U.S. patent application number 15/881344 was filed with the patent office on 2018-06-07 for multifunctional graphene-silicone elastomer nanocomposite, method of making the same, and uses thereof.
This patent application is currently assigned to The Trustees of Princeton University. The applicant listed for this patent is The Trustees of Princeton University. Invention is credited to IIhan A. Aksay, Shuyang PAN, Robert K. Prud'homme.
Application Number | 20180155532 15/881344 |
Document ID | / |
Family ID | 44278013 |
Filed Date | 2018-06-07 |
United States Patent
Application |
20180155532 |
Kind Code |
A1 |
PAN; Shuyang ; et
al. |
June 7, 2018 |
MULTIFUNCTIONAL GRAPHENE-SILICONE ELASTOMER NANOCOMPOSITE, METHOD
OF MAKING THE SAME, AND USES THEREOF
Abstract
A nanocomposite composition having a silicone elastomer matrix
having therein a filler loading of greater than 0.05 wt %, based on
total nanocomposite weight, wherein the filler is functional
graphene sheets (FGS) having a surface area of from 300 m.sup.2/g
to 2630 m.sup.2/g; and a method for producing the nanocomposite and
uses thereof.
Inventors: |
PAN; Shuyang; (Southgate,
MI) ; Aksay; IIhan A.; (Princeton, NJ) ;
Prud'homme; Robert K.; (Lawrenceville, NJ) |
|
Applicant: |
Name |
City |
State |
Country |
Type |
The Trustees of Princeton University |
Princeton |
NJ |
US |
|
|
Assignee: |
The Trustees of Princeton
University
Princeton
NJ
|
Family ID: |
44278013 |
Appl. No.: |
15/881344 |
Filed: |
January 26, 2018 |
Related U.S. Patent Documents
|
|
|
|
|
|
Application
Number |
Filing Date |
Patent Number |
|
|
15225946 |
Aug 2, 2016 |
9908995 |
|
|
15881344 |
|
|
|
|
12945043 |
Nov 12, 2010 |
9441076 |
|
|
15225946 |
|
|
|
|
61260538 |
Nov 12, 2009 |
|
|
|
Current U.S.
Class: |
1/1 |
Current CPC
Class: |
C08J 2383/07 20130101;
C08G 77/20 20130101; C08K 9/00 20130101; C08L 83/04 20130101; C08G
77/38 20130101; B82Y 30/00 20130101; C08G 77/08 20130101; C08G
77/16 20130101; C08J 3/2053 20130101; C08G 77/18 20130101 |
International
Class: |
C08K 9/00 20060101
C08K009/00; C08L 83/04 20060101 C08L083/04; C08G 77/20 20060101
C08G077/20; C08G 77/18 20060101 C08G077/18; B82Y 30/00 20060101
B82Y030/00; C08G 77/16 20060101 C08G077/16; C08J 3/205 20060101
C08J003/205; C08G 77/38 20060101 C08G077/38; C08G 77/08 20060101
C08G077/08 |
Goverment Interests
[0002] This invention was made with government support under Grant
No. CMS-0609049 awarded by the National Science Foundation. The
Government has certain rights in the invention. The invention
described herein was also made in the performance of work under a
NASA contract (No. NNL08AF34P) and is subject to the provisions of
Public Law 96-517 (35 USC 202) in which the Contractor has elected
to retain title.
Claims
1. (canceled)
2. A method for production of a nanocomposite composition
comprising a silicone elastomer matrix and functionalized graphene
sheets having a surface area of from 300 m.sup.2/g to 2,630
m.sup.2/g, comprising: dispersing functional graphene sheets (FGS)
in a polar solvent to form an FGS suspension; combining the FGS
suspension with a vinyl terminated polysiloxane comprising a
diorganosiloxane unit and at least two silicon-bonded alkenyl
groups, the diorganosiloxane unit comprising a phenyl group;
removing the polar solvent; combining the resulting mixture with a
crosslinker and a hydrosilylation catalyst; curing the resulting
mixture to provide the nanocomposite; wherein the functional
graphene sheets have a loading of greater than 0.05 wt % based on
total nanocomposite weight; wherein the hydrosilylation catalyst is
present at a concentration of about 367 ppm to about 5600 ppm;
wherein the nanocomposite comprises a silicon hydride to vinyl
molar ratio of about 1.5 to about 2.1; and wherein the functional
graphene sheets are present within the nanocomposite in a
continuous three-dimensional connected network in a manner wherein
individual functional graphene sheets have nanometer scale
separation at contact point between individual functional graphene
sheets.
3. The method of claim 2, wherein the curing is performed at
elevated temperature for a period of time from 1 to 48 hours.
4. The method of claim 3, wherein the curing temperature is about
100.degree. C.
5. The method of claim 3, wherein the curing is performed for a
period of time from 5 to 30 hours, 20 to 25 hours, or approximately
24 hours.
6. The method of claim 5, further comprising molding the
nanocomposite in a manner to yield fewer functional graphene sheet
contacts point in a transverse direction of the nanocomposite
compared to a longitudinal direction of the nanocomposite.
7. The method of claim 2, wherein the functional graphene sheets
have a loading of from 0.5 to 3 wt %, based on total nanocomposite
weight.
8. The method of claim 2, wherein the silane cross-linker is a
member selected from the group consisting of
tetrakis(dialkylsiloxy)silanes and poly(hydromethyl siloxane)
crosslinkers.
9. The method of claim 2, wherein the silane cross-linker is a
tetrakis(dimethylsiloxy)silane.
10. The method of claim 2, wherein the vinyl-terminated
polysiloxane has a viscosity of from 100 to 300,000 mPas.
11. An article formed from a nanocomposite produced by the method
of claim 2.
12. The article of claim 11, wherein the article is formed by
casting.
13. The article of claim 11, wherein the article is formed by
molding.
14. The article of claim 11, wherein the molding process results in
the article comprising a longitudinal direction and a transverse
direction, the transverse direction comprising fewer functional
graphene sheet contact points compared to functional graphene sheet
contact points of the longitudinal direction.
15. The article of claim 11, wherein the article is a member
selected from the group consisting of coatings, adhesives,
sealants, flexible electrodes, actuators, pressure sensors, printed
circuits, and electromagnetic interference shielding materials.
16. A method for production of a nanocomposite composition
comprising a silicone elastomer matrix and functionalized graphene
sheets having a surface area of from 300 m.sup.2/g to 2,630
m.sup.2/g, comprising: dispersing functional graphene sheets (FGS)
in a polar solvent to form an FGS suspension; combining the FGS
suspension with a vinyl terminated polysiloxane having a viscosity
of from 100 to 300,000 mPas, the vinyl terminated polysiloxane
comprising a diorganosiloxane unit and at least two silicon-bonded
alkenyl groups, the diorganosiloxane unit comprising a phenyl
group; removing the polar solvent; combining the resulting mixture
with a crosslinker and a hydrosilylation catalysts, wherein the
silane cross-linker is a member selected from the group consisting
of tetrakis(dialkylsiloxy)silanes and poly(hydromethyl siloxane)
crosslinkers, and wherein the hydrosilylation catalyst is a member
selected from the group consisting of chloroplantinic acid,
elementary platinum, solid platinum supported on a carrier;
platinum-vinylsiloxane complexes; platinum-phosphine complexes;
platinum-phosphite complexes; Pt (acac).sub.2, wherein (acac)
represents acetylacetonate group; platinum-hydrocarbon conjugates;
platinum alcoholates; RhCl(PPh.sub.3).sub.3; RhCl.sub.3;
Rh/Al.sub.2O.sub.3; RuCl.sub.3; IrCl.sub.3; FeCl.sub.3; AlCl.sub.3;
PdCl.sub.2.2H.sub.2O; NiCl.sub.2; and TiCl.sub.4; and curing the
resulting mixture to provide the nanocomposite, wherein the curing
is performed at elevated temperature for a period of time from 1 to
48 hours; wherein the functional graphene sheets have a loading of
greater than 0.05 wt % based on total nanocomposite weight; wherein
the hydrosilylation catalyst is present at a concentration of about
367 ppm to about 5600 ppm; wherein the nanocomposite comprises a
silicon hydride to vinyl molar ratio of about 1.5 to about 2.1; and
wherein the functional graphene sheets are present within the
nanocomposite in a continuous three-dimensional connected network
in a manner wherein individual functional graphene sheets have
nanometer scale separation at contact points between individual
functional graphene sheets.
17. The method of claim 16, further comprising molding the
nanocomposite in a manner to yield fewer functional graphene sheet
contacts point in a transverse direction of the nanocomposite
compared to a longitudinal direction of the nanocomposite.
18. The method of claim 16, wherein the vinyl terminated
polysiloxane comprises a siloxane unit selected from selected from
the group consisting of R.sub.3SiO.sub.1/2, RSiO.sub.3/2, and
SiO.sub.4/2, where R represents a substituted monovalent
hydrocarbon group or a unsubstituted monovalent hydrocarbon
group.
19. The method of claim 2, wherein the vinyl terminated
polysiloxane comprises a siloxane unit selected from selected from
the group consisting of R.sub.3SiO.sub.1/2, RSiO.sub.3/2, and
SiO.sub.4/2, where R represents a substituted monovalent
hydrocarbon group or an unsubstituted monovalent hydrocarbon group.
Description
REFERENCE TO RELATED APPLICATIONS
[0001] The present application is a Continuation of U.S.
application Ser. No. 15/225,946, filed Aug. 2, 2016, now allowed,
which is a Continuation of U.S. application Ser. No. 12/945,043,
filed Nov. 12, 2010, now U.S. Pat. No. 9,441,076, and also claims
priority on U.S. Provisional Application Ser. No. 61/260,538, filed
Nov. 12, 2009, the entire contents of each of which are hereby
incorporated by reference.
BACKGROUND OF THE INVENTION
Field of the Invention
[0003] The present invention relates to nanocomposites having a
matrix of silicone elastomer with multifunctional graphene sheets
as filler, methods of making the same and their use.
Description of the Related Art
[0004] The effect of filler dispersion on the mechanical properties
of the resulting composite has been studied for decades but a
consensus is yet to be reached. Many have suggested that maximizing
filler dispersion is crucial in achieving good mechanical
properties. For example, for carbon nanotubes (CNT), Ajayan et al.
suggested that load transfer can be limited when the nanotubes are
slipping within the bundles..sup.1 The bundles need to be broken
into individual dispersed tube segments to obtain effective modulus
increase and strengthening. Schandler et al. have proposed that
infiltrating the polymer into the interstices of the nanotube
bundles can create effective load transferring and therefore
mechanical reinforcement..sup.2 Similarly for inorganic fillers,
Lebaron et al. have suggested that the complete dispersion of clay
optimized the number of reinforcing elements for carrying an
applied load and deflecting cracks, allowing for tensile property
improvements..sup.3 .sup.1Ajayan, P. M.; Schadler, L. S.;
Giannaris, C.; Rubio, A. Advanced Materials 2000, 12, (10),
750-.sup.2Schadler, L. S.; Giannaris, S. C.; Ajayan, P. M. Applied
Physics Letters 1998, 73, (26), 3842-3844.sup.3LeBaron, P. C.;
Wang, Z.; Pinnavaia, T. J. Applied Clay Science 1999, 15, (1-2),
11-29
[0005] Large clusters of particles can act as flaws to initiate
premature termination of stretching..sup.4 On the other hand, it
has long been suggested in the automotive tire industry that
aggregated fillers are more effective than primary particles in
enhancing the modulus and tensile strength of the elastomer..sup.5
At large strains, the deformation and irreversible breakdown of
aggregates absorb energy, allowing the composite to tolerate higher
amounts of stress. However, a rigorous understanding of the effect
of breaking up initial filler agglomerates on the mechanical
properties that incorporates the two aforementioned contrasting
views, is lacking. .sup.4Wilbrink, M. W. L.; Argon, A. S.; Cohen,
R. E.; Weinberg, M. Polymer 2001, 42, (26),
10155-10180.sup.5Poovarodom, S.; Hosseinpour, D.; Berg, J. C.
Industrial & Engineering Chemistry Research 2008, 47, (8),
2623-2629
[0006] In achieving the maximum effect with the minimum filler
loading, it is important to understand the correlation between the
spatial distribution of dispersed fillers and the macroscopic
mechanical properties of the composite..sup.6,7 Some understanding
of the structure-property relationship has been developed
previously by others. A larger agglomeration of silica renders a
better improvement in the Young's modulus of the matrix..sup.8 It
has been shown by Akcora et al. that self-assembled nanoparticle
sheet yielded a solid-like rheological behavior in polystyrene
whereas well-dispersed short particle strings did not..sup.9
However, the effect of filler assembly on the tensile properties of
the composites is not yet well-understood. .sup.6Vaia, R. A.;
Maguire, J. F. Chemistry of Materials 2007, 19, (11),
2736-2751.sup.7Balazs, A. C.; Emrick, T.; Russell, T. P. Science
2006, 314, (5802), 1107-1110.sup.8Oberdisse, J. Soft Matter 2006,
2, (1), 29-36.sup.9Akcora, P.; Liu, H.; Kumar, S. K.; Moll, J.; Li,
Y.; Benicewicz, B. C.; Schadler, L. S.; Acehan, D.;
Panagiotopoulos, A. Z.; Pryamitsyn, V.; Ganesan, V.; Ilaysky, J.;
Thiyagarajan, P.; Colby, R. H.; Douglas, J. F. Nature Materials
2009, 8, (4), 354-U121
[0007] Another fundamental issue that has drawn much attention is
the origin of the reinforcements of tensile properties in
composites. Simultaneous improvements in modulus, strength and
elongation at break with the incorporation of fillers have been
observed in poly(methylmethacrylate),.sup.10 epoxy,.sup.11
styrene-butadiene rubber,.sup.12 polyimide,.sup.13 and silicone
rubber..sup.14,15,16,17,18 While the modulus and strength increase
with the filler concentration, the elongation at break in some
cases increases initially and then decreases above a critical
filler concentration..sup.11,13,16,17 .sup.10Sui, X. M.; Wagner, H.
D. Nano Letters 2009, 9, (4), 1423-1426.sup.11Tseng, C. H.; Wang,
C. C.; Chen, C. Y. Chemistry of Materials 2007, 19, (2),
308-315.sup.12Bokobza, L.; Rahmani, M.; Belin, C.; Bruneel, J. L.;
El Bounia, N. E. Journal Of Polymer Science Part B-Polymer Physics
2008, 46, (18), 1939-1951.sup.13An, L.; Pan, Y. Z.; Shen, X. W.;
Lu, H. B.; Yang, Y. L. Journal of Materials Chemistry 2008, 18,
(41), 4928-4941.sup.14Aranguren, M. I.; Mora, E.; Macosko, C. W.;
Saam, J. Rubber Chemistry And Technology 1994, 67, (5),
820-833.sup.15Yuan, Q. W.; Mark, J. E. Macromolecular Chemistry And
Physics 1999, 200, (1), 206-220.sup.16Osman, M. A.; Atallah, A.;
Muller, M.; Suter, U. W. Polymer 2001, 42, (15),
6545-6556.sup.17Bokobza, L.; Rahmani, M. Kgk--Kautschuk Gummi
Kunststoffe 2009, 62, (3), 112-117.sup.18LeBaron, P. C.; Pinnavaia,
T. J. Chemistry Of Materials 2001, 13, (10), 3760-3765
[0008] The increase in modulus is attributed to load transferring
to the stiffer filler material..sup.19,20 Some understanding has
been achieved in the tensile strength and elongation at break
increase. Sui et al. demonstrated using transmission electron
microscopy (TEM) the mechanism responsible for the significant
elongation at break increase in electrospun CNT-poly(methyl
methacrylate) (PMMA) fibers..sup.10 In pure PMMA fiber, sparse and
unstable necking was observed along the fiber under tension,
followed by failure of the fiber. When 1.5 wt. % single wall carbon
nanotubes (SWCNT) were added, multiple necking was initiated but
arrested by SWCNT ropes. Further stretching led to bridging by
SWCNT ropes, which caused a dilation effect in the fiber and an
increase in the elongation at break. The inelastic strain and
energy dissipation introduced by the necking and bridging was
proposed to explain the tensile strength increase of the
nanocomposite. Only one CNT concentration was used. In the same
study, millimeter-sized pure and CNT filled PMMA films were studied
and improvement in the elongation at break was also observed,
although to a lesser extent compared to the electrospun fibers. The
improvement in the films was not addressed in the study.
.sup.19Hashin, Z.; Shtrikman, S. Journal Of The Mechanics And
Physics Of Solids 1963, 11, (2), 127-140.sup.20Nielsen, L. E.
Journal Of Applied Physics 1970, 41, (11), 4626-&
[0009] Load transferring to CNT has been proposed to explain the
strength and elongation at break increase in epoxy..sup.11 When an
amphiphilic block copolymer was incorporated into epoxy, elongation
at break increase was observed..sup.21 The underlying mechanisms
were investigated with optical microscopy and TEM. It was found
that a 15 nm size spherical block copolymer micelle could cavitate
to induce matrix shear banding. It was suggested that the dilation
effect and shear banding introduced by the cavitation led to the
observed increase in the elongation at break. .sup.21Liu, J.; Sue,
H. J.; Thompson, Z. J.; Bates, F. S.; Dettloff, M.; Jacob, G.;
Verghese, N.; Pham, H. Macromolecules 2008, 41, (20), 7616-7624
[0010] When rod-like attapulgite was incorporated into polyimide,
simultaneous improvements in modulus, strength and elongation at
break were observed..sup.13 The enhancement of the interfacial
stress transfer and the resistance to crack propagation induced by
attapulgite was proposed to explain the mechanical
reinforcement.
[0011] Filler agglomerates acting as defects have been proposed to
explain the reversal in the elongation at break..sup.11,22
Incorporation of free volume with the filler has also been
suggested to be causing the reversal effect..sup.13 The addition of
filler increased the free volume or defects in nanocomposites and
the resistance to crack propagation during deformation. Below the
critical concentration, the latter effect dominated and elongation
at break increased. Above the threshold, the increase in the number
of defects dominated and the elongation at break started to
decrease. The reversal effect was also observed with the
incorporation of polystyrene-modified cadmium selenide
nanoparticles to polystyrene (PS)..sup.23 It was proposed that two
competing effects determine the elongation at break of the
composite. Nanoparticles entrapped within the mature craze during
craze widening disrupt the formation of cross-tie fibrils by
increasing the mobility of polymer segments at the craze-bulk
interface. Less cross-tie fibrils reduced the premature rupture of
the craze fibrils and increased the failure strain. On the other
hand, entrapped nanoparticles also reduced the extensibility of the
craze fibrils or the dilation effect of the craze. So the two
competing effects led to a maximum in elongation at break of the
composite as a function of nanoparticle concentrations.
.sup.22Gorga, R. E.; Cohen, R. E. Journal of Polymer Science Part
B-Polymer Physics 2004, 42, (14), 2690-2702.sup.23Lee, J. Y.;
Zhang, Q. L.; Wang, J. Y.; Emrick, T.; Crosby, A. J. Macromolecules
2007, 40, (17), 6406-6412
[0012] The simultaneous improvements are not limited to polymeric
matrices. The incorporation of polymeric fibers increased the
strength and elongation at break of the newly engineered building
material called engineered cementitious composites (ECC)..sup.24
ECCs have been designed to distribute many cracks of small width
throughout the composite rather than only a few large cracks seen
in traditional concrete failure. Such a distributed deformation is
responsible for the observed mechanical reinforcement. Similar
mechanisms have been shown to cause the elongation at break
increase in biological composites such as nacre..sup.25 .sup.24Li,
V. C.; Wang, S. X.; Wu, C. Aci Materials Journal 2001, 98, (6),
483-492.sup.25Wang, R. Z.; Suo, Z.; Evans, A. G.; Yao, N.; Aksay,
I. A. Journal Of Materials Research 2001, 16, (9), 2485-2493
[0013] Despite the aforementioned efforts, some fundamental issues
governing the tensile properties improvements have not been
completely understood. For example, it is not known how the filler
agglomeration and filler concentration influence the interaction
between fillers and tears or cracks, nor how filler length scale
influences the interaction. Further, it is not known how the
interaction is related to the reversal effect or how the local
deformation is directly correlated with the macroscopic tensile
properties in bulk composites. Lastly, it is not known how
mechanical load is being transferred to the filler. These are all
critical questions that need to be addressed in order to gain a
complete understanding of the reinforcement.
[0014] One potential filler that has been suggested is functional
graphene sheets (FGS). FGS is an atomically thin layer of graphite
hundreds of nanometers in the lateral dimension and decorated with
carboxyls at the edges and hydroxyls and epoxides on the planes.
Our group invented a method to produce functionalized graphene
sheet (FGS) on a large scale; see U.S. Patent Application
Publication 2007/0092432, filed Oct. 14, 2005 and published Apr.
26, 2007 (the entire contents of which are hereby incorporated by
reference; hereafter "the '432 application"). It has a wrinkled
geometry with an average aspect ratio of 500 and a surface area
from 300 m.sup.2/g to 2630 m.sup.2/g, typically up to 1800
m.sup.2/g..sup.26,27 It is preferably produced through thermal
exfoliation and reduction of oxidized natural graphite. The '432
application further discloses these FGS products. Stankovich et al.
developed an alternative method to produce graphene..sup.28
Graphene oxide was first obtained by oxidation of natural graphite
and sonication of graphite oxide. Chemical reduction of graphene
oxide yielded graphene with good electrical conductivity. In a
recent study, significant increases in glass transition
temperature, Young's modulus, tensile strength and electrical
conductivity was observed in when 1 weight % of FGS was
incorporated into poly(methyl methacrylate) and
poly(acrylonitrile)..sup.29 An enhancement in the modulus and
electrical conductivity as well as a reduction in the coefficient
of thermal expansion and gas permeability was observed when FGS was
added to poly(ethylene-2,6-naphthalate) and
poly(carbonate)..sup.30,31 When reduced graphene oxide was
incorporated into polystyrene, a low electrical percolation of 0.1
vol. % and good conductivities were obtained..sup.28
.sup.26Schniepp, H. C.; Kudin, K. N.; Li, J. L.; Prud'homme, R. K.;
Car, R.; Saville, D. A.; Aksay, I. A. Acs Nano 2008, 2, (12),
2577-2584.sup.27McAllister, M. J.; Li, J. L.; Adamson, D. H.;
Schniepp, H. C.; Abdala, A. A.; Liu, J.; Herrera-Alonso, M.;
Milius, D. L.; Car, R.; Prud'homme, R. K.; Aksay, I. A. Chemistry
Of Materials 2007, 19, (18), 4396-4404.sup.28Stankovich, S.; Dikin,
D. A.; Dommett, G. H. B.; Kohlhaas, K. M.; Zimney, E. J.; Stach, E.
A.; Piner, R. D.; Nguyen, S. T.; Ruoff, R. S. Nature 2006, 442,
(7100), 282-286.sup.29Ramanathan, T.; Abdala, A. A.; Stankovich,
S.; Dikin, D. A.; Herrera-Alonso, M.; Piner, R. D.; Adamson, D. H.;
Schniepp, H. C.; Chen, X.; Ruoff, R. S.; Nguyen, S. T.; Aksay, I.
A.; Prud'homme, R. K.; Brinson, L. C. Nature Nanotechnology 2008,
3, (6), 327-331.sup.30Kim, H.; Macosko, C. W. Macromolecules 2008,
41, (9), 3317-3327.sup.31Kim, H.; Macosko, C. W. Polymer 2009, 50,
(15), 3797-3809
[0015] U.S. patent application Ser. No. 11/543,872, filed Oct. 6,
2006 (the entire contents of which are hereby incorporated by
reference), discloses the use of the FGS of the '432 application in
the production of various nanocomposite rubbers.
[0016] SE has attracted both scientific and commercial interest for
its thermal stability over a wide range of temperatures (-50 to
over 200.degree. C.), retention of elastomeric properties at low
temperatures due to a low glass transition temperature of
-125.degree. C., its chemical and weathering
resistance..sup.32,33,34 SE is typically made by end-linking
poly(dimethyl siloxane) (PDMS) and therefore its molecular weight
between crosslinks is well-characterized. Due to its relatively
inferior tensile strength in the unfilled state (typically less
than 1 MPa, compared to more than 10 MPa of natural rubber), silica
is generally used to render SE applicable in commercial
applications..sup.34,34 Other fillers including
silica,.sup.14,15,35 clays,.sup.16,16,36 carbon nanotubes
(CNT),.sup.17,37 graphite nanosheet,.sup.38 glass fiber,.sup.39 and
in-situ precipitated alumina,.sup.40 have also been studied as
alternative fillers for SE. .sup.32Mark, J. E. Accounts Of Chemical
Research 2004, 37, (12), 946-953.sup.33Noll, W., Chemistry and
Technology of Silicones. Academic Press, Inc.: New York,
1978.sup.34Butts, M.; et. al. In Kirk-Othmer Encyclopedia of
Chemical Technology-Silicones. Wiley Interscience: New York,
2004.sup.35Mark, J. E.; Jiang, C. Y.; Tang, M. Y. Macromolecules
1984, 17, (12), 2613-2616.sup.36Osman, M. A.; Atallah, A.; Kahr,
G.; Suter, U. W. Journal of Applied Polymer Science 2002, 83, (10),
2175-2183.sup.37Frogley, M. D.; Ravich, D.; Wagner, H. D.
Composites Science And Technology 2003, 63, (11),
1647-1654.sup.38Chen, L.; Lu, L.; Wu, D. J.; Chen, G. H. Polymer
Composites 2007, 28, (4), 493-498.sup.39Park, E. S. Journal of
Applied Polymer Science 2007, 105, (2), 460-468.sup.40Mark, J. E.;
Wang, S. B. Polymer Bulletin 1988, 20, (5), 443-448
SUMMARY OF THE INVENTION
[0017] Accordingly, one object of the present invention is to
provide a nanocomposite based on silicone elastomers that has one
or more of higher modulus, strength, failure strain, electrical
conductivity and lower gas permeability than the unfilled silicone
elastomer.
[0018] A further object of the present invention is to provide a
method for producing such a nanocomposite.
[0019] A further object of the present invention is to provide
articles made from the nanocomposite, including, but not limited to
electrically conductive and low-permeability coating, adhesive and
sealants, as well as flexible electrodes, actuators, pressure
sensor, printed circuits and electromagnetic interference shielding
material.
[0020] These and other objects of the present invention, either
alone or in combinations thereof, have been satisfied by the
discovery of a nanocomposite composition comprising:
[0021] a silicone elastomer matrix having therein a filler loading
of greater than 0.05 wt %, based on total nanocomposite weight;
[0022] wherein the filler is functional graphene sheets (FGS)
having a surface area of from 300 m.sup.2/g to 2630 m.sup.2/g;
[0023] a method for producing the nanocomposite composition and its
use in a variety of end products.
BRIEF DESCRIPTION OF THE DRAWINGS
[0024] The patent or application file contains at least one drawing
executed in color. Copies of this patent or patent application
publication with color drawing(s) will be provided by the Office
upon request and payment of the necessary fee.
[0025] A more complete appreciation of the invention and many of
the attendant advantages thereof will be readily obtained as the
same becomes better understood by reference to the following
detailed description when considered in connection with the
accompanying drawings, wherein:
[0026] FIGS. 1A-1F show SEM images of cryo-fractured unfilled and
FGS-filled SE surfaces.
[0027] FIG. 2 provides a graphical representation of the effect of
filler concentration on the electrical conductivity of various FGS
and graphene filled nanocomposites.
[0028] FIG. 3 provides a graphical representation of stress-strain
curves of FGS-SE nanocomposites at different FGS
concentrations.
[0029] FIGS. 4A-4F provide photographs of the tearing of unfilled
SE (A) and 0.5 wt. % FGS-SE (B)-(D) and SEM images of
tensile-fractured surfaces of unfilled (E) and 0.5 wt % FGS-filled
SE (F).
[0030] FIG. 5 provides a graphical representation of hysteresis
characterization of unfilled and 0.5 wt % FGS-SE nanocomposite.
[0031] FIGS. 6A-6D provide images of the deformed lattice in
unfilled and FGS-SE nanocomposites.
[0032] FIGS. 7A-7C provide graphical representations of (7A).
Simulated normalized stress-strain curves of unfilled and
FGS-filled SE; (7B). Fraction of the matrix torn versus FGS vol. %
at three different strains; and (7C). Average strain of tears
versus FGS vol. %.
[0033] FIGS. 8A and 8B provide graphical representations of a
comparison of FGS with other fillers in the modulus of the
composite and the improvement in the
modulus..sup.14,16,17,18,36,38
[0034] FIGS. 9A and 9B provide graphical representations of a
comparison of FGS with other fillers in the tensile strength of the
composite and the improvement in tensile strength rendered by the
filler..sup.14-18,36,39,40
[0035] FIG. 10 provides a graphical representation of the effect of
catalyst concentration on the modulus of SE at r=1.5 for all
samples.
[0036] FIG. 11 provides a graphical representation of the effect of
silicon hydride to vinyl ratio on the modulus of FGS-SE
nanocomposite.
DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS
[0037] The present invention involves the addition of
functionalized graphene sheets (FGS) to silicone elastomer. Within
the context of the present invention the term "silicone elastomer"
is used to refer to any of a variety of elastomeric modified
silicone polymers as distinct from unmodified polydimethylsiloxane
(PDMS). The invention has higher modulus, strength, failure strain,
electrical conductivity and lower gas permeability than the
unfilled silicone elastomer. The current invention introduces new
applications for silicone elastomer, such as electrically
conductive and low-permeability coating, adhesive and sealants, as
well as flexible electrodes, actuators, pressure sensor, printed
circuits and electromagnetic interference shielding material.
[0038] The FGS-silicone elastomer nanocomposite of a most preferred
embodiment of the present invention simultaneously has superior
mechanical, electrical and barrier properties compared to unfilled
silicone elastomers. Furthermore, the current invention provides a
method to detect internal damage within the material through
conductance measurements. So when used in the industry, it
possesses health-monitoring capability which can be crucial for
applications using the nanocomposite. The product also has lower
density than commercially available silicone and therefore can
reduce the energy cost associated with transporting and using the
product.
[0039] FGS can be produced via a process that has been described in
published articles (H. C. Schniepp, J.-L. Li, M. J. McAllister, et
al., J. Phys. Chem. B 110, 8535-39, 2006; M. J. McAllister, J.-L.
Li, D. H. Adamson, H. C. Schniepp, et al., Chem. Materials 19,
4396-4404, 2007) (the entire contents of each of which are hereby
incorporated by reference) and the '432 application noted
above.
[0040] In a preferred embodiment of the present invention method
for forming the nanocomposite, FGS is dispersed in a polar solvent,
such as tetrahydrofuran, and probe-sonicated. Then the suspension
is combined with a vinyl terminated polysiloxane (preferably a
vinyl terminated polydimethylsiloxane) and the polar solvent is
completely evaporated off. An appropriate crosslinker and
hydrosilylation catalyst (preferably a platinum complex catalyst)
are combined with the resulting mixture and the mixture is cured at
elevated temperature, preferably about 100.degree. C. for a period
of time from 1 to 48 hours, preferably from 5 to 30 hours, more
preferably from 20-25 hours, most preferably approximately 24
hours.
[0041] The crosslinking reaction for the silicone elastomer
involves the reaction between the crosslinker, and the vinyl
terminating groups on the vinyl-terminated polysiloxane in the
presence of a hydrosilylation catalyst. Suitable crosslinking
agents include any conventional crosslinking agent, such as those
disclosed in "The Basics of Silicon Chemistry" (Dow Corning
Publication); W. Noll, Chemistry and technology of Silicones,
Academic Press, New York (1968); T. C. Kendrick, B. Parbhoo, J. W.
White, "Siloxane Polymers and Copolymers," in The Chemistry of
Organic Silicon Compounds Pt. 2 (edited by S. Patai and Z.
Rappoport), 21, p. 1289-1361, John Wiley, Chichester (1989); and S.
J. Clarson, J. A. Semlyen, Siloxane Polymers, Prentice Hall, New
Jersey (1993), the contents of each of which are hereby
incorporated by reference. Preferably, the crosslinking agent is
selected from tetrakis (dimethyl siloxy) silanes, or
poly(hydromethylsiloxane) crosslinkers. The resulting mechanical
properties, electrical properties and gas permeability of the
FGS-silicone elastomer nanocomposite showed increased modulus,
elongation at break, tensile strength and electrical conductivity
and decreased gas permeability, as compared to the same silicone
elastomer without the FGS filler.
[0042] The hydrosilylation catalyst is not particularly restricted,
and can be any conventional hydrosilylation catalyst. Specific
examples include, but are not limited to, chloroplatinic acid,
elementary platinum, solid platinum supported on a carrier such as
alumina, silica or carbon black; platinum-vinylsiloxane complexes
{e.g. Pt.sub.n(ViMe.sub.2SiOSiMe.sub.2Vi).sub.n,
Pt[(MeViSiO).sub.4].sub.m}; platinum-phosphine complexes {e.g.
Pt(PPh.sub.3).sub.4, Pt(PBU.sub.3).sub.4}; platinum-phosphite
complexes {e.g. Pt[P(OPh).sub.3].sub.4, Pt[P(OBu).sub.3].sub.4} (in
the above formulas, Me stands for methyl, Bu for butyl, Vi for
vinyl, Ph for phenyl, and n and m each represents an integer); Pt
(acac).sub.2; and platinum-hydrocarbon conjugates described by
Ashby et al. in U.S. Pat. Nos. 3,159,601 and 3,159,662 as well as
platinum alcoholates described by Lamoreaux et al. in U.S. Pat. No.
3,220,972, the contents of each of which are hereby incorporated by
reference.
[0043] As examples of the catalyst, other than platinum compounds,
there may be mentioned RhCl(PPh.sub.3).sub.3, RhCl.sub.3,
Rh/Al.sub.2O.sub.3, RuCl.sub.3, IrCl.sub.3, FeCl.sub.3, AlCl.sub.3,
PdCl.sub.2.2H.sub.2O, NiCl.sub.2, TiCl.sub.4, etc. These catalysts
may be used singly or two or more of them may be used in
combination. From the viewpoint of catalytic activity,
chloroplatinic acid, platinum-olefin complexes,
platinum-vinylsiloxane complexes, Pt(acac).sub.2 and the like are
preferred, with platinum-cyclovinylmethylsiloxane complex being
most preferred. The amount of the catalyst is not particularly
restricted but the catalyst is preferably used in an amount within
the range of 10.sup.-1 to 10.sup.-8 moles, more preferably
10.sup.-2 to 10.sup.-6 moles, per mole of the alkenyl group in the
vinyl-terminated polysiloxane. Hydrosilylation catalysts are
generally expensive and corrosive and, in some instances, they
induce generation of hydrogen gas in large amount to thereby cause
foaming of cured products. Therefore, it is recommended that their
use in an amount of more than 10.sup.-1 moles be avoided.
[0044] Within the context of the present invention, the term
"vinyl-terminated polysiloxane" is used to represent a component of
the present invention siloxane elastomer that contains at least one
diorganosiloxane unit and has at least two silicon-bonded alkenyl
groups in each molecule. The alkenyl group can be exemplified by
vinyl, allyl, butenyl, pentenyl, hexenyl, and heptenyl and is
preferably vinyl. The non-alkenyl Si-bonded organic groups are
exemplified by alkyl groups such as methyl, ethyl, propyl, butyl,
pentyl, and hexyl; aryl groups such as phenyl, tolyl, and xylyl;
and halogenated alkyl groups such as 3-chloropropyl and
3,3,3-trifluoropropyl, and is preferably methyl and/or phenyl. The
molecular structure of the vinyl-terminated polysiloxane is not
critical as long as it contains at least one diorganosiloxane unit,
i.e., siloxane unit with a general formula R.sub.2SiO.sub.2/2. As
other siloxane units, the vinyl-terminated polysiloxane may contain
small amounts of siloxane unit with a general formula
R.sub.3SiO.sub.1/2, siloxane unit with a general formula
RSiO.sub.3/2, and siloxane unit with a general formula SiO.sub.4/2.
R in the preceding formulas represents a substituted or
unsubstituted monovalent hydrocarbon group and can be exemplified
by the alkyl, alkenyl, aryl, and halogenated alkyl referenced
above. The molecular structure of the vinyl-terminated polysiloxane
can be exemplified by straight chain, branched chain, partially
branched straight chain, and dendritic, wherein straight chain,
branched chain, and partially branched straight chain are
preferred. The viscosity of the vinyl-terminated polysiloxane at
25.degree. C. is not critical, but is preferably 100 to 1,000,000
mPas and more preferably is 100 to 500,000 mPas, most preferably
from 100 to 300,000 mPas. The weight average molecular weight of
the vinyl-terminated polysiloxane is also not particularly critical
and will depend on the end use desired for the finished FGS-SE
composition. Preferably the weight average molecular weight of the
vinyl-terminated polysiloxane is in a range from 5000 to 2,000,000,
more preferably from 5000 to 50,000, most preferably from 8000 to
12,000.
[0045] The vinyl-terminated polysiloxane is preferably a member
selected from dimethylvinylsiloxy-endblocked dimethylpolysiloxanes;
dimethylvinylsiloxy-endblocked dimethylsiloxane-methylvinylsiloxane
copolymers; trimethylsiloxy-endblocked
dimethylsiloxane-methylvinylsiloxane copolymers; branched-chain
dimethylpolysiloxane with molecular chain ends terminated by
dimethylvinylsiloxy and trimethylsiloxy; trimethylsiloxy-endblocked
branched-chain dimethylsiloxane-methylvinylsiloxane copolymers; the
organopolysiloxanes afforded by replacing all or part of the methyl
in the preceding organopolysiloxanes with alkyl such as ethyl or
propyl, aryl such as phenyl or tolyl, or halogenated alkyl such as
3,3,3-trifluoropropyl; the organopolysiloxanes afforded by
replacing all or part of the vinyl in the preceding
organopolysiloxanes with alkenyl such as allyl or propenyl; and
mixtures of two or more of the preceding organopolysiloxanes.
[0046] For convenience, the vinyl-terminated polysiloxane will be
discussed with reference to a vinyl-terminated
poly(dimethylsiloxane). However, this is not intended to be
limiting of the present invention, but merely used in an exemplary
manner for convenience.
[0047] The present invention nanocomposite properties provide the
ability to monitor the structural health of products formed from
the nanocomposite by measuring conductance properties to detect
internal damage in the resulting product.
[0048] In the product of the present invention, tensile properties
improvements are preferably achieved when FGS is percolated in SE.
Within the context of the present invention, the term "percolated"
is intended to indicate that a continuous path is established in
three dimensions through the FGS by formation of a connected FGS
network with nanometer scale separation at the contact point
between individual sheets. Normally, the FGS sheets are
statistically in contact. Indications of percolation are the onset
of the transition from non-electrically conducting to electrically
conducting, or the state in which the storage and loss moduli
measured as a function of frequency (G'(.omega.) and G'(.omega.),
respectively) scale as
G'(.omega.).about.G'(.omega.).about..omega..sup.n. These
characteristics are meant to be indicative of percolation and are
not intended as limiting the present invention. Agglomeration of
FGS can be observed using SEM. The agglomeration facilitates
electrical percolation and therefore tensile properties
improvements. Although the present inventors do not wish to be
bound by any particular mechanistic explanation for the improvement
in properties in the present invention, it is believed that the
increase of tensile strength can be attributed to load transfer to
FGS. The increase of elongation at break is believed to be due to
the dilation effect of tearing and distributed deformation
introduced by the percolated FGS network. The reversal in the
elongation at break is observed and is believed to be due to the
competing effects of the degree of tear opening and the number of
tears with increasing FGS concentration. Multifunctional
reinforcement of SE by FGS is also demonstrated.
[0049] Experimental Section
[0050] 2.1. Materials.
[0051] Vinyl-terminated PDMS with an average molecular weight of
9400, tetrakis(dimethylsiloxy)silane and
platinum-cyclovinylmethylsiloxane complex were obtained from
Gelest, Inc. Tetrahydrofuran (THF) was purchased from Sigma
Aldrich. FGS was produced using a thermal exfoliation method
previously reported using graphite oxide (GO) supplied by Vorbeck
Materials..sup.27,27 The carbon to oxygen ratio of the FGS was
determined to be 15 to 1 using modified classical Pregl and Dumas
method by Atlantic Microlab, Inc..sup.41 .sup.41Patterson, R. K.
Analytical Chemistry 1973, 45, (3), 605-609
[0052] 2.2. Processing of Unfilled SE and FGS-SE Nanocomposite.
[0053] An SE network was prepared by end-linking the di-functional
vinyl-terminated PDMS molecules and the tetra-functional
crosslinker tetrakis(dimethylsiloxy silane) with
platinum-cyclovinylmethylsiloxane complex as the catalyst. The
crosslinking resulted from the reaction of terminating vinyl groups
on the PDMS with silicon hydride groups on the
tetrakis(dimethylsiloxy silane).
[0054] The unfilled SE samples were produced as follows:
predetermined amounts of PDMS, the crosslinking agent
tetrakis(dimethylsiloxy)silane and the catalyst
platinum-cyclovinylmethylsiloxane were mixed by magnetic stirring
for 20 min; the mixture was then poured onto a
polytetrafluoroethylene mold and cured at 100.degree. C. for 12 h.
FGS-SE nanocomposites were produced as follows: an FGS suspension
with a concentration of 1 mg/ml was made by mixing a predetermined
amount of FGS and tetrahydrofuran (THF) in a beaker. The beaker was
immersed in an ice bath while the suspension was probe-sonicated
for 30 min (VirSonic 100, The Virtis Co., NY; with an output power
12 W). After sonication, the suspension was transferred to another
beaker containing a desired amount of PDMS polymer. The mixture
containing the FGS, THF, and PDMS was placed on a stir plate heated
to 60.degree. C. to evaporate off all the THF with magnetic
stirring. After all the THF evaporated, the thixotropic mixture was
cooled to room temperature before tetrakis(dimethylsiloxy)silane
and platinum-cyclovinylmethylsiloxane were added. The mixture was
hand-mixed with a steel spatula for 15 min. The final mixture was
then transferred to a polytetrafluoroethylene mold. A metal plate
was used to shear and spread the mixture evenly across the mold.
The shearing velocity of the plate was 6 cm/s. Finally, the mixture
was cured in an oven at 100.degree. C. for 12 h. For 3 wt % FGS-SE
nanocomposites, the samples were prepared using vacuum molding to
minimize trapped air bubbles.
[0055] 2.3. Mechanical Property Measurements.
[0056] Tensile and mechanical hysteresis measurements were made
under ambient conditions using an Instron tensile testing machine
(Model II22, Instron, MA). The dog-bone-shaped samples used in the
measurements were 22.55 mm long and 4.55 mm wide in the narrow
region. Thickness of the samples varied between 0.2 to 0.6 mm. The
strain rate was set to 50.8 mm/min. For the hysteresis
measurements, samples were stretched to 70%-80% of its average
failure strain, returned to a stress level of zero and were
stretched again to a strain level similar to that of the first
stretch. Samples were then placed in an oven set to 100.degree. C.
to recover for 24 h and then their stress-strain curves were
measured again. The area under the stress-strain curve was
calculated and the difference in the area between the first stretch
and the stretch after recovery was obtained. Hysteresis loss ratio
was computed by dividing the difference with the area of the first
stretch. The reported hysteresis loss ratio is an average from
three samples.
[0057] 2.4. Scanning Electron Microscopy Characterization
(SEM).
[0058] Images of cryo-fractured SE or FGS-filled SE were taken with
two different SEMs. Tescan Vega SEM (Tescan USA, PA) was used to
characterize the sample without conductive coatings at
magnifications up to 3700. To obtain high resolution images, the
samples were coated with 3 nm iridium. An FEI XL-30 field emission
gun SEM (Philips, MA) was used to image the samples.
[0059] 2.5. Electrical Conductivity Measurements.
[0060] The direct current transverse resistivity (the resistivity
across the film thickness direction) of the FGS-SE nanocomposites
was measured with a resistivity test fixture (Keithley 8009,
Keithley Instrument Inc., OH) coupled with a digital multimeter
(Keithley 6517). The composite film was cut into a circular film
with a diameter of 70 mm and placed between the top and guarded
electrodes for the measurement. The DC longitudinal resistivity
(the resistivity along the in-plane direction of the film) was
measured using a standard 4-point technique. The nanocomposite film
was cut into rectangular shape films (1-2 cm in width and 2-4 cm in
length). A film was placed on a polystyrene petri-dish and
conductive copper-nickel adhesive tape (Electron Microscopy
Sciences) was placed near the two ends of the film. Conductive
carbon paste (Electron Microscopy Science) was used to draw
conductive paths between the sample and the copper tape. The
resistance was measured with a DC power supply (Tektronix PS2521G,
Tektronix, OR), digital multimeter (Fluke 27, Fluke Corporation,
WA) and electrometer (Keithley 6514). The conductivity of a sample
film was calculated based on the dimension of the film. The
longitudinal conductivity of SE with FGS concentration less than
0.2 wt. % was below the detection limit of the devices and thus
could not be measured. All the electrical conductivities were the
average from two separately made samples.
[0061] 2.6. Gas Permeation Measurements.
[0062] Oxygen and nitrogen permeability of unfilled and FGS-filled
SE was obtained using a constant pressure/variable volume type
permeation cell from Professor Donald Paul's lab at University of
Texas..sup.42 The amount of gas that has permeated was measured and
plotted as a function of time. The permeability was determined from
the slope of the linear portion of the plot (steady state).
.sup.42Takahashi, S.; Goldberg, H. A.; Feeney, C. A.; Karim, D. P.;
Farrell, M.; O'Leary, K.; Paul, D. R. Polymer 2006, 47, (9),
3083-3093
[0063] 2.7. Two Dimensional Viscoelastic Lattice Model.
[0064] A two dimensional viscoelastic lattice model for the
elastomer matrix with the ability to visualize tearing was utilized
to explain the mechanical reinforcement in FGS-SE nanocomposites.
Detailed description of the model is provided elsewhere..sup.43
Briefly, the model is composed of one dimensional trusses arranged
in a two dimensional (the third dimension is of unit thickness)
triangular lattice. A Zener viscoelastic element is used to model
the behavior of each truss. A tear can be initiated when the axial
stress of one truss element exceeds a prescribed breaking stress.
It is known that like other materials, elastomers have intrinsic
defects tens to hundreds of microns in size, that are possibly
introduced while molding or cutting a test sample..sup.44,45
Tearing is first initiated from these defects. Due to the presence
of these weak links, other parts of the matrix may not even be
sampled mechanically when the failure occurs. Such heterogeneity of
breaking stress within the SE matrix is incorporated in the model
by assigning spatially varying breaking stress across the matrix.
.sup.43Sanborn, S. E.; Pan, S.; Prevost, J. H.; Aksay, I. A.
Submitted to Macromolecules .sup.44Choi, I. S.; Roland, C. M.
Rubber Chemistry And Technology 1996, 69, (4), 591-599.sup.45Hamed,
G. R. Rubber Chemistry And Technology 1983, 56, (1), 244-251
[0065] Two domains were simulated in the model: a small
length-scale model with unpercolated FGS and a large length-scale
model with percolated FGS network. The number of trusses was kept
constant. In the small length-scale model, the representative
volume element (RVE) had a similar length-scale to that of the weak
links in the matrix, whose dimension was set to be an order of
magnitude larger than the length scale of individual FGS. In the
large length-scale model, the percolated FGS had a length scale
comparable to the RVE. The matrix was set to have homogeneous
breaking stress since the percolated FGS has much larger length
scale than the heterogeneities. In the FGS-SE nanocomposite model,
individual FGS and percolated FGS were represented by black lines
with a stiffness four orders of magnitude larger than that of the
matrix. FGS itself and the FGS-SE interface does not fail in the
model. One hundred simulations were run for each FGS
concentration.
[0066] The matrix is deformed in the tensile direction at a strain
rate of 0.0076/s, which is the loading rate used in the
experiment.
[0067] Results and Discussion
[0068] 3.1 Characterization of FGS Dispersion.
[0069] To elucidate the effect of filler agglomeration on the
mechanical properties of FGS-SE nanocomposites, SEM was used to
characterize the FGS dispersion state in SE matrix. The images of
cryo-fractured surfaces of unfilled and FGS-filled SE are shown in
FIGS. 1A-1F. The cryo-fractured surface of the unfilled SE without
conductive coating was smooth (FIG. 1A). The morphology of
FGS-filled SE was very different from that of the unfilled SE. As
shown in FIG. 1B, the back-scattered electron SEM image of an
uncoated cryo-fractured-surface of 0.2 wt. % FGS-SE showed the
presence of rough and smooth morphologies. The rough morphology was
likely due to the presence of FGS. In the secondary electron image
of the same area (FIG. 1C), both bright and dark regions were
observed. The dark regions correlated well with the rough regions
in the back-scattered electron image. As the dark regions did not
appear in the same sample with a conductive coating, they are due
to the conductivity variation across the fractured surface. In SEM
imaging, an electron beam bombards the sample and regions with low
electrical conductivity or without conductive pathways would
accumulate charges due to the lack of charge dissipation mechanism
and therefore appear brighter in the image. When regions with
spatially varying conductivities exist in a sample, regions with
higher conductivity would appear to be darker than less-conducting
regions. Since FGS was the only filler in the nanocomposite, the
dark regions must be the percolated FGS-rich regions and the bright
regions were the FGS-lean regions. As shown in FIG. 1D, the
wrinkled morphology of the FGS rich regions resembled that of the
FGS, confirming the agglomeration of FGSs.
[0070] The above evidence suggests that at a 0.2 wt. % FGS loading,
we have a composite material at two length scales: the first one is
the segregation of FGS-rich and FGS-lean regions with a length
scale of 5-15 .mu.m and the second length scale is the ultimate
FGS-SE nanocomposite in the FGS-rich regions.
[0071] At a high enough FGS concentration, the entire sample is
expected to be composed of FGS-rich regions. That was indeed
observed. When the FGS concentration was increased to 0.8 wt. %,
the conductivity induced contrast in the SEM image disappeared,
indicating the existence of FGS-rich regions across the entire
sample (FIG. 1E). The uniform dispersion of FGS in 1 wt. % FGS-SE
was demonstrated in FIG. 1F.
[0072] 3.2 Electrical Properties of FGS-SE Nanocomposite.
[0073] To characterize the percolation threshold of FGS in SE, the
electrical conductivity as a function of FGS weight percentage was
measured and is shown in Error! Reference source not found. At 0.05
wt. % loading, there was no increase in the transverse electrical
conductivity, indicating a percolated FGS network had not yet
formed. The transverse conductivity increased by almost 6 orders of
magnitude when 0.1 wt. % FGS was added. At an FGS loading of only
0.2 wt. %, the longitudinal conductivity increased by more than 10
orders of magnitude and the transverse conductivity increased by 7
orders of magnitude. Anisotropy in electrical conductivity was
observed. The longitudinal conductivity at 0.2 wt. % loading has
already satisfied the conductivity requirement for electrostatic
dissipation (10.sup.-5 S/m) and also the electrostatic painting
applications (10.sup.-4 S/m)..sup.46 At 0.5 wt. % loading, the
electrical conductivity of SE increased to 4.5.times.10.sup.-3 S/m
in the longitudinal direction and to 1.6.times.10.sup.-6 S/m in the
transverse direction. Further increase in the FGS loading above 0.5
wt. % led to a more gradual enhancement in conductivity. The
conductivity reached 0.89 S/m in the longitudinal direction and
6.6.times.10.sup.-4 S/m in the transverse direction at 3 wt. %
loading. .sup.46Ramasubramaniam, R.; Chen, J.; Liu, H. Y. Applied
Physics Letters 2003, 83, (14), 2928-2930
[0074] When conductive fillers form a network of connected paths
through the insulating matrix, a rapid increase in the electrical
conductivity is expected..sup.47 FIGS. 1A-1F suggest that the
percolation threshold was between 0.05 wt. % and 0.1 wt. % as
evidenced by a rapid increase in the tranverse conductivity (almost
6 orders of magnitude) followed by a more gradual increase in
conductivity (1 order of magnitude increase from 0.1 wt. % to 0.2
wt. % FGS). The observed electrical percolation threshold is, to
the best of our knowledge, among the lowest in filled SE, second
only to one case of multiwall carbon nanotube (MWNT)-SE
nanocomposite..sup.17 The conductivity of FGS-SE as a function of
filler concentration is compared to that of other graphene-based
polymer nanocomposites, as shown in Error! Reference source not
found. A-1F. The observed electrical percolation in FGS-SE is lower
than that of graphene-polymer nanocomposites previously
reported..sup.28,30,31,48,49,50,51The longitudinal conductivities
of FGS-SE are comparable to the best MWNT filled SE and graphene
based polymer nanocomposites..sup.28,52 .sup.47Ruschau, G. R.;
Yoshikawa, S.; Newnham, R. E. Journal Of Applied Physics 1992, 72,
(3), 953-95948 Ansari, S.; Giannelis, E. P. Journal of Polymer
Science Part B--Polymer Physics 2009, 47, (9),
888-897.sup.49Nguyen, D. A.; Lee, Y. R.; Raghu, A. V.; Jeong, H.
M.; Shin, C. M.; Kim, B. K. Polymer International 2009, 58, (4),
412-417.sup.50Lang, J. J.; Wang, Y.; Huang, Y.; Ma, Y. F.; Liu, Z.
F.; Cai, F. M.; Zhang, C. D.; Gao, H. J.; Chen, Y. S. Carbon 2009,
47, (3), 922-925.sup.51Steurer, P.; Wissert, R.; Thomann, R.;
Mulhaupt, R. Macromolecular Rapid Communications 2009, 30, (4-5),
316-327.sup.52Khosla, A.; Gray, B. L. Materials Letters 2009, 63,
(13-14), 1203-1206
[0075] The electrical percolation threshold is influenced by the
filler aspect ratio and shape, as well as filler dispersion in the
matrix..sup.53,54,55 A theoretical percolation threshold of plates
having an aspect ratio of 476 is estimated to be 0.27 vol.
%..sup.53 The experimentally observed percolation was less than one
fourth of the theoretically predicted value. The lower percolation
threshold can be attributed to the agglomeration of FGS which
lowers the percolation threshold..sup.54,55 There exists van der
Waals' attraction between FGSs. Since PDMS has attractive
interaction with FGS by forming hydrogen bonds, PDMS chains can
introduce bridging attraction. Both van der Waals's and bridging
attraction contribute to the agglomeration of FGS. In another study
by our group, homogeneous dispersion of FGS in poly(ethylene oxide)
led to a higher percolation threshold of 1 wt. %, corroborating the
agglomeration-induced percolation of FGS in SE..sup.56 Due to the
FGS agglomeration, the percolation threshold of FGS-SE was not
determined using a typical percolation model which assumes
homogeneous distribution of fillers..sup.57 .sup.53Garboczi, E. J.;
Snyder, K. A.; Douglas, J. F.; Thorpe, M. F. Physical Review E
1995, 52, (1), 819-828.sup.54Pegel, S.; Potschke, P.; Petzold, G.;
Alig, Dudkin, S. M.; Lellinger, D. Polymer 2008, 49, (4),
974-984..sup.55Alig, Lellinger, D.; Engel, M.; Skipa, T.; Potschke,
P. Polymer 2008, 49, (7), 1902-1909.sup.56Korkut, S.; et. al.
Manuscript inpreparation..sup.57McLachlan, D. S.; Chiteme, C.;
Park, C.; Wise, K. E.; Lowther, S. E.; Lillehei, P. T.; Siochi, E.
J.; Harrison, J. S. Journal Of Polymer Science Part B-Polymer
Physics 2005, 43, (22), 3273-3287
[0076] The electrical conductivity of a conductive composite is
governed by the intrinsic conductivity of fillers, constriction and
tunneling resistance at the contact between fillers and the number
of contact spots..sup.47,58 Constriction resistance is associated
with constriction of electron flow through the contact area between
two filler particles and is inversely proportional to the contact
area. The morphology of graphene sheet, which is determined by its
functional groups and defects,.sup.59 can influence the contact
area and therefore the constriction resistance in the composites.
.sup.58Simmons, J. G. Journal Of Applied Physics 1963, 34, (6),
1793-&.sup.59Schniepp, H. C.; Kudin, K. N.; Li, J. L.;
Prud'homme, R. K.; Car, R.; Saville, D. A.; Aksay, I. A. Acs Nano
2008, 2, (12), 2577-2584
[0077] Tunneling resistance is due to the tunneling of electrons
through insulating films covering the fillers and it is
proportional to the work function of the conductor, thickness and
dielectric and thermal properties of the film..sup.47,58 The
dielectric constant of the matrix influences the barrier height,
distance of tunneling and therefore the tunneling resistance. Since
the dielectric constant of most materials is a function of
temperature, thermal properties of the matrix also plays a role in
the tunneling resistance..sup.60 The thermal expansion coefficient
of the matrix is also important. Contact force between the filler,
which is determined by the internal stress inside the composite,
strongly influences the tunneling distance and therefore the
overall conductivity of the composite..sup.61 During heat curing of
SE and subsequent cooling to room temperature, volumetric shrinkage
of SE could occur which induced compressive stress between FGS.
Shrinkage from the processing of composites can alter the tunneling
distance..sup.62 .sup.60von Hippel, A. R., Dielectric Materials and
Applications. Technology Press of MIT: Cambridge, 1961.sup.61Li,
L.; Morris, J. E. In Electrical conduction models for isotropically
conductive adhesive joints, 1997; Ieee-Inst Electrical Electronics
Engineers Inc: 1997; pp 3-8.sup.62Zweifel, Y.; Plummer, C. J. G.;
Kausch, H. H. Journal Of Materials Science 1998, 33, (7),
1715-1721
[0078] The number of contact spots between fillers is influenced by
their dispersion. Better dispersion enables more contacts between
graphene sheets, more conductive paths at a given filler
concentration and therefore higher conductivity.
[0079] It is difficult to compare the conductivity of composites
with different fillers and matrices as aforementioned factors can
be different for different systems. Even for graphene-polymer
nanocomposites, the difference in the dielectric and thermal
properties of the matrix can lead to different tunneling
resistance. The morphology, as well as the dispersion of graphene
can be different depending on the functional groups and defects on
graphene.
[0080] Preferential orientation of FGS during the shear molding
process led to fewer contacts in the transverse direction and more
contacts in the longitudinal direction, causing the observed
anisotropy in the nanocomposite conductivity..sup.63 .sup.63Du, F.
M.; Fischer, J. E.; Winey, K. I. Physical Review B 2005, 72, (12),
4
[0081] 3.3 Mechanical Properties of Graphene-SE Nanocomposite.
[0082] The stress-strain curves of unfilled and FGS-filled SE are
shown in Error! Reference source not found., and the values of
modulus, tensile strength and elongation at break at various FGS
weight and volume percentages are shown in Table 1.
TABLE-US-00001 TABLE 1 Values of the modulus, tensile strength and
elongation at break for FGS-SE nanocomposite at various FGS
loadings. To convert from weight percentage to volume percentage,
SE density of 0.97 g/cm.sup.3 and FGS density of 2.25 g/cm.sup.3
were used. FGS FGS Elongation weight volume Modulus, E at break,
.epsilon..sub.b Strength, .epsilon..sub.b % % (MPa) (%) (MPa) 0 0
1.33 .+-. 0.12 74 .+-. 16 0.57 .+-. 0.09 0.05 0.022 1.42 .+-. 0.13
66 .+-. 11 0.52 .+-. 0.06 0.2 0.088 1.64 .+-. 0.17 138 .+-. 17 0.90
.+-. 0.09 0.5 0.22 1.77 .+-. 0.11 139 .+-. 15 1.30 .+-. 0.10 0.8
0.35 1.93 .+-. 0.12 149 .+-. 40 1.49 .+-. 0.0.36 1 0.43 2.13 .+-.
0.13 138 .+-. 19 1.78 .+-. 0.21 3 1.31 4.86 .+-. 0.44 112 .+-. 9
3.43 .+-. 0.22
[0083] For each FGS loading, the crosslinker and the catalyst
concentration were chosen to yield the highest tensile strength of
the samples. At 0.05 wt. % (0.022 vol. %) FGS loading, no
improvement in mechanical properties was observed. At 0.2 wt. %
(0.086 vol. %), a 23% increase in the modulus, an 87% increase in
the elongation at break and a 58% increase in the tensile strength
were observed. At 0.5 wt. % (0.22 vol. %) FGS, a 33% increase in
the modulus, 87% increase in elongation at break and a 128%
increase in the tensile strength were achieved. At a 3 wt. % (1.34
vol. %) FGS loading, the modulus increased by 265%, the elongation
at break increased by 51% and the tensile strength increased by
over 500%. Above the percolation threshold, the modulus and tensile
strength increased with the FGS concentration whereas the
elongation at break increased initially with FGS content up to 1
wt. % (0.45 vol. %) and then decreased at higher FGS loadings.
There appeared to exist a critical FGS concentration between 0.5
wt. % and 1 wt. % beyond which the elongation at break of the
composite started decreasing.
[0084] To understand the effect of FGS on the tensile properties of
SE, movies of the tearing process of unfilled and FGS-filled SE in
an SEM were recorded to reveal the failure mechanisms. Snapshots of
the movies are shown in Error! Reference source not found. A-4F. As
shown in FIG. 4A, when the unfilled SE was deformed, the notch
gradually opened up. Until a certain level of stress was reached,
tearing was initiated from the tip of the notch due to stress
concentration and it immediately propagated across the specimen
with little resistance, leading to the failure of the specimen. Due
to its lack of ability to crystallize under strain, SE does not
possess mechanism to arrest or deflect tearing and transfer the
mechanical load to other parts of the matrix that is not sampled
mechanically. When percolated FGS network was introduced, the
failure mechanism of SE was altered dramatically. At the initial
stage of the deformation of 0.5 wt. % FGS-SE, the notch opened up
(FIG. 4B). At a certain stress level, tearing was initiated from
the tip of the notch. However, unlike the case of the unfilled SE,
tear propagation was resisted and the sample did not fail upon tear
initiation (FIG. 4C). The percolated FGS network introduced
resistance for tear propagation. Upon further deformation, tear was
further opened, followed by the catastrophic failure of the sample
(FIG. 4D). The observations above clearly illustrated the enhanced
tear resistance in SE introduced by the percolated FGS network.
[0085] SEM was also used to characterize the tensile-fractured
surface of unfilled and FGS-filled. For the unfilled SE (FIG. 4E),
a few ridges were observed and majority of the torn surface was
smooth, indicating that once tearing was initiated, it propagated
across the entire sample with little resistance or distortion. The
fractured surface of 0.5 wt. % FGS-SE is shown in FIG. 4F. The
bright spots with submicron length scales were the FGS. The
morphology of the failure surface was quite different from that of
the unfilled SE. The torn surface had more ridges than the unfilled
SE, indicating the distortion of tear propagation by the presence
of percolated FGS.
[0086] To quantify the degree of tearing in the unfilled and
FGS-filled SE, mechanical hysteresis measurement was undertaken.
Hysteresis loss in filled rubber has been attributed to covalent
bond rupturing in the matrix,.sup.64 viscoelasticity.sup.65 and the
breakdown of the filler network structure.sup.66. While
viscoelasticity induced hysteresis loss is recoverable, covalent
bond rupturing and the breakdown of filler network structure can
lead to irrecoverable hysteresis loss. Given the nature of FGS-FGS
interaction to be weak van der Waals' force, the contribution of
FGS network breakdown to irrecoverable hysteresis loss in FGS-SE is
likely to be small. Therefore, measurements of irrecoverable
hysteresis loss provide a method to quantify the degree of tearing
in unfilled and FGS-filled SE. .sup.64Suzuki, N.; Ito, M.;
Yatsuyanagi, F. Polymer 2005, 46, (1), 193-201.sup.65Roland, C. M.
Rubber Chemistry and Technology 1989, 62, (5),
880-895.sup.66Yamaguchi, K.; Busfield, J. J. C.; Thomas, A. G.
Journal of Polymer Science Part B-Polymer Physics 2003, 41, (17),
2079-2089
[0087] The mechanical hysteresis data is shown in Error! Reference
source not found. Unfilled SE showed little irrecoverable
hysteresis loss after recovery as evidenced by the overlapping of
the stress-strain curves of first stretch and the stretch after
recovery (Error! Reference source not found.), suggesting few
covalent bonds rupturing in the matrix. The 0.5 wt. % FGS-SE sample
showed an observable difference in the two stress-strain curves and
a higher irrecoverable hysteresis loss compared to the unfilled SE,
as shown in Error! Reference source not found. The irrecoverable
hysteresis loss ratio is determined to be 5.5%. The above results
suggest a higher degree of covalent bond rupturing in the matrix
caused by the presence of FGS, confirming the introduction of
distributed deformation by adding FGS.
[0088] Based on the above evidence, the reinforcement mechanism of
SE by FGS can be envisioned. To improve the failure properties of
SE, a mechanism to arrest or distort the tearing initiated from the
intrinsic defects is necessary. The arresting or distortion of
tearing could only be achieved when the length scale of the filler
was much larger than the length scale of the initial tear size,
which is determined by the size of intrinsic defects. An individual
FGS has a lateral size of hundreds of nanometers, much smaller than
the intrinsic defect size in elastomers..sup.44,45 Therefore, only
percolated FGS network leads to a simultaneous increase in tensile
strength and elongation at break of SE. During deformation of the
nanocomposite, tearing is initiated from the intrinsic flaws in the
matrix and arrested or deflected by the presence of percolated FGS.
The arresting or deflection leads to load transferring to the FGS
and other parts of the unstrained matrix, leading to the
enhancement of tensile strength. Deformation and tearing are also
distributed across a larger portion of the matrix compared to
unfilled SE and the opening of tears causes dilation within the
matrix, leading to the observed elongation at break increase.
[0089] The necessity of a percolated FGS network to improve
mechanical properties shown here is in sharp contrast to previous
studies of multiwall carbon nanotube (MWNT) and carbon black (CB)
filled SBR in which the increase in strength and elongation at
break occurred prior to electrical percolation..sup.67,68 The MWNT
was shown to have length up to 5 .mu.m, whereas CB can form
agglomerates up to hundreds of microns..sup.69 Agglomeration of
those fillers can readily achieve a length-scale that is larger
than the critical flaw size and improves the tensile properties.
.sup.67Bokobza, L. Polymer 2007, 48, (17), 4907-4920.sup.68Reffaee,
A. S. A.; El Nashar, D. E.; Abd-El-Messieh, S. L.; Nour, K.
Polymer-Plastics Technology and Engineering 2007, 46, (6),
591-603.sup.69Kohjiya, S.; Kato, A.; Ikeda, Y. Progress in Polymer
Science 2008, 33, (10), 979-997
[0090] 3.4 Modeling of Mechanical Reinforcement in Graphene-SE
Nanocomposite
[0091] To corroborate with the reinforcement mechanism demonstrated
above and more importantly, to understand the reversal in
elongation at break, a lattice-based model is used to study the
deformation mechanics of FGS-SE nanocomposite..sup.43 The deformed
lattice with or without FGS is shown in Error! Reference source not
found. A-6D. Individual FGS or percolated FGS was represented by
the black lines. The torn matrix was represented by green regions.
In the unfilled SE, tearing was initiated from the defects and
propagated across the sample without much resistance due to the
lack of tear arresting mechanism (FIG. 6A). As shown in FIG. 6B,
when individual unpercolated FGSs were present outside of the
defects, tearing was initiated from the intrinsic defects and
propagated across the sample undeterred as in the case of unfilled
SE. No interaction between FGS and tears was observed and failure
properties of SE were therefore not improved. When FGS was
percolated, the percolated network had a length scale much larger
than the intrinsic defect size. When 1.6 vol. % percolated FGS was
added, tearing was initiated and distorted or arrested by the
presence of percolated FGS (FIG. 6C). Through arresting or
distorting of tearing, deformation was distributed to the stronger
parts of the matrix as evidenced by the distributed tearing. When
FGS concentration was increased to 5.2 vol. %, a higher degree of
distributed deformation can be achieved as indicated by the
increased amount of torn matrix (FIG. 6D).
[0092] The simulated and normalized stress-strain curves of
unfilled and FGS-filled SE are shown in Error! Reference source not
found. A-7C. The termination of the stress-strain curve indicated
the strain at the peak stress level. Peak stress was defined as the
tensile strength of the sample and the strain at the peak strength
was defined as the elongation at break. The simulation reproduced
qualitatively the experimental stress-strain curves of unfilled and
FGS-filled FGS. The modulus and strength increased with FGS
concentration whereas the elongation at break increased initially
and decreased above 5.2 vol. %.
[0093] Analysis of the mechanical load carried by FGS demonstrated
the excellent load carrying capacity of percolated FGS..sup.43
Through arresting or distorting of tearing, mechanical load was
transferred to the FGS and stronger parts of the matrix, leading to
a significant increase in the tensile strength of the composite.
The most interesting observation was the increase in the elongation
at break of SE. To gain more insight into the mechanism responsible
for the elongation at break increase, we studied the fraction of
matrix torn and the strain of tears as a function of FGS
concentration. The fraction of matrix torn increased with FGS
concentration, suggesting an increasing degree of distributed
deformation. The strain of tears decreased with FGS concentration,
due to the close proximity of percolated FGS suppressing the
opening of tears. The increase in elongation at break with FGS
concentration can be explained by the dilation effect of tearing.
As tear opened up during straining of the sample, the sample could
be elongated more. Our model showed that the elongation at break of
the nanocomposite was dominated by two factors: the fraction of
matrix torn and the strain of tears. When the FGS concentration was
increased, the two factors were competing with each other and the
reversal effect on the elongation at break with increasing FGS
concentration was a result of the domination of decreasing strain
of tears over increasing fraction of matrix torn.
[0094] The elongation at break increase can also be influenced by
other factors such as the strain and deformability of the untorn
matrix. However, those effects cannot be investigated due to the
technical limitation of the model.
[0095] To demonstrate the superiority of FGS at improving the
mechanical properties of SE, modulus, tensile strength and their
relative improvement of all filled PDMS-based SE were plotted
against the filler volume fraction for the concentration range
studied in the present study (Error! Reference source not found.
A-8B and Error! Reference source not found. A-9B). Only studies
that used pristine SE (containing no fillers) as the base polymer
were chosen for the analysis. For each type of filler, the sets of
data with the largest improvement in mechanical properties were
chosen for comparison.
[0096] To compare the modulus enhancement brought by FGS to that by
other fillers, the modulus and improvement in modulus (calculated
by dividing the difference in modulus between the filled and
unfilled samples by the modulus of the unfilled sample) were
plotted against filler volume percentage in Error! Reference source
not found. A-8B. In terms of the relative improvement of the
modulus, FGS is comparable to or better than all the reported
fillers except for MWNT. One thing needs to be noted is that the
unfilled SE in that MWNT-SE study had a modulus of 0.14 MPa, almost
an order of magnitude lower than that of the unfilled SE used in
the present invention..sup.17 The superiority of FGS at enhancing
modulus can be attributed to the high aspect ratio plate-like
geometry and the high surface area (higher than all the fillers
reported in previous studies) which enabled the low percolation
threshold and offered extensive interfacial interactions with the
matrix and a higher degree of load transferring.
[0097] The strength and the improvement in strength of the FGS-SE
nanocomposite were compared with those of other filled-SE as a
function of filler volume percentage, shown in Error! Reference
source not found. A-9B. In terms of the relative improvement in
strength, FGS performed comparably or better than all other fillers
except for MWNT. In the case of fumed silica and precipitated
silica, which yielded similar strength improvement as FGS, the
unfilled SE in that study had a tensile strength of 0.075
MPa,.sup.15 much lower than that of the unfilled SE in the present
study (0.57 MPa). FGS-SE also has the highest tensile strengths in
the concentration range studied. The superior ability of FGS to
strengthen the matrix is believed to be attributed to: 1. its
plate-like geometry, high aspect ratio and surface area, which
provide a low percolation threshold and large load transferring; 2.
The distributed deformation introduced by the percolated FGS
allowed more regions of the SE matrix to carry loads.
[0098] 3.4 Barrier Properties of FGS-SE Nanocomposite.
[0099] The multi-functionality of FGS as a filler lies in its
ability to simultaneously improve the mechanical and electrical as
well as the barrier properties of SE.
[0100] Oxygen and nitrogen permeabilities of unfilled and
FGS-filled SE were measured and the results are shown in Table
2.
TABLE-US-00002 TABLE 2 Oxygen and nitrogen permeability of unfilled
and FGS-filled SE. Permeability was reduced by half with 3 wt. %
(1.31 vol. %) FGS. Permeability (Barrier) Sample O.sub.2 N.sub.2
Unfilled SE 555 266 1 wt. % (0.43 vol. %) FGS 514 249 3 wt. % (1.31
vol. %) FGS 283 137
[0101] With the incorporation of 1 wt. % (0.43 vol. %) FGS,
permeability for both gases was reduced by 7%. When 3 wt. % (1.31
vol. %) FGS was added, permeability was reduced by half. The
improvement was better than that in the clay-filled SE..sup.18
[0102] The reduction in gas permeability is believed to be
attributed to the presence of FGS acting as impermeable barrier and
increasing the diffusion path for the gas..sup.70 Additionally, it
has been suggested that due to the large interfacial area in the
nanocomposites, the properties of the matrix, such as the fraction
free volume, can be reduced and further decrease in the
permeability can be achieved..sup.71 PDMS can form hydrogen bonding
with hydroxyl groups and the interaction between PDMS and FGS
provides a modification of the matrix permeability and therefore
the overall barrier property of the nanocomposite. .sup.70Nielsen,
L. E. Journal of Macromolecular Science 1967, A1, (5),
929-942.sup.71Wang, Z. F.; Wang, B.; Qi, N.; Zhang, H. F.; Zhang,
L. Q. Polymer 2005, 46, (3), 719-724
[0103] Obviously, numerous modifications and variations of the
present invention are possible in light of the above teachings. It
is therefore to be understood that within the scope of the appended
claims, the invention may be practiced otherwise than as
specifically described herein.
* * * * *