1) Some preferred embodiments of these blends are materials
that are "Inherently Reinforced Thermoplastic Elastomers" in which
nanostructures formed by the aggregation of the oligopeptide
segments act as reinforcing fillers. In particular, these materials
are technologically advantageous because they are thermoplastic
materials with increased mechanical moduli, low creep, and good
thermal processability compared to the corresponding high molecular
weight polymers alone. See in particular the manifold examples
herein.
[0041] (2) Alternatively, further preferred embodiments of these
blends result in the formation of novel "Interpenetrating
Supramolecular Networks". In particular, these materials are
technologically advantageous because they have excellent mechanical
vibration damping properties. See in particular the manifold
examples herein.
[0042] (3) Still, further preferred embodiments of these blends can
be deployed as materials with self-healing properties.
[0043] (4) Lastly, further embodiments of the invention are in the
physical form of composites which comprise the aforesaid polymer
blends as well as reinforcing fillers. Such composites can also be
used as damping materials with, for example, higher strength,
stiffness, yet similar damping properties, showing a potential use
of the present polymer blends materials as matrix materials for
composites containing other desired or required components. These
composites can also show self-healing properties.
[0044] The "Inherently Reinforced Thermoplastic Elastomer"
embodiments can be derived from mixtures of molecules with
oligopeptide termini of the same type and length. These embodiments
are examples of TPEs (in regard to their properties) but are novel
because they, due to the molecularly defined oligopeptide end
groups, already form networks at very short segment lengths of
these end groups, so the inherent properties of the employed
polymer (polyisobutylene) do not alter too much. At the same time,
they are "reinforced" by the nanostructures (the tapes and fibrils)
formed by the oligopeptides when they aggregate, similar to
reinforcing a polymer with a filler (e.g., carbon fibres) to make a
high-performance composite. As a result, we obtain materials with
mechanical properties (moduli) matching or exceeding those of even
high molecular weight polyisobutylene, although we prefer to use
very low molecular weight material. The present materials can be
employed for the same applications as high molecular weight
polyisobutylene but can be thermally processed more easily, have
better mechanical properties, show less creep, and are better from
a sustainability standpoint (recycling elastomers).
[0045] Some preferred embodiments of the aforesaid "Inherently
Reinforced Thermoplastic Elastomer" embodiments can be derived from
mixtures of oligopeptide-terminated flexible polymers such as
polyisobutylene with oligopeptide-terminated glassy polymers such
as polystyrene or its copolymers. In particular, these embodiments
are novel because the phase segregation of the immiscible polymers
competes with the formation of the oligopeptide aggregates in the
blends. Different from the microphase segregation observed in
typical block copolymers with domain sizes on the order of tens of
nanometers and above, this results in separated domains of the
immiscible flexible and amorphous polymers with diameters on the
order of nanometers to tens of nanometers that are, in addition,
connected by the oligopeptide aggregates. This creates a network of
nanoscopic oligopeptide aggregates and nanoscale glassy domains
with a particularly reinforcing effect, even at small weight
fractions of the oligopeptide-terminated glassy polymers.
[0046] The "Interpenetrating Supramolecular Networks" embodiments
are also novel and are obtained from mixtures of molecules with
oligopeptide termini of different type (oligopeptide sequence) or
length, including non-functionalized, higher molecular weight
polymers. In these materials, there are two independent networks
(similar to IPNs), but both are formed by cross-links that are
non-covalent, weak, secondary bonds (e.g. hydrogen bonding). What
this requires is that the network formation relies on two
"specific" (self-sorting) supramolecular interactions that can form
without interfering with each other. In the materials provided by
this invention, both types of networks rely on the same type of
supramolecular interactions (that is, hydrogen-bonding between the
peptide functions) between two just differently long oligopeptides.
The resulting networks can be either a supramolecular,
hydrogen-bonded network of difunctional molecules (such as D0-D5);
or a "percolation network" of tapes and fibrils formed by
hydrogen-bond-driven aggregation of monofunctional molecules
(M2-M5); or a percolation network of aforementioned tapes and
fibrils with "hard" domains formed from oligopeptide-terminated
glassy polymers (S2-S5 or other P2-P5); or an entanglement network
formed from the high molecular weight polymers.
[0047] In any case, these embodiments share structural aspects of
IPNs (two mechanically interlocked networks) and thermoplastic
elastomers (reversibility of network formation). These embodiments
thus extend upon regular supramolecular networks and other examples
of TPE (that contain no interpenetrating networks) and IPNs (that
are formed from covalent networks), and as a result can function as
excellent high performance damping materials whilst simultaneously
allowing beneficial thermoplastic processing (into different
shapes). Their processing is also flexible because it can be
effected either above the temperature where the first network
melts, or above the temperature when also the second network melts,
with different results. Specifically, it is possible to heat above
the melting temperature of both networks, process the material into
the desired shape (by injection moulding, extrusion, or other
required technique), then cool below the melting temperature of one
network, let it form, then cool below the melting temperature of
the second network to have it formed. In this way, an
interpenetrating network can be created just by processing, without
using additional chemical reaction steps (as is required for an
IPN).
[0048] Some preferred embodiments of the aforesaid
"Interpenetrating Supramolecular Networks" materials may also show
self-healing properties. One network may hold the material in place
while the other (the weaker) network may dynamically break/re-form,
either spontaneously or by heating it above that network's melting
temperature, or by "mechanical treatment" (exposing the material
to, e.g., a mechanical vibration).
[0049] Advantageous properties of preferred embodiments of the
present polymer blends include:
[0050] (1) Behaviour as reinforced thermoplastic elastomers, that
is, rubbers with properties similar to or better than high
molecular weight polyisobutylene alone, but yet which can be
processed thermoplastically by melting upon heating. They are
`reinforced` by the oligopeptide nanostructures, resulting in
improved mechanical moduli ("strength") and behave as materials
with no or low "creep" (cold flow; that is, they keep their shape
at room temperature). Commercial applications and industrial uses
can mirror those of regular PIB (e.g., barrier materials) but their
low creep and thermoplastic processing are unexpected
advantages,
[0051] (2) Behaviour as materials with extremely large loss factors
(loss tangents tan delta) over large frequency and/or temperature
ranges, i.e. apparent molecular level properties approaching
liquids (at those mechanical frequencies and temperatures) although
they are in fact solids. This finds potential application in matrix
materials for self-healing applications (materials that can cure
mechanical damages `themselves`),
[0052] (3) Significantly, excellent high-performance vibration
damping materials. Vibration damping uses being abundant in
automotive and aerospace engineering whereby embodiments of the
present blends find substantial application (even in their
non-optimized state, some embodiments of the present invention
already match or out-perform optimised multi-component commercial
formulations.
[0053] The materials disclosed in the present invention implement
useful vibration damping properties in a novel way. Due to the
formation of the aforesaid "interpenetrating supramolecular
networks", these materials themselves combine the beneficial
effects (with respect to energy dissipation upon mechanical
excitation as needed for vibration damping) of low aspect ratio
reinforcing fillers, interpenetrating networks, and low molecular
weight poly(isobutylene) plasticizers. They, therefore, yield
shape-persistent materials with excellent energy dissipation and
damping properties over large frequency and temperature ranges,
without the use of additional components such as additional fillers
or plasticisers.
BRIEF DESCRIPTION OF THE DRAWINGS
[0054] In order that the invention may be further described, more
easily appreciated and readily carried into effect by those skilled
in the art, reference will now be made to embodiments by way of
non-limiting example only and with reference to the accompanying
drawings, in which:
[0055] FIGS. 1A, 1B and 1C represent schematic illustrations of the
selective self-assembly of the monofunctional
oligo(L-alanine)-modified poly(isobutylene)s M0-M5 (n=0-5;
x.apprxeq.20) and the corresponding difunctional derivatives D0-D5
(n=0-5; x 20), FIG. 1D represents a schematic illustration in which
the coexistence of these nanostructures in blends of molecules with
matching oligopeptide termini results in `inherently reinforced
thermoplastic elastomers`, and FIG. 1E represents a schematic
illustration in which the coexistence of these nanostructures in
blends of molecules with different oligopeptide termini results in
`interpenetrating supramolecular networks`,
[0056] FIGS. 2A, 2B, 2C and 2D represent amide A and amide I
regions of the solid-state infrared (IR) spectra of bulk samples of
PIB-Ala.sub.n-Ac M0-M5 as well as Ac-Ala.sub.n-PIB-Ala.sub.n-Ac
D0-D5, and FIGS. 2E and 2F represent corresponding plots of the
position of the global maxima of the amide A and amide I
absorptions,
[0057] FIGS. 3A, 3B, 3C and 3D represent amide A and amide I
regions of the solution-phase infrared (IR) spectra of samples of
PIB-Ala.sub.n-Ac M0-M5 as well as Ac-Ala.sub.n-PIB-Ala.sub.n-Ac
D0-D5 in dilute solution in tetrachlorethane, and FIGS. 3E and 3F
represent corresponding plots of the position of the global maxima
of the amide A and amide I absorptions,
[0058] FIG. 4A represents peak deconvolution of the amide I regions
of samples of PIB-Ala.sub.n-Ac M0-M5 in bulk, FIG. 4B represents
peak deconvolution of the amide I regions of samples of
Ac-Ala.sub.n-PIB-Ala.sub.n-Ac D0-D5 in bulk, FIG. 4C represents
peak deconvolution of the amide I regions of samples of
PIB-Ala.sub.n-Ac M0-M5 in dilute solution in tetrachlorethane, and
FIG. 4D represents peak deconvolution of the amide I regions of
samples of Ac-Ala.sub.n-PIB-Ala.sub.n-Ac D0-D5 in dilute solution
in tetrachlorethane,
[0059] FIG. 5 shows atomic force microscopy (AFM) images of M1-M5
spin-coated from tetrachlorethane solution onto either SiO.sub.2 or
HOPG substrates,
[0060] FIGS. 6A and 6B show thermogravimetric analysis of M0-M5 and
D0-D5 as well as the parent poly(isobutylene)s PIB-NH.sub.2 and
H.sub.2N--PIB-NH.sub.2, FIG. 6C represents differential scanning
calorimetry of M0-M5 as well as the parent poly(isobutylene)
PIB-NH2, and FIG. 6D represents differential scanning calorimetry
of D0-D5 as well as the parent poly(isobutylene) H2N-PIB-NH2,
[0061] FIGS. 7A and 7B represent amide I regions of the
temperature-dependent solid-state IR spectra of M2 and D2,
[0062] FIGS. 8A, 8B and 8C show rheological dynamic frequency sweep
experiments at 25.degree. C. of unmodified PIB of different
molecular weights (1200 for PIB-NH.sub.2, 2500 for
H.sub.2N-PIB-NH.sub.2, 35'000, 75'000, 200'000, 425'000), showing
a) storage moduli G', b) loss moduli G'', and c) viscosity 2 and
D2,
[0063] FIGS. 9A, 9B, 9C, 9D, 9E and 9F show rheological dynamic
frequency sweep experiments at 25.degree. C. of PIB-Ala.sub.n-Ac
M0-M5 and Ac-Ala.sub.n-PIB-Ala.sub.n-Ac D0-D5 as well as the parent
poly(isobutylene)s 1 and 4, showing a,d) storage moduli G', b,e)
loss moduli G'', and c,f) viscosity,
[0064] FIGS. 10A, 10B and 10C show rheological dynamic frequency
sweep experiments at 25.degree. C. of different binary blends
M2/D2, showing a) storage moduli G', b) loss moduli G'', and c) and
viscosity,
[0065] FIGS. 11A, 11B, 11C and 11D shows a comparison of storage
(G') and loss moduli (G'') at 1 rad/s of a) unmodified PIB as a
function of molecular weight, c,d) M0-M5 and D0-D2 as a function of
hydrogen-bonding sites per end group, and b) of different blends of
M2/D2 as a function of composition,
[0066] FIGS. 12A and 12B represents amide A, amide I and amide II
regions of the solid-state infrared (IR) spectra of M2, D2, as well
as the complete series of M2/D2 blends (Examples 1-6),
[0067] FIGS. 13A and 13B represent amide I regions of the
temperature-dependent solid-state IR spectra of the blend M2/D2 9:1
(Example 3), and FIG. 13C shows the parent PIB amine was a viscous
liquid, M2 was a sticky solid, D2 a brittle powder, the blend M2/D2
9:1 (Example 3) was an `inherently reinforced` thermoplastic
elastomer,
[0068] FIGS. 14A, 14B and 14C show temperature-dependent shear
rheology of the blend M2/D2 9:1 (Example 3) and FIGS. 14D, 14E and
14F show temperature-dependent shear rheology of PIB (MW 200'000)
as a reference material,
[0069] FIGS. 15A, 15B and 15C show Temperature-dependent shear
rheology of the blend M2/D1 1:4 (Example 7), compared to pure M2,
pure D1, as well as the blend M2/D2 9:1 (Example 3) including plots
of a) storage modulus G', b) loss modulus G'', and c) viscosity,
and FIG. 15D represents differential scanning calorimetry of M2/D1
1:4,
[0070] FIGS. 16A, 16B and 16C show rheological time-temperature
superposition (TTS) master curves of D1, M2, and M2/D1 1:4 (Example
7),
[0071] FIGS. 17A and 17B show amide A, amide I and amide II regions
of the solid-state infrared (IR) spectra of solution phase IR
spectra of samples in tetrachlorethane and FIGS. 17C and 17D show
amide A, amide I and amide II regions of the solid-state infrared
(IR) spectra of solid state IR spectra of bulk samples of PIB (MW
75'000), PIB (MW 35'000), M2, D2, as well as different binary and
ternary blends of M2 and D2 in PIB (MW 75'000) or PIB (MW 35'000)
(Examples 8-10),
[0072] FIG. 18A represents differential scanning calorimetry of M2,
D2 and their binary and ternary blends M2/PIB (MW 35'000) 5:5
(Example 9) and M2/D2/PIB (Mw 35'000) 4:1:5 (Example 10), and FIG.
18B shows thermogravimetric analysis of M2, D2 and their binary and
ternary blends M2/PIB (MW 35'000) 5:5 (Example 9) and M2/D2/PIB (Mw
35'000) 4:1:5 (Example 10) as well as for PIB-NH2 and
H2N-PIB-NH2,
[0073] FIG. 19 shows atomic force microscopy (AFM) height (left)
and phase (right) images of M2 and M2/PIB (MW 35'000) 5:5 (Example
9),
[0074] FIGS. 20A, 20B, 20C, 20D, 20E and 20F show rheological
time-temperature superposition (TTS) master curves of M2/PIB (MW
75'000) 5:5 (Example 8), M2/PIB (MW 35'000) 5:5 (Example 9),
M2/D2/PIB (MW 35'000) 4:1:5 (Example 10),
[0075] FIGS. 21A, 21B and 21C show an experimental setup for a
random vibration modal analysis test on a sandwich structure
comprising a damping layer, and
[0076] FIGS. 22A, 22B, 22C, 22D and 22E represent Lissajous curves
obtained from oscillatory shear stress-strain test, performed on a
rheometer, and FIGS. 22F and 22H represent plots of the logarithm
of the dissipated energies relative to the logarithm of the strain
applied at -45.degree. C. during an oscillatory shear stress-strain
test, performed on a rheometer.
[0077] FIGS. 23A and 23B show atomic force microscopy (AFM) height
images and differential scanning calorimetry measurements of the
blend of M2 with unmodified polystyrene (9:1) and the blend M2/S2
9:1 (Example 11); FIG. 23C shows atomic force microscopy (AFM)
height images and FIG. 23D shows differential scanning calorimetry
measurements of the blend M3 with unmodified polystyrene (9:1) or
the blend M3/S3 9:1 (Example 12).
[0078] FIGS. 24A and 24B show storage moduli G' and loss moduli G''
determined by rheological dynamic frequency sweep experiments at
25.degree. C. of M2, S2, the blend of M2 with unmodified
polystyrene (9:1) and the blend M2/S2 9:1; FIGS. 24C and 24D show
storage moduli G' and loss moduli G'' determined by rheological
dynamic frequency sweep experiments at 25.degree. C. of M3, S3, the
blend of M3 with unmodified polystyrene (9:1) and the blend M3/S3
9:1.
[0079] FIG. 25 shows a rheological time-temperature superposition
master curve of the ternary blend M3/S3/PIB (MW 35'000) 9:3:12
(Example 13).
DETAILED DESCRIPTION OF EXAMPLES
[0080] As shown in the drawings and referring in particular to
FIGS. 1A, 1B, 1C and 1D:
[0081] FIGS. 1A, 1B and 1C provide Schematic illustration of the
selective self-assembly of the monofunctional
oligo(L-alanine)-modified poly(isobutylene)s M0-M5 (n=0-5; x 20)
and the corresponding difunctional derivatives D0-D5 (n=0-5; x 20)
into small hydrogen-bonded aggregates, flexible single .beta.-sheet
tapes, or rigid stacked .beta.-sheet fibrils. The coexistence of
these nanostructures in blends of molecules with different
oligopeptide termini resulted in `inherently reinforced
thermoplastic elastomers` (see FIG. 1D) or `interpenetrating
supramolecular networks` (see FIG. 1E).
[0082] Differently from all previous examples of supramolecular
networks, the aggregation of the oligopeptide-terminated polymers
that constitute the basis of the present invention that comprise
chiral and monodisperse (molecularly defined) oligopeptides as
hydrogen-bonded ligands results in a highly selective formation of
small hydrogen-bonded aggregates from compounds with short
oligopeptides (such as M0-M1, D0-D1), flexible single .beta.-sheet
tapes from compounds with medium-size oligopeptides (such as M2-M3,
D2-D3), or rigid stacked .beta.-sheet fibrils from compounds with
longer oligopeptides (such as M4-M5, D4-D5), because the helical
conformation of single oligopeptide .beta.-strands, the induced
helical twisting of .beta.-sheets, and finally the number of
stacked .beta.-sheets are intimately interrelated. This
length-dependent self-assembly is even "self-sorting", that is,
specific in the sense that the different nanostructures obtained
from different oligopeptide segments coexist in bulk. It is this
particular feature that has enabled us to tailor the
thermomechanical properties of the blends. Thus, mixtures of
molecules with "matching" oligopeptide termini (identical
oligopeptide length and amino acid sequence) gave rise to
thermoplastic elastomers that were "inherently reinforced" with
.beta.-sheet tapes or fibrils. By contrast, blends of derivatives
with "non-matching" oligopeptide termini (different oligopeptide
length or amino acid sequence, including non-functionalized
polymers) formed novel "interpenetrating supramolecular networks".
It is worth noting that in both cases, network formation allows for
dynamic network reorganization processes and may give rise to
self-healing or thermoresponsive materials. In this regard,
polyisobutylene soft segments have proven to be of high interest,
due to their conformational dynamics and resulting macroscopic
properties. See in particular the manifold examples. Further
examples of blends of non-functionalised and
oligopeptide-terminated derivatives of flexible and hydrophobic
polymers, such as polyisoprene, polybutadiene, polyacrylates,
polysiloxanes, or fluoropolymers, share the same structural
features and properties and are embraced by the current
invention.
[0083] Referring in particular to FIGS. 2A-5:
[0084] FIGS. 2A, 2B, 2C and 2D provide Amide A and amide I regions
of the solid-state infrared (IR) spectra of bulk samples of
PIB-Ala.sub.n-Ac M0-M5 as well as Ac-Ala.sub.n-PIB-Ala.sub.n-Ac
D0-D5, as well as FIGS. 2E and 2F provide corresponding plots of
the position of the global maxima of the amide A and amide I
absorptions as a function of the number of alanine repeating units
n revealed that M2-M5 and D2-D5 exhibited a single amide A
absorption at 3270-3276 cm.sup.-1, a strong and sharp amide I
absorption at 1625-1627 cm.sup.-1 (half-height width.apprxeq.16-17
cm.sup.-1), and a sharp secondary absorption at 1687-1695
cm.sup.-1, all consistent with the presence of highly ordered
antiparallel .beta.-sheet structures. See FIGS. 4A and 4B for peak
deconvolutions.
[0085] FIGS. 3A, 3B, 3C and 3D provide Amide A and amide I regions
of the solution-phase infrared (IR) spectra of samples of
PIB-Ala.sub.n-Ac M0-M5 as well as Ac-Ala.sub.n-PIB-Ala.sub.n-Ac
D0-D5 in dilute solution in tetrachlore thane, as well as FIGS. 3E
and 3F provide corresponding plots of the position of the global
maxima of the amide A and amide I absorptions as a function of the
number of alanine repeating units n revealed that, in solution,
M0-M2 and D0-D2 remained non-aggregated. A sharp transition was
then observed for longer oligopeptides; M3-M5 as well as D3-D5
exhibited a single amide A absorption at 3271-3273 cm-1, a strong
and sharp amide I absorption at 1624-1625 cm-1 (half-height
width.apprxeq.14-17 cm-1), and a sharp secondary absorption at
1690-1695 cm-1, all consistent with the presence of highly ordered
antiparallel .beta.-sheet structures in solution. See FIGS. 4C and
4D for peak deconvolutions.
[0086] Peak deconvolution of the amide I regions of samples of
PIB-Ala.sub.n-Ac M0-M5 as well as Ac-Ala.sub.n-PIB-Ala.sub.n-Ac
D0-D5 are shown in bulk (see FIGS. 4A and 4B) and in dilute
solution in tetrachlorethane (see FIGS. 4C and 4D); global maxima
labelled in blue; .beta.-sheet bands in red; predominant bands in
black. Although the peak fitting was started with the same number
of bands at approximately the same positions, limiting their width
to reasonable values in all cases, the results of the deconvolution
were still sensitive to the exact starting parameters and,
therefore, just served to obtain an estimate for the peak area
A1625 of the absorption bands at around 1625-1630 cm.sup.-1
relative to the total peak area AI, total of the amide I
absorption. The latter is a qualitative assessment for the relative
degree of aggregation. See FIG. 2d (main text) for a plot of
A1625/AI, total as a function of the number of alanine repeating
units n.
[0087] FIG. 5. Atomic force microscopy (AFM) images of M1-M5
spin-coated from tetrachlorethane solution onto either SiO.sub.2 or
HOPG substrates revealed the formation of fibrils for M5; mixtures
of fibrils and tapes for M4 on HOPG and mixtures of fibrils and
drop-like features on SiO.sub.2; no defined aggregates for M0-M3 on
SiO.sub.2, but long tapes for M3, short laterally aggregated tapes
for M2, and continuous films for M1 and M0 on HOPG.
[0088] The monofunctional compounds M0-M5 and the difunctional
compounds D0-D5 exhibited distinctly length-dependent aggregation
properties. According to IR spectroscopy, M4-M5 and D4-D5 gave rise
to highly ordered and strongly aggregated antiparallel .beta.-sheet
structures both in bulk and in solution. M2-M3 and D2-D3 were only
aggregated in bulk materials. The end groups in M0-M1 and D0-D1
were too short to induce .beta.-sheet formation either in the bulk
or in solution. Atomic force microscopy (AFM) imaging then
established a link to the corresponding nanoscopic morphologies for
the monofunctional derivatives M0-M5. Thus, rigid and many
micrometres long fibrils with diameters of a few nanometres were
observed for M5 and M4 on both highly oriented pyrolytic graphite
(HOPG) and SiO.sub.2 substrates. The dimensions of fibrils obtained
from M5 suggested that they were formed from 4-6 stacked
.beta.-sheet tapes. In the case of M4, the fibrils were formed from
2-4 stacked .beta.-sheet tapes, according to their cross-sections
determined by AFM imaging. M3 gave rise to long flexible fibrils or
tapes on HOPG that were thinner than those of M4 and exhibited an
epitaxial orientation with the substrate. In the case of M2, we
observed laterally aggregated tape-like features on HOPG with
lengths on the order of a few hundred nanometres. The epitaxial
orientation of the tape-like features from M2-M4 on HOPG as well as
their absence on SiO.sub.2 substrates suggested that they had not
already been present in solution but formed upon drying of the
sample on the AFM substrate, in agreement with the IR spectroscopic
results. Hence, our results prove that longer oligopeptides did not
only result in the expected increase in aggregation strength but
that superstructure formation was also affected, due to the
molecular chirality of an oligo(L-alanine) segment. We proved that
we selectively obtained rigid stacked .beta.-sheet fibrils from the
"longest" oligopeptides (n 4 alanine residues); single .beta.-sheet
tapes from "medium-size" oligopeptides (n=2-3 alanine residues) in
the bulk; and weak, undefined aggregates from short hydrogen-bonded
end groups (n=0-1 alanine residues).
[0089] Referring in particular to FIGS. 6A-7B:
[0090] FIGS. 6A and 6B providing Thermogravimetric analysis of
M0-M5 and D0-D5 as well as the parent poly(isobutylene)s
PIB-NH.sub.2 and H.sub.2N--PIB-NH.sub.2 revealed that complete PIB
depolymerization occurred at temperatures above 340.degree. C. in
all cases. Whereas derivatives M0-M2 and D0-D2 were stable up to
temperatures of at least 250.degree. C., compounds with longer
oligopeptides noticeably underwent a first stage of decomposition
already at temperatures of around 170-200.degree. C., tentatively
assigned to a degradation of the oligopeptides, supposedly by
ring-closing fragmentation. FIGS. 6C and 6D providing Differential
scanning calorimetry revealed that only M2 and D2 exhibited
reversible thermal transitions at 170.degree. C. (15.5 J/g) and
178.degree. C. (16 J/g), respectively. These could be assigned to
the `melting` (deaggregation) and `crystallization` (aggregation)
of the .beta.-sheet aggregates (see FIGS. 7A and 7B). Derivatives
with longer oligopeptides showed endothermic peaks at increasingly
high temperatures that were already in the range of or above their
decomposition temperatures and, accordingly, did not exhibit any
exothermic peaks upon cooling, except for M3 which exhibited a weak
exothermic peak at 238.degree. C. (3 J/g).
[0091] FIGS. 7A and 7B providing Amide I regions of the
temperature-dependent solid-state IR spectra of M2 and D2 revealed
that the position and intensity of the absorption bands associated
with .beta.-sheet secondary structures remained virtually unchanged
until at least 150.degree. C. but then rapidly decreased at
temperatures above 170.degree. C., proving that the reversible
transitions observed in DSC (see FIGS. 6A, 6B, 6C and 6D) were
associated to `melting` (deaggregation) and `crystallization`
(aggregation) of the .beta.-sheet aggregates.
[0092] Thermogravimetric analysis (TGA), differential scanning
calorimetry (DSC), and temperature-dependent solid-state IR
spectroscopy proved that M0-M2 and D0-D2 were straightforwardly
processable below their degradation temperature of 250.degree. C.
Moreover, M2 and D2 exhibited detectable reversible thermal
transitions at 170.degree. C. and 178.degree. C., respectively,
according to DSC, that solid state IR spectroscopy proved to be
associated with .beta.-sheet deaggregation.
[0093] Referring in particular to FIGS. 8-9:
[0094] FIGS. 8A, 8B and 8C provide Rheological dynamic frequency
sweep experiments at 25.degree. C. of unmodified PIB of different
molecular weights (1200 for PIB-NH.sub.2, 2500 for
H.sub.2N--PIB-NH.sub.2, 35'000, 75'000, 200'000, 425'000), showing
storage moduli G' (FIG. 8A), loss moduli G'' (FIG. 8B), and
viscosity |.eta.*| (FIG. 8C); the shear viscosity formally defined
as |.eta.*|=p.sub.21,0/(.gamma..sub.0.omega.) was used as a
calculated entity to compare both liquid and rubbery materials. The
.gamma. (% strain) values used for these experiments ranged from
3-30% and were selected to be in the linear regime of the
investigated materials. Depending on their molecular weight, the
materials undergo a transition in mechanical properties, from low
viscosity liquid behaviour with a zero-shear viscosity of 32 Pa s
for PIB-NH.sub.2 and 140 Pa s for H.sub.2N--PIB-NH.sub.2 to,
finally, rubbery behaviour for PIB (MW 200'000 and 425'000). Hence,
the latter two materials exhibited a frequency-dependent viscosity
|.eta.*| with a slope of -1, and both their storage (G') and loss
moduli (G'') at 1 rad/s (as a reference) were increased by 7 and 3
orders of magnitude, respectively, as compared to low molecular
weight PIB-NH.sub.2.
[0095] FIGS. 9A-9F providing Rheological dynamic frequency sweep
experiments at 25.degree. C. of PIB-Ala.sub.n-Ac M0-M5 and
Ac-Ala.sub.n-PIB-Ala.sub.n-Ac D0-D5 as well as the parent
poly(isobutylene)s 1 and 4, showing storage moduli G' (FIGS. 9A and
9D), loss moduli G'' (FIGS. 9B and 9E), and viscosity |.eta.*|
(FIGS. 9C and 9F); the shear viscosity formally defined as
|.eta.*|=p.sub.21,0/(.gamma..sub.0.omega.) was used as a calculated
entity to compare both liquid and rubbery materials. The .gamma. (%
strain) values used for these experiments ranged from 0.03-30% and
were selected to be in the linear regime of the investigated
materials. Depending on the number of hydrogen-bonding sites in the
series M0-M5, the materials showed a transition from moderately
viscous liquid behaviour to a rubbery behaviour for M3-M5. The
latter materials exhibited a frequency-dependent viscosity |.eta.*|
with a slope of -1 and both their storage (G') and loss moduli
(G'') at 1 rad/s (as a reference) were increased by 7 and 3 orders
of magnitude, respectively, as compared to the constituent low
molecular weight PIB-NH.sub.2 1. Difunctional derivatives D0-D2 can
give rise to supramolecular networks and, hence, showed an even
more drastic transition of mechanical properties. Thus, D2 was
already brittle hard solid and exhibited a storage modulus of G'=2
MPa, 10 times higher than high molecular weight poly(isobutylene)
(MW.gtoreq.200'000). The pure higher difunctional homologues D3-D5
were hard and brittle powders that could not be processed into
solid discs and, hence, were not investigated by means of shear
rheology.
[0096] Shear rheology on the monofunctional derivatives M0-M5 in
comparison to unmodified polyisobutylenes revealed a transition of
mechanical properties from moderately viscous liquid (M0-M1) to a
rubbery behaviour (M4-M5) as a function of the number of n alanine
residues. Starting with the .beta.-sheet tape-forming derivatives
M2-M3, the materials exhibited shear moduli exceeding those of high
molecular weight polyisobutylenes, although the molecular weight of
the attached polymer (MW 1,200) was far below the entanglement
length of PIB (.apprxeq.15,000) and the monofunctional derivatives
cannot form hydrogen-bonded networks. The storage and loss moduli
at 1 rad/s within the series leveled off towards G'.apprxeq.0.6 MPa
and G''.apprxeq.0.06 MPa, indicating that a further increase of the
oligopeptide length would not substantially affect the materials'
mechanical properties anymore. The network-forming difunctional
derivatives D0-D2 showed even more drastic changes in mechanical
properties as a function of oligopeptide length, and D2 was already
a brittle hard solid (G'=2 MPa). While the notion of a mechanical
reinforcement is certainly well in line with previous examples of
supramolecular networks, the observed drastic dependence of
aggregation behaviour and mechanical properties on the number of
hydrogen-bonding sites allows for further tailoring of the
materials' thermomechanical properties in blends of the
investigated compounds. See in particular the disclosed
examples.
[0097] Referring in particular to FIGS. 10A-14E (Examples 1-6):
[0098] FIGS. 10A, 10B and 10C provide Rheological dynamic frequency
sweep experiments at 25.degree. C. of different binary blends
M2/D2, showing a) storage moduli G' (FIG. 10A), b) loss moduli G''
(FIG. 10B), and c) and viscosity |.eta.*| (FIG. 10C); the shear
viscosity formally defined as
|.eta.*|=p.sub.21,0/(.gamma..sub.0.omega.) was used as a calculated
entity to compare both liquid and rubbery materials. The .gamma. (%
strain) values used for these experiments ranged from 0.05-0.5% and
were selected to be in the linear regime of the investigated
materials. Large changes in the moduli and viscosities were
observed between pure M2 and blends of up to 10 wt % of D2 in M2,
which then level off for compounds with higher content of D2.
[0099] FIGS. 11A, 11B, 11C and 11D provide a Comparison of storage
(G') and loss moduli (G'') at 1 rad/s of unmodified PIB as a
function of molecular weight (FIG. 11A), M0-M5 and D0-D2 as a
function of hydrogen-bonding sites per end group (FIGS. 11C and
11D), and of different blends of M2/D2 as a function of composition
(FIG. 11B). Storage moduli of M3-M5 and D2 as well as blends of at
least 5% D2 in M2 (G'=0.6 MPa) exceed those of high molecular
weight PIB. Blends M2/D2 with more than 10 wt % of D2 show storage
and loss moduli levelling off toward G'.apprxeq.2 MPa and
G''.apprxeq.0.1 MPa as observed for pure D2. In the series of
M0-M5, the storage and loss moduli appear to converge toward
G'.apprxeq.0.6 MPa and G''.apprxeq.0.06 MPa.
[0100] FIGS. 12A and 12B provide Amide A, amide I and amide II
regions of the solid-state infrared (IR) spectra of M2, D2, as well
as the complete series of M2/D2 blends (Examples 1-6). They all
exhibited a single amide A absorption at 3275 cm.sup.-1, a strong
and sharp amide I absorption at 1627 cm.sup.-1, a smaller
absorption at 1687 cm.sup.-1 and an amide II absorption at 1543
cm.sup.-1, consistent with the presence of highly ordered
antiparallel .beta.-sheet structures. Moreover, independent of
their composition, the blends exhibited and amide I regions
indistinguishable from the pure compounds, providing evidence for
the presence of antiparallel .beta.-sheet structures and the
miscibility of D2 in M2.
[0101] FIGS. 13A and 13B providing Amide I regions of the
temperature-dependent solid-state IR spectra of the blend M2/D2 9:1
(Example 3) revealed that the position and intensity of the
absorption bands associated with .beta.-sheet secondary structures
remained virtually unchanged until at least 150.degree. C. but then
rapidly decreased at temperatures above 170.degree. C., exactly
like the individual components M2 and D2. This proves that the
thermal transitions were associated to `melting` (deaggregation)
and `crystallization` (aggregation) of the .beta.-sheet aggregates
and that the two compounds formed a common hydrogen-bonded network
together. FIG. 13C shows that whereas the parent PIB amine was a
viscous liquid, M2 was a sticky solid, D2 a brittle powder, the
blend M2/D2 9:1 (Example 3) was an `inherently reinforced`
thermoplastic elastomer.
[0102] FIGS. 14A, 14B, 14C, 14D, 14E and 14FE show
Temperature-dependent shear rheology of the blend M2/D2 9:1
(Example 3) (FIGS. 14A, 14B and 14C) and PIB (MW 200'000) as a
reference material (FIGS. 14D, 14E and 14F). Plots of storage
modulus G', loss modulus G'', and viscosity |.eta.*| (formally
defined as |.eta.*|=p.sub.21,0/(.gamma..sub.0.omega.) and used as a
calculated entity to compare both liquid and rubbery materials) at
1 rad/s as a function of temperature and examples of rheological
dynamic frequency sweep experiments at 25.degree. C. and
180.degree. C. showed that M2/D2 9:1 (Example 3) experienced a
sharp and single-step decrease of its moduli and viscosity at a
temperature of about 160.degree. C., yielding low viscosity
liquids, that had been found to coincide with .beta.-sheet
deaggregation according to temperature-dependent IR spectroscopy
(see FIG. 13a). By contrast, PIB (MW 200'000) remained in the
rubbery state up to temperatures of at least 250.degree. C. The
.gamma. (% strain) values used for these experiments ranged from
0.1-50% and were selected to be in the linear regime of the
investigated materials.
[0103] Binary blends of monofunctional and difunctional derivatives
with "matching" oligopeptide segments were found to give rise to
supramolecular networks that were "inherently reinforced" by the
incorporated .beta.-sheet aggregates. Specifically, binary blends
of the thermally processable compounds M2 and D2 with the
compositions (by weight) M2/D2 99:1 (Example 1), 95:5 (Example 2),
9:1 (Example 3), 7:3 (Example 4), 5:5 (Example 5), and 1:9 (Example
6) were obtained by dissolving mixtures of the compounds in
tetrachlorethane (TCE), stirring the solutions at room temperature
for 16 h, removing the solvent in vacuo, and drying the resulting
materials in high vacuum at 120.degree. C. for 3 days. Independent
of their composition, the blends exhibited solid state IR spectra
with amide I regions indistinguishable from the pure compounds and
underwent a single-step "melting" transition at 160-170.degree. C.
associated with the deaggregation of all .beta.-sheet structures.
The blends yielded rubbery materials with shear moduli that
exceeded those of even high molecular weight PIB (0.2 MPa) by an
order of magnitude even for low fractions of D2. Thus, the storage
moduli already reached G'=0.6 MPa upon the addition of .gtoreq.5 wt
% D2 (Examples 2-6) and leveled off toward G' 2 MPa for 10 wt % D2
(Examples 3-6). At the same time, the latter materials (Examples
3-6) experienced a sharp decrease of their moduli and viscosities
at their melting temperatures. Hence, we obtained "inherently
reinforced" polyisobutylene-based thermoplastic elastomers that
exhibited superior shear properties and showed lower creep
behaviour at room temperature, but yielded well-processable melts
at elevated temperatures, well below their decomposition
temperature.
[0104] Table 1 shows representative values of storage moduli G',
loss moduli G'', loss factors tan .delta., and viscosities |.eta.*|
for different grades of polyisobutylenes, M0-M5, D0-D2, as well as
Examples 1-6.
[0105] Referring in particular to FIGS. 15A-22H (Examples
7-10):
[0106] FIGS. 15A, 15B and 15C show Temperature-dependent shear
rheology of the blend M2/D1 1:4 (Example 7), compared to pure M2,
pure D1, as well as the blend M2/D2 9:1 (Example 3). Plots of a)
storage modulus G' (FIG. 15A), b) loss modulus G'' (FIG. 15B), and
c) viscosity |.eta.*| (formally defined as
|.eta.*|=p.sub.21,0/(.gamma..sub.0.omega.) was used as a calculated
entity to compare both liquid and rubbery materials) (FIG. 15C) at
1 rad/s as a function of temperature revealed that the material
underwent a two-stage thermomechanical transition, first following
the behaviour of D1 in the temperature range of -45.degree. C. to
above room temperature, and then M2 between 65.degree. C. and the
melting transition at 139.degree. C. FIG. 15D providing
Differential scanning calorimetry showed that the blend M2/D1 9:1
(Example 7) exhibited a transition at about 25.degree. C. (assigned
to the melting of the D1 network) as well as a reversible
transitions at 139.degree. C. (onset at 128.degree. C.; 0.5 J/g)
that we assigned to the reversible deaggregation of the
.beta.-sheet aggregates of M2. While the apparent `melting point
depression` as compared to pure M2 (170.degree. C.) suggests a
certain interaction between D1 and M2 in that temperature range,
the pronounced effect of the minority component M2, the two-stage
temperature transition, and the superimposed rheological properties
of the pure components in the blend provide sufficient evidence for
the presence of two independent, `interpenetrating supramolecular
networks` that do not undergo macrophase segregation.
[0107] FIGS. 16A, 16B and 16C provide Rheological time-temperature
superposition (TTS) master curves of D1, M2, and M2/D1 1:4 (Example
7) at T.sub.ref=25.degree. C. D1 showed an entanglement point at
25.degree. C. M2 exhibited a large tan .delta. peak at an unusual
temperature as compared to high molecular weight PIB. M2/D1 1:4
(Example 7) possessed a broad region with pronounced `liquid-like`
behaviour and a large loss factor of up to tan .delta.=2.0.
[0108] FIGS. 17A, 17B, 17C and 17D provide Amide A, amide I and
amide II regions of the solid-state infrared (IR) spectra of
solution phase IR spectra of samples in tetrachlorethane (FIGS. 17A
and 17B) and sold state IR spectra of bulk samples of PIB (MW
75'000), PIB (MW 35'000), M2, D2, as well as different binary and
ternary blends of M2 and D2 in PIB (MW 75'000) or PIB (MW 35'000)
(Examples 8-10) (FIGS. 17C and 17D). All mixtures were deaggregated
in solution, but bulk materials exhibited a single amide A
absorption at 3276 cm.sup.-1, a strong and sharp amide I absorption
at 1627 cm.sup.-1, a smaller absorption at 1687 cm.sup.-1 and an
amide II absorption at 1543 cm.sup.-1, consistent with the presence
of highly ordered antiparallel .beta.-sheet structures. Moreover,
independent of their composition, the blends exhibited and amide I
regions indistinguishable from the pure compounds, providing
evidence for the presence of antiparallel .beta.-sheet structures
dispersed in a PIB matrix.
[0109] FIG. 18A providing Differential scanning calorimetry
revealed that M2, D2 and their binary and ternary blends M2/PIB (MW
35'000) 5:5 (Example 9) and M2/D2/PIB (Mw 35'000) 4:1:5 (Example
10) exhibited reversible thermal transitions at the onset
temperatures of 170.degree. C. (16 J/g), 178.degree. C. (17 J/g),
172.degree. C. (6.2 J/g) and 169.degree. C. (6.1 J/g) respectively.
These could be assigned to the `melting` (deaggregation) and
`crystallization` (aggregation) of the .beta.-sheet aggregates.
FIG. 18B providing Thermogravimetric analysis revealed that all
materials except PIB-NH2 and Smactane.TM. stable up to temperatures
of at least 250.degree. C. Complete PIB depolymerization occurred
at temperatures above 340.degree. C. for M2, D2 and their
blends.
[0110] FIG. 19 Atomic force microscopy (AFM) height (left) and
phase (right) images of M2 and M2/PIB (MW 35'000) 5:5 (Example 9)
drop-cast from tetrachlorethane solution onto SiO.sub.2 substrates
revealed the formation of .beta.-sheet fibrils, while AFM images of
pure PIB (MW 35'000) did not show such features.
[0111] FIGS. 20A, 20B, 20C, 20D, 20E and 20F provide Rheological
time-temperature superposition (TTS) master curves of M2/PIB (MW
75'000) 5:5 (Example 8), M2/PIB (MW 35'000) 5:5 (Example 9),
M2/D2/PIB (MW 35'000) 4:1:5 (Example 10), in comparison to
unmodified higher molecular weight PIB (MW 35'000, MW 75'000, MW
200'000) as well as Smactane.TM. (hollow symbols in other graphs)
as reference materials at T.sub.ref=25.degree. C. The blends gave
rise to soft materials with a loss factor of tan .delta.>0.6
over almost the whole frequency range investigated.
[0112] FIG. 21A provides an Experimental setup for a random
vibration modal analysis test on a sandwich structure comprising a
damping layer. The first resonance frequency of the steel structure
at 32.9 Hz and its intensity decrease in the sandwich structure for
PIB (MW 200'00), M2, for M2/D2/PIB (75 k) 5:5 (Example 8),
M2/D2/PIB (35 k) 5:5 (Example 9), M2/D2/PIB (35 k) 4:1:5 (Example
10) as well as Smactane.TM.. FIG. 21B provides Finite element
simulations of the same sandwich configurations for damping layers
based on the same materials as well as additional commercial
damping materials. FIG. 21C provides experimental (circles) and
calculated (squares) loss moduli G'' considered as the "figure of
merit" for the vibration damping ability in constrained layers
plotted relative to the damping ratio with an exponential fit
presented here only as a guide. These experimental G'' values were
taken at 200 rad/s from a classical rheological frequency sweep
test at 25.degree. C.
[0113] FIGS. 22A, 22B, 22C, 22D, 22E, 22F, 22G and 22H provide
Lissajous curves obtained from oscillatory shear stress-strain
test, performed on a rheometer at -45.degree. C. and 25 rad/s with
a theoretical .gamma. (% strain) value imposed of 0.05%. FIG. 22A
is for Smactane.TM. with .gamma.=0.05%, FIG. 22B is for PIB (Mw
200'000) with .gamma.=0.045%, FIG. 22C is for M2 with
.gamma.=0.05%, FIG. 22D is for M2/PIB (MW 35'000) 5:5 (Example 9)
with .gamma.=0.068%, and FIG. 22E is for M2/D2/PIB (MW 35'000)
4:1:5 (Example 10) with .gamma.=0.043%. The areas of the latter
represent dissipated energies during one cycle of these tests.
FIGS. 22F, 22G and 22H provide Plots of the logarithm of the
dissipated energies relative to the logarithm of the strain applied
at -45.degree. C. during an oscillatory shear stress-strain test,
performed on a rheometer. Based on the mathematical equation:
W.sub.d=nG'' .epsilon..sub.0.sup.2, fitting equations of PIB (Mw
200'000) (see FIG. 22F), M2/PIB (MW 35'000) 5:5 (Example 9) (see
FIG. 22G), and M2/D2/PIB (MW 35'000) 4:1:5 (Example 10) (see FIG.
22H) were used in order to correct the values of energies
dissipated for a .gamma. (% strain) value of 0.05% for all the five
materials.
[0114] Binary and ternary blends of compounds with "non-matching"
oligopeptides (different oligopeptide length or amino acid
sequence, including non-functionalised polymers) were found to give
rise to novel "interpenetrating supramolecular networks".
Specifically, binary blends (compositions by weight) of M2/D1 1:4
(Example 7), M2/polyisobutylene MW 75'000 (Example 8),
M2/polyisobutylene MW 35'000 (Example 9), as well as the ternary
blend M2/D2/polyisobutylene (MW 35'000) 4:1:5 (Example 10) were
obtained by dissolving mixtures of the compounds in
tetrachlorethane (TCE), stirring the solutions at room temperature
for 16 h, removing the solvent in vacuo, and drying the resulting
materials in high vacuum at 120.degree. C. for 3 days. The
annealing temperature was chosen such that it was below the melting
transition of the tape-forming component M2 but above the softening
temperature of the second network-forming component (D1 or
polyisobutylene). In the case of Example 7, the gelation point of
M2 and the entanglement point of D1 were superimposed in their
blends, resulting in materials with both a gelation and an
entanglement point within a similar frequency range (in a classical
rheological frequency sweep at 25.degree. C.). Thus the material
exhibited a large frequency region with pronounced liquid-like
behaviour (that is, G''>G') at room temperature, confined by two
regions of solid elastomer-like (G'>G'') behaviour at higher and
lower shear frequencies, as seen from a rheological
time-temperature superposition (TTS) master curve at 25.degree. C.
(all shift factors log a.sub.T and log b.sub.T. for the TTS master
curves listed in Table 2). Whilst high molecular weight
polyisobutylene materials may exhibit such regions of "liquid-like"
behaviour at temperatures just above their glass transition
temperature of T.sub.g.apprxeq.-65.degree. C., Example 7 showed
such behaviour over a large temperature range and a broad frequency
range at room temperature, exhibiting a large loss factor of up to
tan .delta.=2.0 in this region, which is unprecedented in related
materials. Upon heating Example 7 above the materials'
glass-transition temperature of about T.sub.g=-55.degree. C., both
its storage and loss moduli as well as viscosity first closely
followed those of D1. At about 65.degree. C., where pure D1 is
already in its liquid regime, the storage and loss moduli became
similar to those of M2 and remained constant up to the melting
transition at above 139.degree. C. The pronounced effect of the
minority component M2, the two-stage temperature transition, and
the superimposed rheological properties of the pure components in
the blend provide sufficient evidence for the presence of two
independent hydrogen-bonded superstructures, resulting in an
"interpenetrating supramolecular network". Whereas the high
frequency boundary of the "liquid-like" region can be assigned to
the hydrogen-bonded network formed by D1, the low frequency
boundary is associated to a percolation network of the M2
.beta.-sheet tapes.
[0115] The observed pronounced "liquid-like" behaviour (in a
certain temperature and mechanical frequency range), that is, the
apparent molecular level properties approaching liquid-like
properties (such as flow) in a solid and macroscopically
shape-persistent polymer materials is the prerequisite for the
self-healing properties of the materials disclosed here.
[0116] For the various blends of M2, D2, and polyisobutylenes
(Examples 8-10), Solution-phase IR spectra in chlorinated solvents
showed that all mixtures remained non-aggregated in solution. This
enabled us to obtain homogenous blends from solution, so as to
obtain hydrogen-bonded aggregates from M2 dispersed in PIB as a
matrix material (Examples 8-9) that can be cross-linked using the
difunctional network-forming difunctional D2 (Example 10).
Solid-state infrared (IR) spectroscopy of the bulk materials
revealed strongly aggregated and highly ordered antiparallel
.beta.-sheet structures. Thermogravimetric analysis (TGA) and
differential scanning calorimetry (DSC) proved that the materials
were thermally stable (against degradation) up to at least
250.degree. C. All materials exhibited sharp and reversible thermal
transitions at temperatures of 169-178.degree. C. Comparing the
enthalpies of fusion of pure M2 and D2 (16-17 J/g) to those of the
blends (6-6.5 J/g), we concluded that 75-80% of the
oligopeptide-modified components were aggregated into .beta.-sheet
tapes or fibrils in the blends. Moreover, temperature-dependent
solid-state IR spectroscopy on the materials proved that the
observed transition was associated to .beta.-sheet deaggregation in
all cases. Visualization of the nanoscopic morphologies of the
obtained aggregates by means of atomic force microscopy (AFM)
imaging of continuous 1 .mu.m thick films drop-cast from TCE
solution onto SiO.sub.2 substrates proved that .beta.-sheet tape or
fibril structures were present in those bulk materials.
[0117] In order to evaluate the mechanical properties of Examples
8-10, we tested their rheological properties in comparison to
unmodified higher molecular weight polyisobutylenes (MW 35'000,
75'000, and 200,000). Compared to Smactane.TM., PIB (MW 200'000)
showed lower storage (G') and loss moduli (G'') over the whole
range of frequencies (10.sup.-4-10.sup.6 rad/s) or temperatures
(-45.degree. C.-105.degree. C.) investigated, but a slightly higher
and broader peak of the loss factor tan .delta. (as a function of
frequency), corresponding to its glass transition. Likewise, pure
M2 and the blends M2/PIB (Examples 8-9) showed significantly lower
storage and loss moduli over a frequency range of
10.sup.-4-10.sup.3 rad/s, but significantly higher loss factors of
tan .delta..apprxeq.1 for a large frequency range. Examples 8 and
9, for instance, gave rise to a soft rubber-like material with
storage and loss moduli G' and G'' that were very similar to one
another over almost the complete range of investigated frequencies,
as determined from a TTS master curve at room temperature (all
shift factors log a.sub.T and log b.sub.T. for the TTS master
curves listed in Table 2). As a result, the loss factor of Examples
8 and 9 peaked at tan .delta.=1.1 at a reduced frequency of about
a.sub.T.omega.=5.times.10.sup.4 rad/s and never fell below tan
.delta.=0.6 in the reduced frequency range of
a.sub.T.omega.=10.sup.-3-10.sup.6 rad/s. The ternary blend
M2/D2/PIB (MW 35'000) 4:1:5 (Example 10) possesses higher storage
and loss moduli as well as similar loss factors compared to
Smactane.TM., but with even improved moduli and loss factor at low
frequencies, due to the addition of D2 which acts as a network
forming crosslinker, resulting in an extra reinforcement of the
materials. Moreover, the obtained master curves had a substantially
different shape as compared to either pure M2 or unmodified
polyisobutylenes. Specifically, the rubbery plateau in the
low-frequency regime was absent, indicating that neither does M2
just serve as a filler, nor does the polyisobutylene matrix just
act as a diluting "solvent". One can therefore attribute the large
temperature and frequency range of high loss factor tan .delta.
values to an interpenetration of the PIB entanglement network and a
percolation network formed by the M2 .beta.-sheet tapes. The
resulting supramolecular networks exhibit an improved vibration
damping performance was attributed to improved energy dissipation
by the high fraction of pendant polymer chains incorporated into
the network.
[0118] Such "interpenetrating supramolecular networks" as described
here provide an alternative to traditional IPNs for the preparation
of high-performance vibration damping materials. In order to
evaluate the performance of Examples 8-10 with other damping
materials, we tested their shear vibration damping characteristics
of the in comparison to unmodified higher molecular weight
polyisobutylenes (MW 35'000, 75'000, and 200,000), as well as
Smactane.TM., a commercially available high performance damping
material with excellent damping properties specifically at low
temperatures. To this end, we employed a random vibration modal
analysis test on a sandwich structure representing a typical
constrained damping layer application. The test structure was
designed specifically to investigate the structural damping
performance of the material in low frequency vibration (30-40 Hz),
which is typical for the first vibration modes of many steel or
aluminium panels used in automotive or aerospace applications. The
specimen with a free length of 54 mm was fixed at one end while a
weight of 4.6 g was clipped to the other one. The beam was excited
using a pseudo random signal using a vibration shaker driven
through an open loop random vibration controller. Accelerometers
were used to monitor the base and tip accelerations and reconstruct
the frequency response function of the system around its first
resonance peak. The modal damping ratio was obtained by
single-degree-of-freedom modal curve fitting of the resonance peak
in the complex domain.
[0119] All results of the damping tests and derived damping ratios
are listed in Table 3.
[0120] The first resonance frequency for steel, in this particular
set up, occurred at 32 Hz and was slightly damped by steel itself
and its clamping on the base (0.4%). However, damping ratio
significantly increased to 2.9% once Smactane.TM. was used as the
damping layer in the sandwich structure. By comparison, while
unmodified high molecular weight PIB (MW 200'000) exhibited a low
damping ratio of 1.4%, the damping ratios were 3.2% for pure M2,
2.6% for the binary blend M2/PIB (MW 75'000) (Example 8), 2.5% for
the binary blend M2/PIB (MW 35'000) (Example 9), and 3.4% for the
ternary blend M2/D2/PIB (MW 35'000) (Example 10). Examples 9 and 10
thus showed excellent damping ratios, even exceeding those of the
commercially available high-performance damping material
Smactane.TM. and by far surpassing those of unmodified PIB that is
considered to possess good damping properties and is already used
in damping applications on a technological scale.
[0121] Moreover, we complemented our results with detailed finite
element (FE) simulations of the sandwich beam vibration tests. We
performed the simulations also on other high-performance damping
materials as a reference, including Smactane.TM., Soundcoat.TM.
Dyad 601 and 3M ISD.TM. 130.54 using their rheology data as the
input to the FE simulations. In qualitative agreement with the
experimental results, the finite element simulations resulted in
damping ratios of Examples 8-10 above those of the reference
materials.
[0122] The damping properties of the investigated materials at low
temperatures were obtained using the loss moduli G'' obtained from
rheological frequency sweep experiments at those temperatures and
calculating the dissipated energies during one cycle of oscillatory
stress-strain test from the area of the corresponding Lissajous
curves (corrected for the imposed strain).
[0123] The calculated damping properties at low temperatures are
listed in Table 4.
[0124] All of the investigated materials and, in particular, the
ternary blend M2/D2/PIB (MW 35'0000) (Example 10) exhibit excellent
damping properties at temperatures of -45.degree. C. and below,
down to their glass transition temperatures at about
T.sub.g.apprxeq.-65.degree. C., even exceeding those of
Smactane.TM. and in marked contrast to PIB (MW 200,000) that
possesses lower loss moduli over the whole range of frequencies
(0.1-100 rad/s) tested at -45.degree. C.
[0125] All commercial reference materials (except PIB) are
composites with formulations highly optimized for damping
performance. It is worth noting that, as a consequence, our
materials were light (with a density of 0.92 g/cm.sup.3 compared to
1.18 g/cm.sup.3 for Smactane.TM.) and did not require any
additional fillers or low molecular weight plasticizers. Moreover,
as the "ideal" damping characteristics depend on the application,
the versatility offered by the use of oligopeptide-modified
polymers as additives to commercial elastomers appears to provide
an excellent pathway towards lightweight, low-creep, and
high-performance constrained layers for vibro-acoustic damping.
[0126] Referring in particular to FIGS. 23A-25 (Examples
11-13):
[0127] FIGS. 23A, 23B, 23C and 23D show AFM phase images and
differential scanning calorimetry measurements that prove that
oligopeptide aggregation competes with polymer phase segregation in
blends of M2/S2 and M3/S3 and leads to nanoscale phase segregation.
As shown in FIG. 23A, blends of M2 with unmodified polystyrene
(9:1) as a reference material domains (white) with average sizes of
hundreds of nanometers in diameter, only very few such polystyrene
domains could be observed in blends M2/S2 (9:1) (Example 11) that
were otherwise mostly homogeneous. As shown in FIG. 23B, in
differential scanning calorimetry measurements, the glass
transition of PS could be identified in both cases, demonstrating
that even in the blend M2/S2 (9:1) (Example 11) nanoscale
polystyrene domains are present. As shown in FIG. 23C, blends of M3
with unmodified polystyrene (9:1) as a reference material forms
polystyrene domains (white) with average sizes of dozens of
nanometers in diameter, and no such polystyrene domains were
observed in blend M3/S3 (9:1) (Example 12). As shown in FIG. 23D,
in this case, the polystyrene glass transition was only observed in
the differential scanning calorimetry measurements of the reference
materials and not in the blend M3/S3 (9:1) (Example 12), proving
that the polystyrene domain size was now so small that no
cooperative properties such as a glass transition were
observed.
[0128] FIGS. 24A and 24B show a) storage moduli G' (FIG. 24A) and
b) loss moduli G'' (FIG. 24B) determined by rheological dynamic
frequency sweep experiments at 25.degree. C. of M2, S2, the blend
of M2 with unmodified polystyrene (9:1) and the blend M2/S2 9:1
(Example 11); FIGS. 24C and 24D show storage moduli G' and loss
moduli G'' determined by rheological dynamic frequency sweep
experiments at 25.degree. C. of M3, S3, the blend of M3 with
unmodified polystyrene (9:1) and the blend M3/S3 9:1 (Example 12).
Notably, whereas the storage and loss moduli of the blends of
either M2 or M3 with unmodified polystyrene (9:1) increased
slightly compared to the pure M2 and M3, respectively, a strong
increase in the storage and loss moduli was observed for the two
blends M2/S2 9:1 (Example 11) and M3/S3 9:1 (Example 12). In
particular, in the low frequency range, the addition of 10 wt % of
S3 in the blend M3/S3 9:1 (Example 12) is sufficient to achieve a
ten-fold increase in storage modulus G'. This proves that the
polystyrene domains in these blends served to "glue together" the
oligopeptide aggregates (tapes and nanofibrils), giving rise to a
new percolation network of hard domains within the material.
[0129] FIG. 25 shows a rheological time-temperature superposition
master curve at a reference temperature of 25.degree. C. of the
ternary blend M3/S3/PIB (MW 35'000) 9:3:12 (Example 13). The
storage moduli G', loss moduli G'' and the loss factor tan .delta.
are reported for a large range of reduced frequency
(10.sup.6-10.sup.-4 rad s.sup.-1), as well as temperature range
(-45.degree. C. to 65.degree. C.). The .gamma. (% strain) values
used for these experiments was 0.1% and were selected to be in the
linear regime of the investigated materials. The TTS master curve
revealed that the tan .delta. was higher than 0.35 over a large
frequency range of 10.sup.-3-10.sup.-5 rad s.sup.-1, and the G''
was higher than 1 Mpa for frequencies higher than 10 rad s.sup.-1.
This combination of a large temperature range of a large loss
factor and a high storage modulus is ideal for good,
temperature-invariant damping properties. Accordingly, the material
showed a very high damping ratio of =3.3% in constrained layer
damping tests on a sandwich structure at a resonance frequency of
32 Hz, as described in previous examples.
2. PREPARATIVE EXAMPLES
2.1 Instrumentation and Methods
[0130] NMR Spectroscopy was carried out on a Bruker Avance 300
spectrometer operating at a frequency of 300.23 MHz for .sup.1H and
75.49 MHz for .sup.13C nuclei, or on a Bruker Avance 400
spectrometer operating at a frequency of 400.23 MHz for .sup.1H and
100.63 MHz for .sup.13C nuclei. Deuterated solvents were purchased
from Cambridge Isotope Laboratories. The spectra were calibrated to
the respective residual proton peaks of the deuterated solvents
(.sup.1H NMR: 7.26 ppm CDCl.sub.3, 6.0 ppm TCE-d.sub.2, 5.32 ppm,
DMSO-D.sub.6, 3.31 ppm CD.sub.3OD; .sup.13C NMR: 77.16 ppm
CDCl.sub.3, 49.00 ppm CD.sub.3OD, 39.52 ppm DMSO-D.sub.6).
[0131] Solution--Phase FTIR Spectra were recorded on a "Spectrum
One" IR spectrometer from Perkin Elmer using a solution-phase
cuvette with KBr windows and a light path of 0.5 mm, or on a Jasco
FT/IR-6300 Fourier, using a KBr window 32.times.3 mm (by Pike
Technology). The materials were dissolved in either TCE or
CHCl.sub.3 (5 mg/mL). The TCE solutions were stirred for 1 h at
100.degree. C. and left to slowly cool to room temperature. The
CHCl.sub.3 solutions were stirred at RT for 1-16 h.
[0132] Solid-State FTIR Spectra were recorded on a Bruker ALPHA
FTIR spectrometer or on a JASCO FT/IR 6300 spectrometer using the
Miracle ATR accessory from PIKE, as well as on a Varian Fourier
Transform spectrometer equipped with a Golden Gate diamond ATR with
temperature control up to 200.degree. C.
[0133] High Resolution Mass Spectra were recorded at the Mass
Spectrometry Service of EPFL on either an AXIMA Performance device
from Shimadzu Biotech for MALDI-TOF and a Q-TOF Ultima from Waters
for ESI-TOF, or at the Mass Spectrometry Service of ETH Zurich on a
Bruker Daltonics maXis for HiRes-ESI-MS.
[0134] Thermogravimetric Analyses were performed on a TGA Q500
device from TA Instruments, loaded with samples of more than 2 mg.
The measurement range was 50-710.degree. C. in a nitrogen
atmosphere or 50-914.degree. C. in air. A heating rate of
10.degree. C./min was applied in all cases.
[0135] Differential Scanning calorimetry was performed on a DSC
Q1000 from TA Instruments in a nitrogen atmosphere, loaded with
samples of more than 2 mg. For the measurement of glass transition
or melting temperatures, three measurements were performed in the
range of -80 to 400.degree. C. Both the heating and cooling rates
were 10.degree. C./min. All data were collected from the second
heating cycle.
[0136] Combustion Elemental Analyses were carried out as service
measurements at EPFL using EA 1100 CHN Instrument or at the
Institute of Organic Chemistry at ETH Zurich using a LECO CHN/900
instrument.
[0137] AFM Imaging was performed on a Nanoscope IIIa instrument.
Samples were prepared from stock solutions of the compounds in
tetrachlorethane (TCE) at an initial concentration c=10.sup.-3
mol/L. The solutions were placed into sealed tubes, vigorously
stirred (400 rpm) and heated in an oil bath to 180.degree. C. for 2
h, followed by stepwise cooling 160.degree. C. (1 h), 140.degree.
C. (1 h), 120.degree. C. (1 h), and 100.degree. C. (1 h) under
continued stirring (100 rpm). Afterwards, the heating was switched
off and the solution was allowed to cool to room temperature at the
same stirring rate. The solutions were then diluted to a
concentration of c=1.times.10.sup.4 mol/L or c=5.times.10.sup.-5
mol/L and spin-coated onto SiO.sub.2 substrates treated with
ethanol and ultrapure water (3000 rpm) or onto freshly cleaved HOPG
(1800 rpm). The obtained samples were analyzed in tapping mode at
room temperature in air, using cantilevers with an average
resonance frequency of 75 kHz and scan rates of 0.5-1.5 Hz. The
image resolution was 512.times.512 pixels.
[0138] Dynamic Shear Rheology Measurements were carried out on
parallel plate rheometers AR 2000, ARES LR2 or ARES from TA
Instruments. Disc shaped sample specimen from all materials that
were shape-persistent were prepared on a Rittal table press. A
force of 2 kN was applied for 30 min at 100.degree. C., after which
the specimen were cooled to 20.degree. C. at a force of 1.1 kN for
30 min. Depending on the samples quantity, aluminium plates of 15
mm or 25 mm diameter, as well as stainless steel plates of 25 mm
diameter were used. Discs of 12 mm diameter were prepared as well
and placed in the centre of the stainless-steel plates (25 mm
diameter) with a centring tool. The gap between the plates was in
the range of 0.4-2 mm. In the case of non-adhesive samples, the
plates were covered with emery paper to avoid wall slipping.
Measurements were carried out at temperatures of -45.degree. C. to
250.degree. C. Once the desired temperature was reached, the system
was equilibrated for 2 min. Frequency sweeps ranging from 100 rad
s.sup.-1 to 0.01 rad s.sup.-1 were carried out under controlled
strain. Depending on the sample composition and temperature, the
strain amplitude ranged from 0.03% to 50%. The applied strain was
defined such that the sample stayed in its linear viscoelastic
domain during the complete frequency sweep.
[0139] Modal Damping Tests were carried out on a RMS 3000 vibration
shaker using an HP 35670A vibration controller and signal analyzer
in open loop pseudo random vibration analysis. The acceleration of
the base and of the tip of the specimen were monitored using two
Bruel&Kaer 4517 accelerometer through a B&K Nexus 2692
amplifier. The frequency transfer functions of the different
specimens were measured in a frequency range of 10 to 110 Hz with a
resolution of 0.125 Hz. The test specimen consisted of a sandwich
structure representing a constrained damping layer application. The
base substrate was a steel plate with dimensions of
60.times.6.times.0.5 mm onto which a damping layer with dimensions
of 40.times.0.6.times.2.2 mm was superimposed. A thin steel plate
with dimensions of 40.times.0.6.times.0.2 mm was used to constrain
the top of the damping layer. The sandwich test structure was then
clamped with one end to the vibration table over a length of 6 mm,
and a mass of 4.6 g was added to the free end of the beam over a
length of 5 mm. The added mass has been calculated such that the
first bending mode of the sandwich beam is in the range of 30-40
Hz. In order to determine the modal damping ratio the first peak of
the frequency response function was first fitted using a complex
polynomial fraction least square method integrated in the HP 35670A
signal analyzer. The modal damping ratio was then calculated from
the real part X and the imaginary part .omega. of the first complex
pole of the polynomial fraction using the definition
.xi.=.lamda./.omega..
[0140] Finite Element Simulations of the sandwich beam vibration
tests have been carried out to compare the damping performance of
M2/PIB with other high-performance damping materials in this
particular application. The chosen specimen geometry was the same
as the sandwich beam specimen used in the experimental modal
damping tests (60.times.6.times.0.5 mm steel base plate with a
40.times.0.6.times.2.2 mm damping layer constrained on top by a
40.times.0.6.times.0.2 mm steel sheet). The base plate was modelled
as being clamped to the shaker on one side (imposed displacement,
no rotation, over 6 mm) and attached to two steel blocks of
5.times.5.times.15 mm. The steel plates were modelled as linear
elastic with a Young modulus of 210 GPa, Poisson ratio of 0.3 and
mass density of 7,800 kg/m.sup.3. The damping layer materials were
all considered incompressible and modelled using an Arruda-Boyce
hyperelastic potential (power exponent of 7) with complex shear
moduli taken from rheology measurements or literature data (Table
3). The whole specimen was modelled using 2640 3D hybrid quadratic
hexahedric elements (14345 nodes) in Simulia Abaqus.COPYRGT. 6.10
(FIG. 4e in the main paper) and subjected to a steady state
harmonic simulation in a frequency range of 20-70 Hz with a
resolution of 0.333 Hz to capture in detail the first resonance
peak at around 34 Hz. The simulated frequency response functions
(tip vs base acceleration in the direction of the Y axis) were then
processed to extract the modal damping ratio .xi. (Table 3) using a
single degree of freedom complex curve fitting method.
2.2 Materials and General Synthesis Procedures
Materials.
[0141] Reagents were purchased as reagent grade from commercial
sources and used without further purification. Poly(isobutylene) 1
(Kerocom.TM. PIBA) was obtained from BASF and purified from
non-functionalized poly(isobutylene) by column chromatography prior
to use. Poly(isobutylene) diamine 4 containing about 10% of
monofunctional poly(isobutylene) amine was obtained from BASF SE,
Germany, and used without further purification. THF, acetonitrile,
toluene, dichloromethane and triethylamine were purchased as HPLC
grade and dried using a solvent purification system from Innovative
technologies. Other solvents were purchased as reagent grade and
distilled once prior to use. Thin Layer Chromatography (TLC)
Analyses were performed on TLC plates from Merck; UV-light (254 nm)
or standard colouring reagents were used for detection. Column
Chromatography was conducted on Geduran.RTM. Silica gel Si 60 from
Merck (40-60 .mu.m).
[0142] Sample Preparation. For the preparation of films and solid
samples of either single compounds or blends, the
oligopeptide-polymer derivatives and/or PIB (MW 75,000) were
dissolved in either TCE or CHCl.sub.3, the solutions were stirred
at room temperature for 1-16 h, and concentrated in vacuo. The
resulting materials were dried in HV at 120.degree. C. for 3
days.
[0143] General Procedure A: Peptide Coupling. The carboxylic acid
derivative was dissolved in THF. The amine (1 equiv) was added, as
well as N-ethyldiisopropylamine (DIEA; 3 equiv) and
(benzotriazol-1-yloxy)tripyrrolidinophosphonium hexafluorophosphate
(PyBOP; 1.2 equiv). The solution was stirred for 3-16 h, and the
reaction progress was monitored by TLC. The crude product was
typically purified by precipitation into water (see General
Procedures C or D). Specific purification or sample preparation
procedures were performed before further characterization in some
cases.
[0144] General Procedure B: Fmoc Deprotection. The Fmoc-protected
amine derivatives were dissolved in CHCl.sub.3. Then, a large
excess of piperidine (.gtoreq.15 equiv) was added, and the solution
was stirred overnight. The reaction progress was monitored by TLC.
After completion of the reaction, the solvents were removed in
vacuo. Unless otherwise noted, the crude product was purified by
column chromatography.
[0145] General Procedure C: Precipitation of Compounds Soluble in
THF. After completion of the reaction affording the desired
compound, the reaction mixture was concentrated to half of its
original volume. A large excess of aqueous 1 M HCl solution was
added. The resulting precipitate was filtered off, re-dissolved in
THF and precipitated again, following the same procedure as
described above. After three repetitive precipitations, the crude
product was finally dissolved in CH.sub.2Cl.sub.2, CHCl.sub.3 or
THF. The solution was dried over MgSO.sub.4 and concentrated in
vacuo at 40.degree. C.
[0146] General Procedure D: Precipitation of Compounds Insoluble in
THF. After completion of the reaction affording the desired
compound, the reaction mixture was diluted with a large excess of
aqueous 1 M HCl solution. The precipitate was collected and
re-dispersed in THF at 60.degree. C. The product was precipitated
again using the same procedure as described above two more times.
The precipitate was finally re-dispersed in THF and concentrated in
vacuo at 40.degree. C.
2.3 Synthesis Procedures and Analytical Data for 2-3, 5-6, M0-M5
and D0-D5
[0147] Synthesis of PIB.sub.19-Ala.sub.3-Fmoc 2. Following General
Procedure A, PIB.sub.19-NH.sub.2 1 (14.55 g, 12.03 mmol) and
N-(9-fluorenylmethyloxycarbonyl)-L-alanyl-L-alanyl-L-alanine (6.0
g, 13.23 mmol) were dissolved in THF (250 mL). DIEA (6.18 mL, 36.08
mmol) and PyBOP (7.51 g, 14.43 mmol) were added. The reaction
mixture was stirred overnight. The product was precipitated
following General Procedure C. The final product (19.2 g, 97%) was
obtained as a slightly yellow wax. .sup.1H NMR (400 MHz,
CDCl.sub.3) .delta.=7.75 (d, J=7.5 Hz, 2H, aromatic H), 7.57 (d,
J=7.5 Hz, 2H, aromatic H), 7.52 (m, 1H, NH), 7.38 (d, J=7.4 Hz, 2H,
aromatic H), 7.29 (t, J=7.4 Hz, 2H, aromatic H), 7.13 (m, 1H, NH),
6.76 (m, 1H, NH), 5.84 (m, 1H, NH), 4.73-4.29 (m, 5H,
Fmoc-CO.sub.2CH.sub.2, 3 CHCH.sub.3), 4.20 (t, J=7.0 Hz, 2H,
fluorenyl CH), 3.31-3.07 (m, 2H, CH.sub.2NH), 1.81-0.53 (m, 178H,
aliphatic H, 3 CHCH.sub.3). MS (MALDI-TOF, DCTB/NaTFA 10:1): calcd
for C.sub.73H.sub.126N.sub.4O.sub.5Na: (n=10[M+Na]+) 1161.9620;
found: 1161.8517. R.sub.f: 0.45 (CH.sub.2Cl.sub.2/MeOH 10:1). DSC
(10.degree. C./min, N.sub.2) T.sub.g=-66.degree. C.
[0148] Synthesis of PIB.sub.19-Ala.sub.3-H 3. Following General
Procedure B, PIB.sub.19-Ala.sub.3-Fmoc 2 (17.2 g, 10.46 mmol) was
dissolved in CHCl.sub.3 (250 mL). Piperidine (10.35 ml, 104.57
mmol) was added, and the reaction mixture was stirred at room
temperature overnight. The crude product was purified by column
chromatography (silica gel, gradient
CH.sub.2Cl.sub.2.fwdarw.CH.sub.2Cl.sub.2/MeOH 5:1). The final
product (10 g, 67%) was obtained as a slightly yellow wax. 1H NMR
(400 MHz, CDCl.sub.3) .delta.=7.83 (d, J=7.3 Hz, 1H, NH), 7.10 (m,
1H, NH), 6.46 (m, 1H, NH), 4.45 (m, 2H, CHCH.sub.3), 3.51 (q, J=6.9
Hz, 1H, CHCH.sub.3NH.sub.2), 3.4-3.1 (m, 2H, CH.sub.2NH), 1.82-0.65
(m, 178H, aliphatic H, 3 CHCH.sub.3). MS (MALDI-TOF, DHB): calcd
for C.sub.58H.sub.116N.sub.4O.sub.3Na: (n=10 [M+Na].sup.+)
939.8940; found: 940.0981. R.sub.f: 0.15 (CH.sub.2Cl.sub.2/MeOH
10:1). DSC (10.degree. C./min, N.sub.2) T.sub.g=-68.degree. C.
[0149] Synthesis of Fmoc-Ala.sub.3-PIB.sub.40-Ala.sub.3-Fmoc 5.
Following General Procedure A,
N-(9-fluorenylmethyloxycarbonyl)-L-alanyl-L-alanyl-L-alanine (4.36
g, 9.61 mmol) and NH.sub.2-PIB.sub.40-NH.sub.2 4 (10.9 g, 4.28
mmol) were dissolved in THF (400 mL). DIEA (2.47 mL, 14.42 mmol)
and PyBOP (5.50 g, 10.58 mmol) were added. After 16 h, the crude
product was precipitated following General Procedure D. The final
product (14.18 g, 94%) was obtained as a white solid. .sup.1H NMR
(400 MHz, C.sub.2D.sub.2Cl.sub.4 at 110.degree. C.) .delta.=7.81
(d, J=7.5 Hz, 4H, aromatic H), 7.63 (d, J=7.3 Hz, 4H, aromatic H),
7.50-7.41 (m, 5H, aromatic H), 7.4-7.3 (m, 4H, aromatic H), 7.21
(m, 3H, aromatic H), 6.52 (m, 2H, NH), 6.34 (m, 2H, NH), 6.00 (m,
2H, NH), 5.16 (m, 2H, NH), 4.63-4.08 (m, 12H, 6 CHCH.sub.3, 2
Fmoc-CO.sub.2CH.sub.2, 2 fluorenyl CH), 3.37-2.94 (m, 4H,
2CH.sub.2NH), 1.94 (s, 4H, 2CH.sub.2C(CH.sub.3).sub.2Ph), 1.62-0.98
(m, 346H, aliphatic H, 6 CHCH.sub.3), 0.95 (s, 12H,
2PhC(CH.sub.3)2). MS (MALDI-TOF, DCTB): calcd for
C.sub.104H.sub.158N.sub.8O.sub.10Na: (n+m=9 [M+Na].sup.+)
1702.1993; found: 1702.3004. R.sub.f: 0.4 (CH.sub.2Cl.sub.2/MeOH
10:1). DSC (10.degree. C./min, N.sub.2) T.sub.g=-55.degree. C.
[0150] Synthesis of H-Ala.sub.3-PIB.sub.40-Ala.sub.3-H 6. Following
General Procedure B, Fmoc-Ala.sub.3-PIB.sub.40-Ala.sub.3-Fmoc 5
(11.00 g, 3.2 mmol) was dissolved in CHCl.sub.3 (200 mL).
Piperidine (200 ml, 2.02 mol) was added, and the reaction mixture
was stirred at room temperature overnight. The next day, the
solvent was evaporated in vacuo, and the mixture was washed three
times with cold heptane. The crude product was then dispersed in
DCM and concentrated in vacuo at 40.degree. C. Finally, the product
(7.76 g, 82%) was obtained as a white solid. .sup.1H NMR (400 MHz,
CDCl.sub.3 and TFA) .delta.=7.96 (m, 2H, NH), 7.71 (m, 6H,
NH2+TFA), 7.47 (m, 2H, NH), 7.38 (m, 1H, aromatic H), 7.21-7.06 (m,
3H, aromatic H), 6.88 (m, 2H, NH), 4.7-4.2 (m, 4H, CHCH.sub.3),
3.41-2.79 (m, 6H, 2 CH.sub.2NH, 1 CH.sub.2NH.sub.2), 1.84 (s, 4H,
2CH.sub.2C(CH.sub.3).sub.2Ph), 1.77-0.87 (m, 346H, aliphatic H,
6CHCH.sub.3), 0.80 (s, 12H, 2 PhC(CH.sub.3).sub.2). MS (MALDI-TOF,
DHB): calcd for C.sub.82H.sub.155N.sub.8O.sub.6: (n+m=11
[M+H].sup.+) 1348.2065; found: 1348.5963. R.sub.f: 0.05
(CH.sub.2Cl.sub.2/MeOH 10:1). DSC (10.degree. C./min, N.sub.2)
T.sub.g=-65.degree. C.
[0151] Synthesis of PIB.sub.19-Ac M0. PIB.sub.19--NH.sub.2 1 (3.85
g, 3.18 mmol) was dissolved in THF (80 mL). Acetyl chloride (0.454
mL, 6.37 mmol) and pyridine (0.642 mL, 7.96 mmol) were added. The
reaction mixture was stirred overnight. The crude product was
precipitated following General Procedure C. The final product (3.53
g, 88%) was obtained as a slightly yellow viscous oil. .sup.1H NMR
(400 MHz, CDCl.sub.3) .delta.=5.36 (m, 1H, NH), 3.47-3.08 (m, 2H,
CH.sub.2NH), 1.96 (s, 3H, C.dbd.OCH.sub.3), 1.73-0.60 (m, 169H,
aliphatic H). MS (MALDI-TOF, DCTB/NaTFA 10:1): calcd for
C.sub.51H.sub.103NONa (n=10 [M+Na].sup.+) 768.7932; found 768.9546.
R.sub.f: 0.85 (CH.sub.2Cl.sub.2/MeOH 10:1). DSC (10.degree. C./min,
N2) T.sub.g=-67.degree. C.
[0152] Synthesis of PIB.sub.19-Ala-Ac M1. Following General
Procedure A, PIB.sub.19-NH.sub.2 1 (3.69 g, 3.05 mmol) and
N-acetyl-L-alanine (0.4 g, 3.05 mmol) were dissolved in THF (100
mL). DIEA (1.57 mL, 9.15 mmol) and PyBOP (1.9 g, 3.66 mmol) were
added.
[0153] The reaction mixture was stirred overnight. The product was
precipitated following General Procedure C. The final product (3.9
g, 97%) was obtained as a slightly yellow viscous oil. .sup.1H NMR
(400 MHz, CDCl.sub.3 and TFA) .delta.=7.91 (d, J=7.6 Hz, 1H, NH),
6.87 (m, 1H, NH), 4.75-4.36 (m, 1H, CHCH.sub.3), 3.30 (m, 2H,
CH.sub.2NH), 2.16 (s, 3H, C.dbd.OCH.sub.3), 1.64-0.84 (m, 172H,
aliphatic H, 1 CHCH.sub.3). MS (MALDI-TOF, DCTB/NaTFA 10:1): calcd
for C.sub.54H.sub.108N.sub.2O.sub.2Na: (n=10 [M+Na].sup.+)
839.8303; found: 839.6927. R.sub.f: 0.55 (CH.sub.2Cl.sub.2/MeOH
10:1). DSC (10.degree. C./min, N.sub.2) T.sub.g=-65.degree. C.
[0154] Synthesis of PIB.sub.19-Ala2-Ac M2. Following General
Procedure A, PIB.sub.19--NH.sub.2 1 (17.48 g, 14.46 mmol) and
N-acetyl-L-alanyl-L-alanine (3.8 g, 18.79 mmol) were dissolved in
THF (250 mL). DIEA (4.95 mL, 28.91 mmol) and PyBOP (9.03 g, 17.35
mmol) were added. The reaction mixture was stirred overnight. The
product was precipitated following General Procedure C. The final
product (19.5 g, 97%) was obtained as a slightly yellow wax.
.sup.1H NMR (400 MHz, CDCl.sub.3) .delta.=7.44 (m, 1H, NH), 7.00
(m, 1H, NH), 6.74 (m, 1H, NH), 4.70 (m, 1H, CHCH.sub.3), 4.59 (m,
1H, CHCH.sub.3), 3.45-3.09 (m, 2H, CH.sub.2NH), 2.05 (s, 3H,
C.dbd.OCH.sub.3), 1.62-0.64 (m, 175H, aliphatic H, 2 CHCH.sub.3).
MS (MALDI-TOF, CHCA/NaTFA 1:1): calcd for
C.sub.57H.sub.113N.sub.3O.sub.3Na: (n=10 [M+Na].sup.+) 910.8674;
found: 910.7816. R.sub.f: 0.4 (CH.sub.2Cl.sub.2/MeOH 10:1). DSC
(10.degree. C./min, N2) T.sub.g=-68.degree. C., T.sub.m=170.degree.
C.
[0155] Synthesis of PIB.sub.19-Ala.sub.3-Ac M3.
PIB.sub.19-Ala.sub.3-H 3 (2.99 g, 2.1 mmol) was dissolved in THF
(150 mL). Acetyl chloride (0.3 mL, 4.2 mmol) and pyridine (0.424
mL, 5.25 mmol) were added and the reaction mixture was stirred
overnight. The next day, the product was precipitated following
General Procedure C. The final product (3.0 g, 98%) was obtained as
a slightly yellow wax. .sup.1H NMR (400 MHz, CDCl.sub.3 and TFA)
.delta.=7.58 (m, 2H, NH), 7.13 (m, 1H, NH), 6.90 (m, 1H, NH),
4.66-4.55 (m, 3H, CHCH.sub.3), 3.26 (m, 2H, CH.sub.2NH), 2.11 (s,
3H, C.dbd.OCH.sub.3), 1.70-0.57 (m, 178H, aliphatic H, 3
CHCH.sub.3). MS (MALDI-TOF, CHCA/NaTFA 1:1): calcd for
C.sub.60H.sub.118N.sub.4O.sub.4Na: (n=10 [M+Na].sup.+) 981.9045;
found: 981.8898. R.sub.f: 0.25 (CH2Cl2/MeOH 10:1). DSC (10.degree.
C./min, N2) T.sub.g=-67.degree. C.
[0156] Synthesis of PIB.sub.19-Ala.sub.4-Ac M4. Following General
Procedure A, PIB.sub.19-Ala.sub.3-H 3 (2.0 g, 1.41 mmol) and
N-acetyl-L-alanine (184.36 mg, 1.41 mmol) were dissolved in THF
(200 mL). DIEA (0.722 mL, 4.22 mmol) and PyBOP (877.98 mg, 1.69
mmol) were added. The reaction mixture was stirred overnight. The
product was precipitated following General Procedure C. The final
product (2.0 g, 90%) was obtained as a white rubber. .sup.1H NMR
(400 MHz, CDCl.sub.3 and TFA) .delta.=7.83 (d, J=7.4 Hz, 1H, NH),
7.60 (m, 1H, NH), 7.34 (d, J=6.1 Hz, 1H, NH), 7.27 (m, 1H, NH),
6.80 (m, 1H, NH), 4.89-4.22 (m, 4H, CHCH.sub.3), 3.45-3.09 (m, 2H,
CH.sub.2NH), 2.14 (s, 3H, C.dbd.OCH.sub.3), 1.65-0.74 (m, 181H,
aliphatic H, 4 CHCH.sub.3). MS (MALDI-TOF, CHCA/NaTFA 1:1): calcd
for C.sub.63H.sub.123N.sub.5O.sub.5Na: (n=10 [M+Na].sup.+)
1052.9416; found: 1052.8131. R.sub.f: 0.1 (CH.sub.2Cl.sub.2/MeOH
10:1). DSC (10.degree. C./min, N2) T.sub.g=-66.degree. C.
[0157] Synthesis of PIB.sub.19-Ala.sub.5-Ac M5. Following General
Procedure A, PIB.sub.19-Ala.sub.3-H 3 (2.0 g, 1.41 mmol) and
N-acetyl-L-alanyl-L-alanine (284.29 mg, 1.41 mmol) were dissolved
in THF (200 mL). DIEA (0.722 mL, 4.22 mmol) and PyBOP (877.98 mg,
1.69 mmol) were added. The reaction mixture was stirred overnight.
The product was precipitated following General Procedure C. The
final product (2.2 g, 95%) was obtained as a white rubber. .sup.1H
NMR (400 MHz, C.sub.2D.sub.2Cl.sub.4 and TFA at 65.degree. C.)
.delta.=7.44 (d, J=6.9 Hz, 1H, NH), 7.16 (m, 3H, NH), 6.81 (m, 1H,
NH), 6.56 (m, 1H, NH), 4.67-4.42 (m, 5H, CHCH.sub.3), 3.45-3.09 (m,
2H, CH.sub.2NH), 2.17 (s, 3H, C.dbd.OCH.sub.3), 1.84-0.68 (m, 184H,
aliphatic H, 5 CHCH.sub.3). MS (MALDI-TOF, CHCA): calcd for
C.sub.66H.sub.128N.sub.6O.sub.6Na: (n=10 [M+Na].sup.+) 1123.9788;
found: 1124.3684. R.sub.f: 0.05 (CH.sub.2Cl.sub.2/MeOH 10:1). DSC
(10.degree. C./min, N2) T.sub.g=-69.degree. C.
[0158] Synthesis of Ac-PIB.sub.40-Ac D0.
NH.sub.2-PIB.sub.40-NH.sub.2 4 (5.00 g, 1.96 mmol) was dissolved in
THF (200 mL). Acetyl chloride (0.80 mL, 11.22 mmol) and pyridine
(0.9 mL, 11.22 mmol) were added, and the solution was stirred
overnight. The next day, the crude product was precipitated
following General Procedure C. The final product (4.4 g, 85%) was
obtained as a colourless viscous oil. .sup.1H NMR (400 MHz,
CDCl.sub.3) .delta.=7.37 (m, 1H, aromatic H), 7.23-7.06 (m, 3H,
aromatic H), 5.44 (m, 2H, NH), 3.2-2.9 (m, 4H, 2 CH.sub.2NH), 1.98
(s, 6H, C.dbd.OCH.sub.3), 1.84 (s, 4H,
2CH.sub.2C(CH.sub.3).sub.2Ph), 1.75-0.86 (m, 328H, aliphatic H),
0.80 (s, 12H, 2 PhC(CH.sub.3).sub.2). MS (MALDI-TOF, DCTB): calcd
for C.sub.76H.sub.144N.sub.2O.sub.2Na: (n+m=13 [M+Na].sup.+)
1140.1120; found: 1139.5068. R.sub.f: 0.7 (CH.sub.2Cl.sub.2/MeOH
10:1). DSC (10.degree. C./min, N2) T.sub.g=-55.degree. C.
[0159] Synthesis of Ac-Ala-PIB.sub.40-Ala-Ac D1. Following General
Procedure A, N-acetyl-L-alanine (810.0 mg, 6.17 mmol) and
NH.sub.2-PIB.sub.40-NH.sub.2 4 (7.00 g, 2.75 mmol) were dissolved
in THF (200 mL). DIEA (1.58 mL, 9.26 mmol) and PyBOP (3.53 g, 6.79
mmol) were added. After 16 h, the crude product was precipitated
following General Procedure C. The final product (5.12 g, 66%) was
obtained as a yellow glue. .sup.1H NMR (400 MHz, CDCl.sub.3)
.delta.=7.37 (m, 1H, aromatic H), 7.23-7.06 (m, 3H, aromatic H),
6.17 (m, 4H, NH), 4.45 (m, 2H, CHCH.sub.3), 3.25-2.9 (m, 4H, 2
CH.sub.2NH), 2.00 (s, 6H, C.dbd.OCH.sub.3), 1.84 (s, 4H,
2CH.sub.2C(CH.sub.3).sub.2Ph), 1.76-0.85 (m, 334H, aliphatic H, 2
CHCH.sub.3), 0.80 (s, 12H, 2 PhC(CH.sub.3).sub.2). MS (MALDI-TOF,
DCTB): calcd for C.sub.86H.sub.162N.sub.4O.sub.4Na: (n+m=14
[M+Na].sup.+) 1338.2488; found: 1338.2838. R.sub.f: 0.5
(CH.sub.2Cl.sub.2/MeOH 10:1). DSC (10.degree. C./min, N.sub.2)
T.sub.g=-54.degree. C.
[0160] Synthesis of Ac-Ala.sub.2-PIB.sub.40-Ala.sub.2-Ac D2.
Following General Procedure A, N-acetyl-L-alanyl-L-alanine (700.0
mg, 3.46 mmol) and NH.sub.2-PIB.sub.40--NH.sub.2 4 (3.93 g, 1.54
mmol) were dissolved in THF (300 mL). DIEA (0.89 mL, 5.19 mmol) and
PyBOP (1.98 g, 3.81 mmol) were added, and the reaction mixture was
stirred overnight. The next day, the crude product was precipitated
following General Procedure D. The final product (4.0 g, 89%) was
obtained as an off-white solid. .sup.1H NMR (400 MHz, CDCl.sub.3
and TFA) .delta.=7.64-7.3 (m, 6H, NH), 7.37 (m, 1H, aromatic H),
7.23-7.06 (m, 3H, aromatic H), 4.7-4.5 (m, 4H, CHCH.sub.3),
3.25-2.9 (m, 4H, 2 CH.sub.2NH), 2.07 (s, 6H, C.dbd.OCH.sub.3), 1.84
(s, 4H, 2CH.sub.2C(CH.sub.3).sub.2Ph), 1.77-0.86 (m, 300H,
aliphatic H, 4 CHCH.sub.3), 0.80 (s, 12H, 2 PhC(CH.sub.3).sub.2).
MS (MALDI-TOF, DCTB): calcd for C.sub.92H.sub.172N.sub.6O.sub.6Na:
(n+m=14 [M+Na].sup.+) 1480.3231; found: 1480.6152. R.sub.f: 0.4
(CH.sub.2Cl.sub.2/MeOH 10:1). DSC (10.degree. C./min, N2)
T.sub.g=-57.degree. C., T.sub.m=178.degree. C.
[0161] Synthesis of Ac-Ala.sub.3-PIB.sub.40-Ala.sub.3-Ac D3.
H-Ala.sub.3-PIB.sub.40-Ala.sub.3-H 6 (1.50 g, 0.50 mmol) was
dissolved in THF (200 mL). Acetyl chloride (0.16 mL, 2.23 mmol) and
pyridine (0.18 mL, 2.23 mmol) were added, and the reaction mixture
was stirred overnight. The next day, the crude product was
precipitated General Procedure D. The final product (1.43 g, 92%)
was obtained as an off-white solid. .sup.1H NMR (400 MHz,
CDCl.sub.3 and TFA) .delta.=8.1-7.5 (m, 5H, NH), 7.38 (m, 2H,
aromatic H, NH), 7.21-7.06 (m, 3H, aromatic H), 6.83 (m, 2H, NH),
4.7-4.3 (m, 6H, CHCH.sub.3), 3.35-2.95 (m, 4H, 2 CH.sub.2NH), 2.12
(s, 6H, C.dbd.OCH.sub.3), 1.84 (s, 4H,
2CH.sub.2C(CH.sub.3).sub.2Ph), 1.77-0.86 (m, 346H, aliphatic H, 6
CHCH.sub.3), 0.80 (s, 12H, 2 PhC(CH.sub.3).sub.2). MS (ESI-TOF):
calcd for C.sub.74H.sub.135N.sub.8O.sub.8Na: (n=8 [M+H+Na].sup.2+)
643.5145; found: 646.5165. R.sub.f: 0.3 (CH.sub.2Cl.sub.2/MeOH
10:1). DSC (10.degree. C./min, N2) T.sub.g=-55.degree. C.
[0162] Synthesis of Ac-Ala.sub.4-PIB.sub.40-Ala.sub.4-Ac D4.
Following General Procedure A, N-acetyl-L-alanine (146.0 mg, 1.11
mmol) and H-Ala.sub.3-PIB.sub.40-Ala.sub.3-H 6 (1.50 g, 0.50 mmol)
were dissolved in THF (200 mL). DIEA (215.83 mL, 1.67 mmol) and
PyBOP (637.3 mg, 1.22 mmol) were added. After 16 h, the crude
product was precipitated following General Procedure D. The final
product (1.49 g, 91%) was obtained as an off-white solid. .sup.1H
NMR (400 MHz, CDCl.sub.3 and TFA) .delta.=8.1-7.5 (m, 6H, NH), 7.38
(m, 2H, aromatic H, NH), 7.21-7.06 (m, 4H, 3 aromatic H, NH), 6.83
(m, 2H, NH), 4.7-4.35 (m, 8H, CHCH.sub.3), 3.35-2.95 (m, 4H, 2
CH.sub.2NH), 2.12 (s, 6H, C.dbd.OCH.sub.3), 1.84 (s, 4H,
2CH.sub.2C(CH.sub.3).sub.2Ph), 1.77-0.86 (m, 352H, aliphatic H, 8
CHCH.sub.3), 0.80 (s, 12H, 2 PhC(CH.sub.3).sub.2). MS (ESI-TOF):
calcd for C.sub.92H.sub.170N.sub.10O.sub.10: (n+m=11 [M+2H].sup.2+)
787.6545; found: 787.1820. R.sub.f: 0.3 (CH.sub.2Cl.sub.2/MeOH
10:1). DSC (10.degree. C./min, N2) T.sub.g=-55.degree. C.
[0163] Synthesis of Ac-Ala.sub.5-PIB.sub.40-Ala.sub.5-Ac D5.
Following General Procedure A, N-acetyl-L-alanyl-L-alanine (195.1
mg, 0.96 mmol) and H-Ala.sub.3-PIB.sub.40-Ala.sub.3-H 6 (1.30 g,
0.43 mmol) were dissolved in THF (200 mL). DIEA (0.25 mL, 1.45
mmol) and PyBOP (552.3 mg, 1.06 mmol) were added. After 16 h, the
crude product was precipitated following General Procedure D. The
final product (1.30 g, 88%) was obtained as off-white solid.
.sup.1H NMR (400 MHz, CDCl.sub.3 and TFA) .delta.=8.1-7.5 (m, 10H,
NH), 7.38 (m, 2H, aromatic H, NH), 7.21-7.06 (m, 3H, aromatic H),
6.85 (m, 1H, NH), 4.7-4.2 (m, 10H, CHCH.sub.3), 3.35-2.95 (m, 4H, 2
CH.sub.2NH), 2.12 (s, 6H, C.dbd.OCH.sub.3), 1.84 (s, 4H,
2CH.sub.2C(CH.sub.3).sub.2Ph), 1.77-0.86 (m, 358H, aliphatic H, 10
CHCH.sub.3), 0.80 (s, 12H, 2PhC(CH.sub.3).sub.2). MS (ESI-TOF):
calcd for C.sub.98H.sub.178N.sub.12O.sub.12Na: (n+m=11
[M+Na].sup.+) 1739.3613; found: 1740.5134. R.sub.f: 0.25
(CH.sub.2Cl.sub.2/MeOH 10:1). DSC (10.degree. C./min, N.sub.2)
T.sub.g=-54.degree. C.
[0164] Synthesis of PS.sub.15-NH.sub.2 7. Styrene (12.19 g, 117.1
mmol) was freshly distilled from CaH.sub.2 prior to use. A
rigorously dried 250 mL Schlenk flask was filled with dry
cyclohexane (50 mL), and sec-butyl lithium (5.6 mL, 7.8 mmol, 1.4 M
solution in cyclohexane) was added slowly via a syringe at
10.degree. C. Under vigorous stirring, styrene was added via a
syringe as fast as possible, leading to a yellow to orange color of
the solution. The cooling bath was removed and the mixture was
stirred for 1 h. Then, 2.5 excess of
1-.beta.-bromopropyl)-2,2,5,5-tetramethyl-1-aza-2,5-disilacyclopentane
(5.47 g, 19.5 mmol) dissolved in dry THF (20 mL) was added via a
syringe, causing an immediate decoloration. The mixture was stirred
for 2 h at room temperature, then was concentrated in vacuo and
taken up in THF (100 mL). Then, 1M HCl (35 mL) was added, and
stirring was continued overnight. The mixture was concentrated in
vacuo, taken up in CH.sub.2Cl.sub.2, washed twice with 1M KOH and
once with sat. NaCl solution. The combined organic phases were
dried over MgSO.sub.4, filtered, and concentrated in vacuo. The
crude product was purified by column chromatography (silica gel,
gradient CH.sub.2Cl.sub.2.fwdarw.CH.sub.2Cl.sub.2/MeOH 20:1). The
amine terminated polystyrene (9.2 g, 70%) was obtained as a white
solid. .sup.1H NMR (400 MHz, CDCl.sub.3) .delta.=7.25-6.3 (m, 75H,
Ph-H), 2.48 (m, 2H, CH.sub.2NH.sub.2), 2.36-0.5 (m, 58H, 15
CH.sub.2CHPh, 3 CH.sub.2, 1 CHCH.sub.3, 2 CH.sub.3). MS (ESI-TOF):
calcd for C.sub.127H.sub.138N: (n=15 [M+H].sup.+) 1678.0858; found:
1677.9187. R.sub.f: 0.35 (CH.sub.2Cl.sub.2/MeOH 10:1).
[0165] Synthesis of PS.sub.15-Ala.sub.3-Fmoc 8. Following General
Procedure A, PS.sub.15-NH.sub.2 (5.04 g, 3.01 mmol) and
N-(9-fluorenylmethyloxycarbonyl)-L-alanyl-L-alanyl-L-alanine (1.50
g, 3.31 mmol) were dissolved in THF (250 mL). DIEA (1.54 mL, 9.02
mmol) and PyBOP (1.88 g, 3.61 mmol) were added. The reaction
mixture was stirred overnight. The product was precipitated
following General Procedure C. The final product (6.0 g, 95%) was
obtained as a pinkish solid. .sup.1H NMR (400 MHz, CDCl.sub.3)
.delta.=7.76 (d, J=7.8 Hz, 2H, aromatic H), 7.55 (d, J=7.5 Hz, 2H,
aromatic H), 7.41 (t, J=7.4 Hz, 2H, aromatic H), 7.32 (t, J=7.3 Hz,
2H, aromatic H), .delta.=7.3-6.3 (m, 75H, Ph-H), 6.23 (s, 1H, NH),
5.20 (s, 1H, NH), 4.50-4.29 (m, 5H, Fmoc-CO.sub.2CH.sub.2, 3
CHCH.sub.3), 4.17 (m, 2H, fluorenyl CH), 3.02 (m, 2H, CH.sub.2NH),
2.5-0.5 (m, 67H, 15 CH.sub.2CHPh, 5 CH.sub.2, 3 CHCH.sub.3,
C.dbd.OCH.sub.3).
[0166] Synthesis of PS.sub.15-Ala.sub.3-H 9. Following General
Procedure B, PS.sub.15-Ala.sub.3-Fmoc (5.4 g, 2.56 mmol) was
dissolved in CHCl.sub.3 (250 mL). Piperidine (5.06 ml, 51.11 mmol)
was added, and the reaction mixture was stirred at room temperature
overnight. The crude product was purified by column chromatography
(silica gel, gradient CH.sub.2Cl.sub.2.fwdarw.CH.sub.2Cl.sub.2/MeOH
50:1). The final product (3.55 g, 75%) was obtained as yellowish
powder. .sup.1H NMR (400 MHz, CDCl.sub.3) .delta.=7.70 (m, 1H, NH),
.delta.=7.3-6.3 (m, 75H, Ph-H), 5.90 (m, 1H, NH), 4.31 (m, 2H,
CHCH.sub.3), 3.45 (m, 1H, CHCH.sub.3NH.sub.2), 3.02 (m, 2H,
CH.sub.2NH), 2.5-0.5 (m, 67H, 15 CH.sub.2CHPh, 5 CH.sub.2, 3
CHCH.sub.3, C.dbd.OCH.sub.3).
[0167] Synthesis of PS.sub.15-Ala.sub.2-Ac S2. Following General
Procedure A, PS.sub.15-NH.sub.2 (3.0 g, 1.79 mmol) and
N-acetyl-1-alanyl-1-alanine (0.36 g, 1.79 mmol) were dissolved in
THF (250 mL). DIEA (1.22 mL, 7.15 mmol) and PyBOP (1.4 g, 2.68
mmol) were added. The reaction mixture was stirred overnight. The
product was precipitated following General Procedure C. The final
product (2.2 g, 67%) was obtained as pinkish solid. .sup.1H NMR
(400 MHz, CDCl.sub.3) .delta.=7.3-6.3 (m, 75H, Ph-H), 6.10 (d,
J=6.1 Hz, 1H, NH), 5.79 (s, 1H, NH), 4.42 (m, 1H, CHCH.sub.3), 4.29
(m, 1H, CHCH.sub.3), 3.02 (m, 2H, CH.sub.2NH), 2.5-0.5 (m, 67H, 15
CH.sub.2CHPh, 5 CH.sub.2, 3 CHCH.sub.3, C.dbd.OCH.sub.3).
[0168] Synthesis of PS.sub.15-Ala.sub.3-Ac S3.
PS.sub.15-Ala.sub.3-H (1.2 g, 0.63 mmol) was dissolved in THF (50
mL). Acetyl chloride (136 mL, 1.9 mmol) and pyridine (0.205 mL,
2.54 mmol) were added and the reaction mixture was stirred
overnight. The next day, the product was precipitated following
General Procedure C. The final product (1.2 g, 98%) was obtained as
white solid powder. .sup.1H NMR (400 MHz, CDCl.sub.3 and TFA)
.delta.=7.3-6.3 (m, 75H, Ph-H), 5.0-4.6 (m, 3H, NH), 3.75 (m, 3H,
CHCH.sub.3), 3.03 (m, 2H, CH.sub.2NH), 2.5-0.5 (m, 70H, 15
CH.sub.2CHPh, 5 CH.sub.2, 4 CHCH.sub.3, C.dbd.OCH.sub.3).
[0169] Synthesis of PS.sub.15-Ala.sub.4-Ac S4. Following General
Procedure A, PS.sub.15-Ala.sub.3-H (602.8 mg, 0.319 mmol) and
N-acetyl-1-alanine (46.0 mg, 0.351 mmol) were dissolved in THF (30
mL). DIEA (0.17 mL, 0.957 mmol) and PyBOP (602.9 mg, 1.69 mmol)
were added. The reaction mixture was stirred overnight. The product
was precipitated following General Procedure C. The final product
(0.6 g, 95%) was obtained as a white solid powder. .sup.1H NMR (400
MHz, CDCl.sub.3 and TFA) .delta.=7.3-6.3 (m, 75H, Ph-H), 5.5-4.6
(m, 4H, CHCH.sub.3), 3.03 (m, 2H, CH.sub.2NH), 2.5-0.5 (m, 73H, 15
CH.sub.2CHPh, 5 CH.sub.2, 5 CHCH.sub.3, C.dbd.OCH.sub.3).
TABLE-US-00001 TABLE 1 Representative rheological data of M0-M5,
D0-D2, polyisobutylenes of different molecular weights, as well as
the blends M2/D2 99:1 (Example 1), M2/D2 95:5 (example 2), M2/D2
9:1 (Example 3), M2/D2 7:3 (Example 4), M2/D2 5:5 (Example 5), and
M2/D2 1:9 (Example 6); storage moduli G', loss moduli G'', loss
factors tan .delta., and shear viscosities |.eta.*|. Materials
G'/Pa G''/Pa tan .delta. |.eta.*|Pa s) M0 0.22 139 623 139 M1 249
1250 5.02 1270 M2 42800 14300 0.334 45200 M3 251000 42900 0.171
254000 M4 510000 40200 0.079 511000 M5 633000 63900 0.101 636000 D0
25 2690 109 2690 D1 12300 70700 5.75 71800 D2 2000000 117000 0.059
2010000 PIB-NH.sub.2 (MW 1200) 0.19 32 168 32 H.sub.2N-PIB-NH.sub.2
(MW 1200) 4.99 234 46.9 235 PIB (MW 35'000) 13400 24600 1.84 28000
PIB (MW 75'000) 64500 42800 0.664 77500 PIB (MW 200'000) 231000
25270 0.109 232000 PIB (MW 425'000) 188100 11840 0.063 188400 M2/D2
99:1 123000 21800 0.177 125000 M2/D2 95:5 564000 47310 0.084 565000
M2/D2 9:1 495000 29610 0.060 496000 M2/D2 7:3 798000 52780 0.066
800000 M2/D2 5:5 1760000 124000 0.070 1760000 M2/D2 1:9 1500000
93880 0.063 1500000
TABLE-US-00002 TABLE 2 Shift factors log a.sub.T and log b.sub.T
and activation energies E.sub.a obtained from the Arrhenius
equation for the rheological measurements used in the
time-temperature superposition master curves of polyisobutylenes of
different molecular weights, pure M2 and D1, as well as the blends
M2/D1 1:4 (Example 7), M2/PIB (MW 75'000) 5:5 (Example 8), M2/PIB
(MW 35'000) 5:5 (Example 9), and M2/D2/PIB(MW 35'000) 4:1:5
(Example 10). Materials T/.degree. C. log a.sub.T log b.sub.T
E.sub.a/J mol.sup.-1 PIB (200k) -45 4.474 0.119 1055.8 -10 1.677
0.028 25 0 0 105 -2.199 -0.111 PIB (75k) -45 3.889 -0.17 986.91 -10
1.684 0.021 25 0 0 65 -1.317 -0.0048 105 -2.271 -0.099 PIB (35k)
-45 4.492 0.133 1080.16 -10 1.689 0.04 25 0 0 65 -1.302 -0.025 105
-2.356 -0.166 M2 -45 4.317 -0.089 1040.6 -10 1.65 -0.066 25 0 0 105
-2.255 0.015 D1 -25 3.901 -0.033 1735.8 -10 2.991 0.137 25 0 0 -25
3.93 -0.024 M2/D1 1:4 -10 2.998 0.112 1585.7 25 0 0 65 -1.918 0.115
105 -3.928 -0.119 -25 3.901 -0.033 M2/PIB (75k) 5:5 -45 4.305 0.033
1078.1 -10 1.701 0.035 25 0 0 65 -1.282 -0.024 105 -2.573 -0.216
M2/PIB (35k) 5:5 -45 4.075 -0.145 1111.9 -10 1.478 -0.065 25 0 0 65
-1.901 -0.364 105 -2.881 -0.464 M2/D2/PIB (35k) 4:1:5 -45 4.4 0.037
1136.3 -10 1.6 0.057 25 0 0 65 -1.841 -0.186 105 -2.677 -0.265
Smactane -45 4.342 -0.059 1194.3 -10 1.746 0.05 25 0 0 65 -1.8
-0.035 105 -3.208 -0.045
TABLE-US-00003 TABLE 3 Storage moduli G' and loss moduli G'' used
as input in the finite element simulations, and the resonance
frequencies .omega. and modal damping ratios .xi. resulting from
these simulations for the reference materials Smactane .TM., PIB
(MW 200'000), Soundcoat .TM. Dyad 601, and 3M ISD 130, as well as
M2 and the blends M2/PIB (MW 75'000) 5:5 (Example 8), M2/PIB (MW
35'000) 5:5 (Example 9), and M2/D2/PIB(MW 35'000) 4:1:5 (Example
10). Materials G'/MPa G''/MPa .omega./Hz .xi. Smactane .TM. 1.01
0.47 36.1 1.8% PIB (200k) 0.317 0.148 35 0.70% Soundcoat Dyad
601.sup.2 1.26 0.756 36.5 2.6% 3M ISD 130.sup.1 0.094 0.036 34.6
0.20% M2 0.424 0.524 35.3 2.3% M2/PIB (75k) 5:5 0.509 0.442 35.396
1.89% M2/PIB (35k) 5:5 0.377 0.371 35.1 1.7% M2/D2/PIB (35k) 4:1:5
1.48 0.83 36.8 2.8%
TABLE-US-00004 TABLE 4 Calculation of the low-temperature damping
properties. Loss modulus G'' and loss factor tan .delta. at 200
rad/s (.apprxeq.32 Hz) from a rheological frequency sweep at room
temperature, first resonance frequency .omega. in the `forced
vibration tests` on sandwich structures, modal damping ratio .zeta.
at room temperature, and calculated dissipated energy W.sub.d at
-45.degree. C. and 25 rad/s for steel, Smactane .TM., PIB (MW
200'000), M2, M2/PIB (MW 35'000) 5:5 (Example 9), and M2/D2/PIB(MW
35'000) 4:1:5 (Example 10). G''/MPa tan .delta. .omega./Hz .zeta.
W.sub.d/J m.sup.-3 Steel -- -- 32.2 0.4% -- Smactane .TM. 0.47 0.45
34.9 2.9% 5.98 PIB (MW 200'000) 0.15 0.45 32.4 1.1% 3.95 M2 0.54
1.3 33.4 3.2% 13.29 M2/PIB (35k) 5:5 0.37 1.0 33.4 2.5% 11.09
M2/D2/PIB (35k) 4:1:5 0.83 0.55 35.6 3.4% 18.95
TABLE-US-00005 TABLE 5 Shift factors log a.sub.T and log b.sub.T
for the rheological measurements used in the time-temperature
superposition master curves of blends M3/S3/PIB (MW 35'000) 9:3:12
(Example 13). Materials T/.degree. C. log a.sub.T log b.sub.T
M3/S3/PIB (35k) -45 4.220 -0.220 -25 2.856 -0.115 -10 1.600 -0.170
25 0 0 65 -2.454 -0.033
* * * * *