**Interpenetrating Networks of Elastomers**

**Exhibiting 300% Electrically-Induced Area Strain**

**Soon Mok Ha^1^, Wei Yuan^1^, Qibing Pei\*^1^, Ron Pelrine^2^ and Scott
Stanford^2^**

**^1^Department of Materials Science and Engineering, UCLA, Los Angeles,
CA, USA 90095;**

**^2^SRI International, 333 Ravenswood Avenue, Menlo Park, CA, USA
942025**

> **E-mail : <soonmok@ucla.edu>, <weiyuan@seas.ucla.edu>,
> <qpei@seas.ucla.edu>, <ronald.pelrine@sri.com> and
> <scott.stanford@sri.com>**

**Abstract**

Mechanical prestrain is generally required for most electroelastomers to
obtain high electromechanical strain and high elastic energy density.
However, prestrain can cause several serious problems, including a large
performance gap between the active materials and packaged actuators,
instability at interfaces between the elastomer and prestrain-supporting
structure, and stress relaxation. Difunctional and trifunctional liquid
additives were introduced into 400% biaxially prestrained acrylic films
and subsequently cured to form the second elastomeric network. The goal
of this research was to determine the effect of different functional
additives and concentrations on the microstructure, the mechanical
properties, and the actuation of composite films. In the as-obtained
Interpenetrating Polymer Networks (IPN), the additive network can
effectively support the prestrain of the acrylic films and as a result,
eliminate the external prestrain-supporting structure. However, the
large amount of additive used to completely preserve prestrain was found
to make the films too stiff, causing damage to IPN composite films.
Furthermore, the interpenetrating network formed from a trifunctional
monomer is more effective than that formed from the difunctional monomer
in supporting the high tension of the VHB network. This high efficiency
trifunctional additive leads to the enhancement of the breakdown field,
due to less damage on the microstructure. The IPN composite films
without external prestrain exhibit electrically-induced strains up to
300% in area, comparable to those of VHB 4910 films under high prestrain
conditions.

**1. Introduction**

In recent studies of large electromechanical strain from
electroelastomers, based on acrylic copolymer elastomers (e.g., 3M VHB
4910), it was found that prestrain significantly enhances performance by
increasing the dielectric breakdown strength \[1-6\]. Because an
electrically-induced strain is proportional to the square of the applied
electric field, the maximum strain is significantly increased as the
breakdown field is enhanced. Mechanisms to account for the enhancement
of the dielectric breakdown strength have not been well understood; they
have been attributed to the increased probability of hot electrons
colliding with polymer chains realigned in parallel to the film
surfaces. Prestrain is a key factor in transferring electroelastomer
performance. Prestrain realigns defects, such as fibrous impurities,
non-spherical voids, and gel particles that may be responsible for
premature dielectric breakdown. However, the overall performance of the
electroelastomer is observed to substantially decrease when packaged to
support the required high prestrain. Disparities in performance between
the active material and the packaged actuators are due to the
prestrain-supporting structures, which significantly occupy more space
and weigh more than the acrylic films themselves \[7,8\]. In addition,
the lifetimes of the actuators are limited due to the stress
concentration at the interfaces between the soft polymer film and the
rigid supporting structure. The shock tolerance of the actuators is also
reduced because of the rigid structural components introduced. The
prestrained films exhibit stress relaxation which affects the subsequent
actuation \[6\]. Consequently, it is highly desirable to eliminate
mechanical prestrain while retaining its performance benefits.

Recently, we reported interpenetrating polymer networks (IPN) as new
electroelastomers that exhibit high actuation performance without
mechanical prestrain \[9,10\]. An IPN is defined as a combination of two
or more polymers that form a network, in which at least one polymer is
polymerized and/or crosslinked in the immediate presence of the others
\[11\]. In this approach, acrylic elastomers were used as the initial
elastomeric network in order to take advantage of the high dielectric
breakdown field obtained in highly prestrained acrylic elastomer films.
Polymerizable and crosslinkable liquid monomers were introduced into
highly-prestrained acrylic films and cured to form the second
elastomeric network. When the interpenetrating elastomeric network films
were allowed to relax at zero external stress, the acrylic network was
observed to contract, compressing the additive network. The resulting
free-standing films consist of two networks in balance; one in high
tension and the other in high compression. Consequently, the molecular,
nanometer, and micrometer-scale alignments may be preserved in the
resulting free-standing films. A schematic explanation of this approach
is shown in Figure 1.

**Figure** 1. Schematic illumination of the fabrication process of an
IPN electroelastomer film: (a) an acrylic film before processing, (b)
after 400 % biaxial prestrain, the film area expands by 24 folds, (c)
curable additives are added into the prestrained film and cured, forming
an interpenetrating network of a highly crosslinked polymer, (d) after
external stress is removed, the interpenetrating network preserves most
of the prestrain of the acrylic film

In this study, two functional acrylic monomers, 1,6 hexanediol
diacrylate (HDDA) and trimethylolpropane trimethacrylate (TMPTMA), are
used to form the interpenetrating network. The density of crosslinks in
the interpenetrating network depends on the monomer type and the degree
of functionality. At equal weight fractions, TMPTMA, a trifunctional
acrylate, has a concentration of crosslinkable double bonds that is
33.68% higher than that of the HDDA, a difunctional acrylate. The
network formed from the trifunctional monomer is more effective than
that formed from the difunctional monomer in supporting the high tension
of the VHB network in the IPN composite films.

**2. Experimental**

*2.1. Materials*

An acrylic copolymer elastomer, which is commercially available as a 1
mm-thick adhesive film (3M, VHB 4910), was used as the electroelastomer
network. The film is extremely compliant and has a quite uniform
surface. In forming an additive network in the acrylic network, a
difunctional acrylate monomer or a trifunctional methacrylate monomer
was used. The thermally-curable 1,6-hexanediol diacrylate (HDDA), was
purchased from Sartomer Co. It has two acrylate functional groups per
molecule, a molecular weight of 226 Dalton, and a viscosity of 9 cps at
room temperature. The thermally-curable, trifunctional
trimethylolpropane trimethacrylate (TMPTMA) purchased from Sigma
Aldrich, has a molecular weight of 338 Dalton. The chemical structures
of the HDDA and TMPTMA are shown in Figure 2. Benzoyl peroxide, was used
as a thermally-activated initiator for free-radical polymerization.
Benzoyl peroxide is soluble in HDDA and TMPTMA.

![](media/image1.png){width="4.747916666666667in"
height="0.5194444444444445in"}

\(a\) 1,6-Hexanediol diacrylate (HDDA) monomer

![](media/image2.png){width="2.6493055555555554in"
height="1.7958333333333334in"}

b\) Trimethylolpropane trimethacrylate (TMPTMA) monomer

**Figure 2**. Chemical structures of difunctional and trifunctional
acrylic monomers

*2.2. Preparation of IPN in highly prestrained elastomer films***.**

The fabrication of IPN composite films began with the biaxial stretching
of acrylic films. The structures to support the stretched films are made
of Mesonite to minimize thermal deformation. The films were stretched to
5 times its original length and width, i.e., 400% biaxial prestrain. The
film thickness after the prestrain was 40 μm for VHB 4910. The
diacrylate and trimethacrylate monomers were added into the prestrained
VHB films. Benzoyl peroxide, a thermal initiator for free-radical
polymerization, was also added. The films were then placed in a vacuum
oven and heated up to 85 ^o^C to induce the free-radical polymerization
of the monomers. Any nonuniform heating or additive degradation gave
rise to stress concentration near the edges. Therefore, the edges of the
composite films were trimmed to eliminate this problem.

*2.3. Analysis of microstructure and FTIR spectroscopy*

The microstructure of the IPN formed films was observed by a field
emission scanning electron microscope (S4700, Hitatch Inc.). The samples
were prepared and attached onto high-density polyethylene. To improve
the image quality, the sample surface was coated up to 10 nm with gold
by sputtering. FTIR spectroscopy was used to determine the chemical
structure change during the thermal curing process. Infrared spectra of
IPN electroelastomer films were measured with FTIR spectroscopy (M2200,
MIDAC Corporation) in the wavenumber range of 4000-400 cm^-1^.

*2.4. Measurement of mechanical properties*

A simple setup was built to measure the biaxial deformation of clamped
electroelastomer films under hydraulic load. Acrylic films after thermal
polymerization of admixed diacrylate were carefully mounted onto one end
of an acrylic tube with an inner diameter of 25.4 mm and a tube length
of 1.22 m. Water was added into the tube to provide a hydraulic pressure
to biaxially stress the acrylic film \[12\]. Deflection of the central
point in the acrylic film was measured by a non-contacting technique
using a CCD camera and image analysis software. Care was taken to
prevent delamination of the IPN formed film from the acrylic tube. The
deflection was used to calculate the biaxial strain of the acrylic film
by the hydraulic pressure. The thickness of the IPN films was measured
using a digital thickness gauge (Mitutoyo).

*2.5. Actuator fabrication*

Diaphragm actuators were fabricated from free-standing composite films.
Conductive carbon grease was smeared onto both sides of the film. Thin
strips of aluminum foil were used as electrical contacts.
Electrically-induced strain was monitored using a CCD camera to capture
the changes in diameter and height. Strain in area was computed from
geometric relations found between the diameter and height of the bulge.
High DC voltages up to 10 kV were supplied from a high voltage power
source built in-house. When expanded at a high driving voltage, the IPN
films bulged or wrinkled without directional preference. To measure the
strain, a small pneumatic pressure was applied to deform the expanding
films into a spherical geometry.

**3. Results and Discussion**

*3.1. Preserved prestrain with curable additive*

The acrylic copolymers in the 3M VHB films are very soft with a Young's
modulus on the order of 500 kPa. This low modulus is due to the
long-chain and branched alkyl groups attached to the acrylate polymer
chains. The highly crosslinked polymer of HDDA, with a modulus on the
order of 930 MPa \[13\], is more rigid than the VHBs. However, the
flexible hexylene segment between the two acrylate end groups of HDDA
allows the polymer to contract or stretch within a certain range. The
modulus of poly(TMPTMA) is expected to be higher than that of the
poly(HDDA), as the trifunctional TMPTMA results in a higher density of
crosslinks. Therefore, the same percentage of preserved prestrain
obtained by poly(HDDA) can be accomplished by using a smaller amount of
poly(TMPTMA).

In the experiment, various amounts of HDDA and TMPTMA were added
separately into a highly prestrained VHB film. Upon curing, an
interpenetrating network of poly(HDDA) or poly(TMPTMA) formed in the
highly prestrained VHB network. As the IPN composite film was released
from the prestrain structure, a contraction occurred in the prestrained
VHB network, compressing the poly(HDDA) or poly(TMPTMA) network. The
composite film reached a stable geometry when the tension in the VHB
network balanced with the compression in the poly(HDDA) or poly(TMPTMA)
network. This VHB network remained substantially strained compared to
its original geometry. The preservation of the prestrain of the VHB
network was measured using the distance between black dots marked on the
cured film. The resulting preserved prestrain versus weight percentage
of additive is plotted in Figure 3.

![](media/image3.wmf)

**Figure 3**. Preserved prestrain of VHB 4910 films after being released
from prestrain structure. The preserved prestrain is calculated based on
the original geometry of the VHB films before prestrain. Preserved
prestrains of 0% and 400% correspond to zero and full prestrain
preservation, respectively.

As expected, the prestrain preservation is strongly dependent on the
amount of additive used. The preserved prestrain increases rapidly above
a threshold concentration of additive . This may be explained by the
percolation theory \[14\]. When the amount of additive is increased to a
critical value, the polymer forms a crosslinked network of its own. This
network effectively supports the prestrain of the preexistent VHB
network. The threshold concentrations, the onset of the rapid increase
of preserved prestrain, are approximately 14% for HDDA and 8.4% for
TMPTMA. The lower threshold concentration of TMPTMA can be attributed to
its trifunctionality that facilitates the formation of its 3-dimensional
network in the acrylic elastomer.

![](media/image4.wmf){width="3.2319444444444443in"
height="2.0597222222222222in"}

![](media/image5.wmf){width="3.2284722222222224in"
height="2.040277777777778in"}

![](media/image6.wmf){width="3.2284722222222224in"
height="2.0305555555555554in"}

**Figure 4**. Scanning electron micrograph (SEM) images of the VHB films
for (a) 400% biaxial prestrained VHB 4910 without additive, (b) cured
VHB 4910 with 35.94% by weight of poly(HDDA) before releasing, (c) VHB
4910 with 23.38% by weight of poly(TMPTMA) before releasing.

Scanning electron micrographs displayed in Figure 4 showed sub-micron
size cracks. These cracks may be due to the broken additive network,
which is less compliant than the VHB network. Figure 4 (b) and (c)
compare the VHB-poly(HDDA) and the VHB-poly(TMPTMA) films containing the
same molar concentration of crosslinking acrylate or methacrylate
functional groups. In an HDDA molecule, the first acrylate group is used
for polymerization and the second acrylate is the crosslinkable group
that bonds polymer chains together. In TMP-TMA, each molecule contains 2
crosslinkable groups. In Figure 4, the VHB-poly(HDDA) film exhibits a
much rougher surface morphology and larger cracks than the
VHB-poly(TMPTMA). Consequently, the difunctional HDDA system is more
damaged than the trifunctional TMPTMA system containing the same amount
of crosslinkable vinyl double bonds. The IPN films with high
concentrations of additive appear turbid, indicating that the two
networks are chemically independent from each other. For the same
concentration of crosslinkable groups, the VHB-poly(HDDA) film is found
to be more turbid than the VHB-poly(TMPTMA) film. The turbidity can be
attributed to the phase separation of the additive network from the VHB
network. The degree of phase separation is observed to increase
proportionally with the turbidity of the IPN film.

*3.2. Infrared spectra of IPN films*

The curing of the additive and its network formation can be monitored by
recording the decrease of IR absorption peaks that are characteristic of
the carbon-carbon double bonds in the additive. Figures 5 and 6 show the
comparison of FTIR spectra from (a) a 400% biaxial prestrain VHB 4910
without additive, (b) VHB 4910 with additive before curing, and (c) VHB
4910 with additive after curing. Upon thermal curing, significant
changes are observed in the IR spectra of the IPN films. The absorption
peak at 810 cm^-1^ from alkene C-H stretching in the additive, decreases
to near disappearance after thermal curing. The other relatively intense
absorption peak at 1635 cm^-1^ (acrylic group. C=C) shown in Figures 5
and 6 (b) also drastically decreases after curing. From these results,
it is evident that the HDDA and TMPTMA has polymerized and crosslinked
(because they are di- and trifunctional), and that the interpenetrating
network of the cured additive has formed in the preexistent VHB network.

![](media/image7.wmf)

**Figure 5**. FTIR spectra of the VHB films for (a). 400% biaxially
prestrained VHB 4910, (b). (a) with 18.3 wt% HDDA, (c). (b) after
thermally curing.

![](media/image8.wmf)

**Figure 6**. FTIR spectra of the VHB films for (a). 400% biaxially
prestrained VHB 4910, (b). (a) with 9.7 wt% TMPTMA, (c). (b) after
thermally curing

*3.3. Mechanical properties of IPN composite films before release*

The stress-strain behavior of VHB-poly(HDDA) and VHB-poly(TMPTMA) was
measured, assuming little chemical crosslinking occured between the two
networks. Results show that the acrylic network remained intact during
the formation of the interpenetrating polymer network. Any variation in
the stress-strain behavior would occur due to the presence of an
additive network. Figure 7 shows the stress-strain response of the cured
composite films before the films were released to free-standing. The VHB
network is under 400% biaxial prestrain, whereas the additive network is
at zero prestrain. There is no experimental data to prove that
poly(HDDA) and poly(TMPTMA) aligns itself with the prestrained VHB
network.

![](media/image9.wmf)![](media/image10.wmf)

\(a\) VHB-poly(HDDA) (b) VHB-poly(TMPTMA)

**Figure 7**. Hydraulic pressure versus stretch ratio curves for
prestrained VHB 4910 (a) containing different concentrations of
poly(HDDA) and (b) poly (TMPTMA) before being released from prestrain
structure. The numbers in the insets are in weight percent of the
additive network in the IPN films.

As shown in Figure 7, the stress-strain response is nonlinear for VHB
without any additive, but it becomes increasingly linear with higher
concentrations of poly(HDDA) and poly(TMPTMA). The non-linearity may be
due to a combination of chain uncoiling and bond stretching \[15\]. At
higher stretching ratios, the chemical bonds are stretched very far
apart and the VHB network begins to break. With increasing
concentrations of additive, the films become stiffer, particularly at
high stretch ratios. At small stretch ratios (\<1.05), the
VHB-poly(HDDA) films in Figure 7 (a) are not much stiffer, or even
softer, than the pure VHB films. This is further discussed in the next
paragraph.

The theory for large deflection in circular thin plates developed by
Timoshenko \[16\] was used to evaluate the effect of additive on film
stiffness. At small stretch ratios, the Young's modulus E, of a clamped
circular plate under a uniform applied pressure P is given by

$E = \frac{3\text{Pa}^{4}(1 - \nu^{2})}{\text{16}h^{4}\left\lbrack \frac{\omega_{0}}{h} + \frac{0\text{.}\text{442}}{1 - \nu^{2}}\left( \frac{\omega_{0}}{h} \right)^{3} \right\rbrack}$
(1)

where *h*, *ω*~0~, a, and ν are film thickness, deflection at the film
center, and radius of plate and Poisson's ratio, respectively.

![](media/image11.wmf)

**Figure 8**. Calculated pseudo-Young's moduli of prestrained VHB 4910
films with or without additive. The VHB films were biaxially prestrained
by 400%, admixed with additive and cured. The moduli are calculated
based on a stretch ratio of 1.03.

Figure 8 shows the intriguing effects of additive content on the
pseudo-Young's modulus of the prestrained films before release. Note
that at zero deflection or at a stretch ratio of 1, the films have
built-in tension. The Young's moduli, which by definition are values at
minimal external tension, thus only reflect the stiffness of the films
at this condition, and do not correspond to the Young's moduli of the
films. Therefore, we call the calculated values "pseudo-Young's moduli".
The psuedo-Young's modulus of VHB 4910 under 400% biaxial prestrain is
18.42 MPa, which increases to 38.76 MPa when the poly(TMPTMA) content is
increased to 23.49%. This characteristic effect should result from
augmentative proportions of additive network. The additive monomers are
polymerized in the presence of the preexistent VHB matrix. As shown in
Figure 8, the psuedo-Young's modulus of VHB-poly(HDDA)first decreases
from 18.42 MPa to 13.3 MPa as the poly(TMPTMA) content is increased from
zero to 21.91%. The psuedo-Young's modulus then continues to increase
with increasing weight percentages of poly(HDDA). The poly(HDDA) content
corresponding to the modulus minimum falls near the high end of the
percolation transition concentration range shown in Figure 3. The
abnormal modulus decrease may be caused by the unreacted and locally
polymerized monomers that increase the total film thickness without
contributing much to the total elastic modulus \[17-19\].

*3.4. Mechanical behavior of IPN composite films after release*

The stress-strain behavior of the IPN composite films after they were
released to free-standing were measured. In the free-standing composite
films, the VHB network is under tension due to preserved prestrain.
Figure 3 shows the additive network under compression due to the
contraction of the films from a 400% prestrain to the preserved
prestrain.

![](media/image12.wmf) ![](media/image13.wmf)

\(a\) VHB-poly(HDDA) (b) VHB-poly(TMPTMA)

**Figure 9**. Hydraulic pressure versus stretch ratio curves for (a)
different concentrations of poly(HDDA) and (b) poly (TMPTMA) after being
released from prestrain structure. The numbers in the insets are weight
percent for poly(1,6-hexanediol diacrylate) in the composite films.

Figure 9 shows the biaxial stretch ratios by the hydraulic pressure of
the prestrain preserved, of free-standing composite films. The results
of the VHB-poly(HDDA) and VHB-poly(TMPTMA) are similar to those of the
unreleased films shown in Figure 7. The VHB-poly(HDDA) films in Figure 9
(a) seem to be considerably softer than the VHB-poly(TMPTMA) films shown
in Figure 9 (b) even with respect to equivalent preserved prestrain.
These results could be explained by comparing the bridge-length between
functional groups (or crosslinks) (see Figure 2). The flexible hexylene
segment in poly(HDDA) makes the network more pliable than that of the
TMPTMA, which has a much shorter bridge-length. At small stretch ratios,
the VHB-poly(HDDA) films appear much softer than the VHB films without
poly(HDDA). However, the IPN films are stiffer when film thickness is
considered.

![](media/image14.wmf)

**Figure 10**. Calculated Young's moduli of IPN films as a function of
weight percentage of additive. The VHB films were biaxially prestrained
by 400%, admixed with HDDA or TMP-TMA, cured, and released to
free-standing. The moduli were calculated based on =1.03.

As shown in Figure 10, the Timoshenko equation for free-standing
composite films was used to calculate the Young's moduli. The Young's
modulus of non-prestrained VHB 4910 is 181.53 kPa. Although the Young's
moduli of IPN composite films do not theoretically depend on the weight
percentage of poly(HDDA) in the range of 18.3-23.6%, experimental data
obtained show that the Young's moduli of VHB-poly(TMPTMA) films increase
rather rapidly with an increasing amount of poly(TMPTMA).

*3.5. Actuation of IPN films without external prestrain*

To accomplish a complete preservation of the original prestrain, a large
amount of additive is required. However, according to passive mechanical
analysis, the films with a larger amount of additive would require a
much higher energy to obtain large strains (stretch ratio greater than
1.2). The electrically-induced strains would thus be diminished. To
support this argument, the electrical strains of free-standing IPN
composite films with various amounts of additive were measured. The
results are displayed in Figure 11. The strains were measured with
increasing driving voltage until dielectric breakdown is reached in the
film.

![](media/image15.wmf)![](media/image16.wmf)

\(a\) VHB-poly(HDDA) (b) VHB-poly(TMPTMA)

**Figure 11**. Electric field induced strain (area expansion) for (a)
VHB-poly(HDDA) and (b) VHB-poly(TMPTMA) films containing various
concentrations of additive. The films have been released from the
prestrain structure and mounted, at minimal external stress, onto a
diaphragm structure for the strain measurement.

As expected, the IPN composite films show lower electrically-induced
strains with increasing amounts of additive. This is consistent with the
increasing film stiffness observed with increasing additive content at
large stretch ratios shown in Figure 9. For the VHB-poly(HDDA) shown in
Figure 11 (a), the IPN film with 18.33% poly(HDDA) show the highest
electrical strain of 233 %, with a breakdown electric field of 300 MV/m.
Its preserved prestrain was found to be 275%. The IPN formed film with
35.94% poly(HDDA) is too rigid to obtain high strain. The
VHB-poly(TMPTMA) film with 9.7 wt% additive exhibits a preserved strain
of 243.75%, with a maximum electrically-induced strain of 300%. This
result deserves further attention. As shown in Figure 9, although the
VHB-poly(TMPTMA) film is more rigid than the VHB-poly(HDDA) film with
18.33 wt% of additive, it exhibits a higher electrical strain. It also
has a dielectric breakdown strength of 420 MV/m, substantially higher
than that of the VHB-poly(HDDA) film. The improved breakdown field in
VHB-poly(TMPTMA) film can be attributed to the more uniform IPN film
morphology and the smaller, fewer cracks present in the additive network
(compare Figure 4 (b) and (c)). The rather rough surface and
submicron-size cracks in the VHB-poly(HDDA) shown in Figure 4 (b), are
responsible for the diminished dielectric strength \[20\]. The
electrically induced strain up to 300% in area is comparable to that of
the highly-prestrained, VHB 4910 films. Figure 12 shows the expansion of
a diaphragm actuator using free-standing IPN composite films.

![](media/image17.jpeg){width="2.667361111111111in"
height="1.5402777777777779in"}
![](media/image18.jpeg){width="2.6638888888888888in"
height="1.5402777777777779in"}

VHB-poly(HDDA)

![](media/image19.jpeg){width="2.6638888888888888in"
height="1.5402777777777779in"}
![](media/image20.jpeg){width="2.671527777777778in"
height="1.5381944444444444in"}

VHB-poly(TMPTMA)

**Figure 12.** Diaphragm actuators based on IPN composite films. The
films are under a constant, small tension maintained by the air pressure
in the diaphragm chamber. (a) VHB-poly(HDDA) containing 18.3 wt% HDDA at
0 MV/m, (b) at 300 MV/m, (c) VHB-poly(TMPTMA) film containing 9.7 wt%
TMP-TMA at 0 MV/m, (d) at 420 MV/m.

**4. Conclusion**

High performance interpenetrating polymer networks based on VHB acrylic
elastomers, poly(HDDA), and poly(TMPTMA) have been prepared for
electroelastomer artificial muscles. The interpenetrating polymer
networks are formed in the highly prestrained VHB acrylic elastomer
network to effectively support the prestrain of the VHB network. In the
free-standing composite films, the VHB network is under tension,
balanced by the poly(HDDA) or poly(TMPTMA) network that is in
compression. The preserved prestrain of the VHB network can be varied in
a wide range depending on the concentration of functional diacrylate or
trimethacrylate monomers. Small amounts of additive monomers are not
effective to lock the prestrain, while large amounts can substantially
increase the stiffness of the IPN films at large strains. The
VHB-poly(HDDA) composite film with 18.3 wt% poly(HDDA) preserves 275%
prestrain of the acrylic elastomer and exhibits an electrically-induced
strain up to 233% in area expansion, whereas the VHB-poly(TMPTMA)
composite film with 9.7 wt% poly(TMPTMA) retains 243.75% prestrain and
exhibits a 300% electrical strain. Studies are underway to use such
"prestrain-less" electroelastomer films to fabricate new dielectric
elastomer actuators without (or minimal) supporting structure/mass and
to deliver greater actuator performance.

**Acknowledgements**

This work was financially supported by the Office of Naval Research, ONR
Contract No. N00014-05-C-0005.

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