COMPOSITES

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Composites Science and Technology 66 (2006) 2135–2142

SCIENCE ANDTECHNOLOGY

Multiscale fiber-reinforced nanocomposites:Synthesis, processing and properties

Derrick Dean a,*, Apollo M. Obore b, Sylvester Richmond b, Elijah Nyairo c

a Department of Materials Science and Engineering, University of Alabama Birmingham, 1530 3rd Ave. S., Birmingham, AL 35294, United Statesb Center for Advanced Materials, Tuskegee University, 101 Chappie James Center, Tuskegee, AL 36088, United States

c Department of Physical Sciences, Alabama State University, 915 South Jackson Street, Montgomery, AL 36101, United States

Received 27 June 2005; received in revised form 2 November 2005; accepted 15 December 2005Available online 6 March 2006

Abstract

Multiscale fiber-reinforced nanocomposites have been manufactured using a vacuum assisted resin infusion molding (VARIM) pro-cess. The nanocomposites prepared were epoxy resin prepolymers dispersed with layered silicates. The effect of silicate loading on theflow, isothermal cure behavior and solid state properties was studied. The addition of silicate results in an increase in resin viscosityby a factor of 10, although it is still within the range of processability by VARIM. A slight decrease in resin gel times is also observed.X-ray diffraction studies of Epoxy 815C nanocomposites indicated an exfoliated morphology for the 2% silicate and an intercalated mor-phology for the 4% and 6% specimens. Dynamic mechanical analyses were conducted to establish the Tg of the specimens. An improve-ment of 31% in flexural modulus and 24% in flexural ultimate strength for the 2% silicate fiber-reinforced nanocomposites was achieved.Interlaminar shear and fracture toughness studies were also conducted, however no enhancement was observed.� 2006 Published by Elsevier Ltd.

Keywords: Nanocomposite; Rheology; Morphology

1. Introduction

Fiber-reinforced polymeric composites (FRC) haveshown great promise as high strength structural materialsdue their high stiffness to weight ratio and ease of process-ability and have found use in applications ranging fromconstruction to aerospace materials, and recreationalequipment [1]. Research in nanostructured composites ornanocomposites, in which a resin is dispersed with nano-particles, such as layered silicates, carbon nanotubes orcarbon nanofibers, has shown explosive growth in the pastdecade [2–8]. Incorporation of nanoscale constituents intocomposite matrices leads to new or modified propertyenhancements significantly greater than that attainableusing conventional fillers or polymer blends. Improvementsinclude [1], for example, decreased permeability to gases,

0266-3538/$ - see front matter � 2006 Published by Elsevier Ltd.

doi:10.1016/j.compscitech.2005.12.015

* Corresponding author. Tel.: +1 205 975 4666; fax: +1 205 934 8485.E-mail address: deand@uab.edu (D. Dean).

water and hydrocarbons, thermal stability and heat distor-tion temperature, flame retardancy and reduced smokeemissions, chemical resistance, surface appearance, electri-cal conductivity and optical clarity. A synergistic combina-tion of polymeric nanocomposites results in a multiscalecomposite with a hierarchal structure ranging from nano-scale particles to micron size fibers. These multiscale com-posite offer a route by which multifunctionality (enhancedthermal stability, lower coefficient of thermal expansion,high thermal and electrical conductivity) can be impartedto the FRC. One route for fabricating multiscale FRC isto infuse fiber preforms with nanostructured prepolymers,or nanocomposites. While a significant amount of workhas been studied on thermoplastic-based polymer nano-composites, comparatively few studies on thermoset-basedsystems have been extensively investigated [9,3,10–28]. Forboth systems, the precise mechanism(s) by which the prop-erty enhancements reported are achieved remains a topic ofintense research.

2136 D. Dean et al. / Composites Science and Technology 66 (2006) 2135–2142

Layered silicate nanoparticles are well suited for thedesign of hybrid composites because they have high in planestrength, stiffness and a high aspect ratio. The platelet ratioexceeds 300, giving rise to a high degree of polymer–silicatesurface interaction which results in barrier and mechanicalproperties that are far superior to those of the base material.

Relatively few studies of fiber-reinforced nanocompositeshave been reported. Chen et al. prepared an epoxy-silicatenanocomposite using an aerospace grade epoxy resin andcarbon fibers [29,30]. The nanocomposites were processableusing vacuum assisted resin infusion; no improvements in themechanical properties were reported, however. We haverecently reported on the synthesis and fabrication of glassfiber reinforced nanocomposites based on vinyl ester andepoxy thermoset nanocomposites, respectively [31,32]. TheS2-glass/epoxy-silicate nanocomposites were manufacturedthrough an affordable vacuum assisted resin infusionmethod (VARIM). Basic correlations between polymer mor-phology, strength, modulus, toughness, and thermal stabil-ity of thermoset nanocomposites were investigated as afunction of layered silicate content. Transmission electronmicroscopy (TEM) and wide angle X-ray diffraction(WAXD) were used to characterize the morphology of thedispersed silicate particles. The thermal properties such asonset of decomposition and glass transition temperatureswere determined by thermo gravimetric analysis (TGA)and dynamic modulus analyzer (DMA). Mechanical proper-ties such as interlaminar shear strength, flexural propertiesand fracture toughness were also determined for both con-ventional S2-glass/epoxy composites and S2-glass fiber-rein-forced nanocomposites. The results show significantimprovements in mechanical and thermal properties of con-ventional fiber-reinforced composites with low loading oforgano silicate nanoparticles. By dispersing 1% by weightnanosilicates, S2-glass/epoxy-silicate nanocomposites con-tributed to almost 44%, 24% and 23% improvement in inter-laminar shear strength, flexural strength and fracturetoughness in comparison to conventional S2-glass/epoxycomposites. Similarly, the nanocomposites exhibit approxi-mately 26 �C higher decomposition temperatures than thatof conventional composites. These improved properties offiber-reinforced polymer nanocomposites are achievedmostly due to increased interfacial surface areas, improvedbond characteristics and intercalated/exfoliated morphol-ogy of the epoxy-silicate nanocomposites.

Herein we report on the synthesis and fabrication ofmultiscale fiber-reinforced composites using carbon fiberpreforms infused with nanodispersed epoxy resins. Wehave utilized rheology to investigate the effect of the nano-particles on the processability and correlated this with ulti-mate properties (e.g., strength, stiffness, toughness andinterlaminar shear strength).

2. Experimental

Epon 815C (Modified Bisphenol A), an Epichlorohydrinbased Epoxy resin and Epicure Curing Agent W (Aromatic

diamine) were purchased from the Miller-StephensonChemical Company. Cloisite 30B, montmorillonite silicateorganically modified with a ternary ammonium salt (alkylquarternary ammonium montmorillonite) was suppliedby Southern Clay Company, and 8 Harness satin weavecarbon fabric was supplied by NASA Glenn Research Cen-ter. Ten gram of epoxy resin was weighed out and the cor-responding ratio of silicate by weight (2–6 wt%) added. Theresin and silicate were initially mixed for a period of 20 minusing a mechanical stirrer. Two gram of the curing agentwas added and the mixture was mixed for an additional5 min. This mixture was then infused into a carbon fabriclay up by the VARIM process and left under vacuum for1 h. The fabric was then removed and consolidated usinga Carver hydraulic hot press, at 150 �C for 4 h. The pristineepoxy-silicate nanocomposites were prepared by takingportions of mixed resin, silicate, and curing agent, castinginto a plastic mold, degassing them in a vacuum ovenand then curing at 150 �C for 4 h.

Flexural tests were conducted on the Minimat 2000instrument. A crosshead speed of 0.5 mm min�1 and spanlength 18.5 mm was used according to ASTM D790. Inter-laminar shear tests on the carbon–epoxy composites weredone according to ASTM D-2344 test method. A crossheadspeed of 1.3 mm (0.05 in.)/min was used.

2.1. Morphology/characterization

An M8 optical microscope equipped with a Kodak 290zoom digital camera was used to capture optical micro-graphs. Microstructure and fracture surfaces of the com-posites were evaluated by scanning electron microscopy(Model JSM5800) at an accelerating voltage of 3 kV witha tungsten filament and a coating of Au–Pd.

X-ray diffraction scans were conducted using a RigakuD/MAX-2200 diffractometer generator with Cu Ka radia-tion (k = 1.54 A) at an accelerating voltage of 40 kV/30 mA. Rheological experiments were performed on thenanodispersed prepolymers using a TA Instruments AR2000 Rheometer. Both flow experiments and time sweepswere performed. The glass transition temperature wasobtained using a Dynamic Mechanical Analyzer (TAInstrument DMA 2980 Dynamic Mechanical Analyzer)using a single cantilever beam fixture. The temperaturewas ramped from 250 to 250 �C at 3.0 �C/min. The Tg

was determined by the corresponding peak of the loss mod-ulus (G00) curve.

3. Results and discussion

3.1. Morphology-wide angle X-ray diffraction (WAXD)

The interaction between silicate and the polymer matrixresults in the incorporation of the polymer between the lay-ers of the silicate, which increases the interlayer distance.The extent of the dispersal was determined by wide angleX-ray diffraction (WAXD). Bragg’s equation is used to

D. Dean et al. / Composites Science and Technology 66 (2006) 2135–2142 2137

measure the interlayer distance using the diffraction peakand the position of d0 0 1 in the WAXD patterns and is typ-ically shown as sinh = nk/2d0 0 1, where the subscripts indi-cate that specific diffraction angles correspond to specificinterlayer spacing. d0 0 1 is the interlayer distance of the(001) diffraction face or plane, h is the diffraction angleand k is the wavelength (1.54 A). The WAXD data forepoxy is displayed in X–Y plots of the intensity of thereflections (in counts or arbitrary units) versus the scatter-ing angle, 2h, in Fig. 1. The scan for the Cloisite 30B cor-responds to a basal spacing of 17.83A (2h = 5) which is theinterlayer distance between the silicate layers of the mont-morrilonite silicate and has been indexed as the 001 reflec-tion. A sharp peak for the 4% and the 6% silicate samplescorresponds to d-spacing of 36.7 A and 33.5 A, respec-tively. These are indicative of an intercalated morphology.The additional peak in the scan for the 4% sample can beattributed to some phase separation domains of the pristinesilicate. This is conceivable at higher loadings of the silicatein the epoxy matrix. There is no peak for the 2% silicate.This is presumably attributable to an exfoliated morphol-ogy, although TEM is needed to confirm this proposition.As the intergallery spacing increases, the characteristicbasal spacing shifts to smaller angles. This provides a con-venient measure, up to approximately 9.0 nm, after which apeak is no longer visible. At this point, transmission elec-tron microscopy (TEM) which provides local, real spaceimaging becomes useful for characterizing the dispersionof the silicates on a more local scale. The information pro-

Fig. 1. X-ray diffractograph o

vided by this technique complements the WAXD data,which provides globally averaged reciprocal space imaging.

3.2. Rheology

While the morphology of polymer nanocomposites andits development has been studied extensively, their rheol-ogy has received considerably less attention. Furthermore,most of that attention has been focused on the rheologicalbehavior of thermoplastic-based nanocomposites [33–36],with a few concentrated on thermoset systems [18,21,27].Krishnamoorti et al. have studied the effect of silicate load-ing on the melt state rheological properties of polycarbon-ate nanocomposites [37]. A decrease in frequencydependence of the dynamic moduli was observed for highsilicate loadings. In addition, the highest silicate loadingsexhibited a transition in relaxation behavior from liquid-like to pseudo-solid like. In thermosets, the chemical andphysical structure of the network is closely related to theformation process (the reaction between the small moleculeprepolymers to form the final crosslinked macromolecule),which in turn may be affected by the presence of nanopar-ticles. Therefore, understanding issues such as the effect ofthe silicates on the gel points and the kinetics of the cure isessential to developing the processing parameters. We haveconducted a series of rheological experiments designed toprovide insight into the cure behavior of the system underinvestigation. Fig. 2 shows viscosity versus shear rate, anddepicts the effect of the nanoparticles on the viscosity of the

f epoxy nanocomposites.

0.01000 0.1000 1.000 10.00 100.0shear rate (1/s)

0.1000

1.000

10.00

100.0

visc

osity

(P

a.s)

Epon 815C002-0005f, Stepped flow stepEpon 815C_4%C30-0001f, Stepped flow stepEpon 815C_2%C30-0001f, Stepped flow stepEpon 815C_6%C30-0001f, Stepped flow step

Fig. 2. Viscosity of pure epoxy resin and samples containing 2%, 4% and6% silicate, respectively.

150000 2500 5000 7500 10000 12500

time (s)

1.000E7

0.01000

0.1000

1.000

10.00

100.0

1000

10000

1.000E5

1.000E6

G' (

Pa)

815CNEAT@120-2 815CTS2%30B#2-,3 EPON815C-4%CLO30B#2@1204 815C-6%CLAY@120-

22222222222222222222222

22

2

2222222222222

222222222222222222222222222222222222222

2222222

222222222222

222222222222222222222222

222222222222222222222222222222222222

2222222222222222222222222222222222222222222222222222222222

3 3

3

33

3333333333333333

3333333

333333

333333

333333

333333

3333333333333333333333333333333333

333333333

333

44444444

4

4444

444444444444

4444444444

44444444

44444444

44444444

4444444

4444444444444444444444444444444

444444

444444444

4444

Fig. 3. Storage modulus overlay for various silicate loadings at 120 �C.

150000 2500 5000 7500 10000 12500

time (s)

1.000E7

0.1000

1.000

10.00

100.0

1000

10000

1.000E5

1.000E6G

' (P

a)

815CNEAT@150-2 815C-2%C30B@150-3 815C-4%30B#2@1504 815C-6%C30B#2@150

22

222222

2

2

2

2

2222222222222222222222

2222222222222222

222222222222222222222222222222222222222222222222222222222222222222222222

2222222222222222222222222222222222222222222222222222222222222222222222222222222222222222222222222222222222222

333333

3333333333333333333

3

3

3

3

33333333333

3333333333

33333333333333333333333333333333333333333

33333333333333333333333333333

4444

444444

44444444444444444

4

4

4

4

44444444444

44444444

4444444444444444444

Fig. 4. Storage modulus overlay for various silicate loadings at 150 �C.

2138 D. Dean et al. / Composites Science and Technology 66 (2006) 2135–2142

epoxy resin. The neat resin prepolymer behaves as a vis-cous fluid, and can be described by Newton’s law, in whichthe shear stress is directly proportional to the shear rate(i.e. viscosity is independent of the shear rate), and exhibitsa value of 1 Pa s (10 poise), which is well within the rangesuitable for the VARIM process [1]. Examination of thelow frequency data allows delineation of differences instructure as a function of composition. As the silicate load-ing increases, the viscosity increases significantly, howeverits magnitude still suggests processability via VARIM. Ofparticular note is the transition from Newtonian flow topseudo-plastic or shear thinning behavior as the silicateloading increases to 2% and 4%, respectively, in the fre-quency range of 0.05–1 Hz. Shear thinning is typicallyexhibited by polymers, however, the presence of the layeredsilicates and the resulting prepolymer–nanoparticle interac-tions may cause this behavior in the relatively low molecu-lar weight neat resin. Isothermal rheological studies wereconducted to determine the effect of the nanoparticles onthe cure behavior. The curves, shown in Figs. 3 and 4, showa viscosity minimum, followed by a rapid increase and agradual plateau, due to gelation and vitrification, respec-tively. The crossing of the storage and loss storage moduli,G 0 and G00, respectively was taken as the gelation point[21,24,38]. The viscous behavior of the oligomeric materialdominates in the initial part of the polymerization. Withincreasing molecular weight, the loss modulus increaseswhile the storage modulus rises sharply until it intersects,then exceeds the loss modulus, becoming more elastic innature. As the silicate loadings increase, the viscosity min-imum is no longer observed, and the magnitude of G 0

increases, indicative of more elastic behavior. The G 0 curvesfor the 4% and 6% samples exhibit an additional effect notseen in the curves for the pure and 2% sample. The slopechange in the G 0 curves at 2000 s for the sample cured at150 �C and 4500 s for the sample cured at 120 �C suggestsa two-stage curing mechanism may be occurring. Thiseffect may possibly be due to a partitioning effect of the

resin and curing agent caused by the larger amount of sil-icate present in the system. This effect could potentially off-set the reaction stoichiometry and hence affect the crosslinktopology, which in turn would affect the ultimate proper-ties. The slope changes may also be due to intragalleryexpansion, which is halted by the more global extragalleryexpansion. Once the extragallery expansion starts, theincrease in modulus will prevent the silicate layers fromexpanding further, and this intercalated structure will beessentially locked in, once the system gels. This reasoningcould explain the intercalated morphologies for the 4%and 6% systems, as determined by XRD in Fig. 1. It shouldbe noted that the entire curing process in the presence ofthe silicates is complicated, and probably involves a combi-nation of the two processes discussed above.

The gel times for Epon resin 815C with different silicateloadings were determined at two different temperatures byisothermal time sweeps. The gel points were obtained fromthe intersection of the storage and loss modulus as dis-cussed above. Table 1 shows the results of the rheologytests. It is observed that the gel times at the higher temper-ature are considerably shorter than at 120 �C. This isattributable to the higher reactivity of the epoxy systems

Table 1Gelation time (min) at 120 and 150 �C

Sample Time at 120 �C Time at 150 �C

Neat 94.9 27.22% Silicate 90.6 24.24% Silicate 92.3 26.26% Silicate 87.2 26.2

100

1000

10000

Loss

Mod

ulus

(M

Pa)

1000

10000

1.0E5

Sto

rage

Mod

ulus

(M

Pa)

0 50 100 150 200 250

Temperature (˚C)

–––––– 6815C_6%– – – 0815C_Neat–––– · 2815C_2%–– – – 4815C_4%

Fig. 6. DMA overlay of fiber-reinforced epoxy nanocomposites.

128.83˚C

124.98˚C

124.41˚C

123.33˚C

10

100

1000

10000

Loss

Mod

ulus

(M

Pa)

0 50 100 150 200 250

Temperature (˚C)

– – – 6815C_6%.001– –– 0815C_neat.001– –– 2815C_2%.001– –– 4815C_4%.001

D. Dean et al. / Composites Science and Technology 66 (2006) 2135–2142 2139

at higher temperature. Figs. 3 and 4 show a markeddifference in times at which the storage modulus starts torise rapidly, probably indicative of varying duration tothe onset of extragallery polymerization. Increasing theamount of silicate catalyzes this reaction slightly. Theincrease in storage modulus is mostly due to the vitrifica-tion of the polymer matrix [38].

3.3. Dynamic mechanical analysis (DMA)

Dynamic mechanical analysis (DMA) is a powerfultechnique for studying the viscoelastic behavior of polymerbased materials. It is also the preferred method of measur-ing the glass transition temperature (Tg), particularly forpolymers with rigid backbones. Experiments involve apply-ing an oscillatory stress to the sample while monitoring theresulting strain, which consists of both in-phase and out-of-phase components. The strain can then be used to calculatethe in-phase response, or storage modulus (G 0), and theout-of-phase, or the loss modulus (G00). The ratio of G 0/G00 is known as the tan delta. The Tg is determined bythe corresponding peak of the loss modulus curve. TheDMA data for the epoxy/silicate and epoxy/silicate fiberssamples is shown in Figs. 5–7.

Fig. 5 shows the temperature dependence of the storagemodulus G 0, for the neat composite and the epoxy-silicatecomposites containing 2%, 4%, and 6% silicate. Below

0.0001

0.001

0.01

0.1

1

10

100

1000

Loss

Mod

ulus

(M

Pa)

1

10

100

1000

10000

Sto

rage

Mod

ulus

(M

Pa)

0 50 100 150 200 250

Temperature (˚C)

–––––– 815C6%CLAY&RESIN– –– 815CNEAT–––– · 815C2%CLAY&RESIN–– – – 815C4%CLAY&RESIN

Fig. 5. Storage and loss modulus curves overlay of 815C resin/silicatenanocomposite specimens.

Fig. 7. Loss modulus curves overlay of 815C fiber-reinforced composites(staggered for clarity).

Tg, the 2% and 6% composites exhibit higher storage mod-ulus than the neat polymer; this continuing to the rubberystate, while the 4% sample exhibits no change in modulus.The Tg increases from 125 �C to a high of 136 �C for the2% sample, with values of 131 and 130 for the 4% and6% samples, respectively. These values are summarized inTable 2. The large Tg increase at 2% silicate loading is pre-sumably due to better dispersion of the silicate layers asinferred from the X-ray diffraction data in Fig. 1. The large

Table 2Results of Tg values measured from the DMA tests

Sample Tg (�C) fiber-reinforcedcomposites

Tg (�C) resin/silicatespecimen

Neat 124.98 125.162% Silicate 124.41 136.824% Silicate 123.33 131.466% Silicate 128.83 130.83

Table 4Flexural test results for 815C-4 ply nanocomposites cured at 150 �C for 4 hunder hot press

Sample E (Gpa) Standarddeviation

Ultimatestrength (Mpa)

Standarddeviation

Neat 69 0.56 1057 822% C30B 91 9.88 1310 764% C30B 59 6.5 878 1106% C30B 72 6.65 808 70

2140 D. Dean et al. / Composites Science and Technology 66 (2006) 2135–2142

surface area of the exfoliated structures presumably leadsto more nanoparticle–polymer interactions which serve torestrict the polymer segmental mobility, leading to a higherTg. Very little change in Tg is observed for the fiber-rein-forced samples, as summarized in Table 2. The only sampleexhibiting an increase is the 6% sample.

3.4. Flexural tests

The results of the flexural test of Epon 815C resin sili-cate nanocomposites are listed in Table 3. There is an 8%drop in modulus at 2% silicate loading from that of thepristine polymer, followed by a 12% and 18% increase inmodulus for the 4% and 6% samples, respectively. Thestrength exhibits a 9–16% monotonic decrease with silicateloading. A similar effect has been observed by others, and isgenerally attributed to poor dispersion of the nanoparti-cles, since the strength is very sensitive to the presence offlaws in the sample, and the inhomogeneous nanoparticledispersion is considered as a flaw. In addition, variationsin crosslink topology could lead to molecular scale defects,

Table 3Mechanical test results of 815 resin silicate

Sample(%)

Modulus(Gpa)

Standard deviationmodulus

Strength(Mpa)

Standard deviationstrength

0 3.5 0.4 228.6 14.62 3.2 0.16 196 104 4.0 0.14 194 216 4.3 0.35 190 12.1

Fig. 8. SEM of flexural

such as dangling chains, which could also result in strengthdecreases [7,18].

The flexural properties for the fiber-reinforced compos-ites are shown in Table 4. The flexural modulus increasesby 31% for the 2% sample, followed by a slight decreaseand a slight increase for the 4% and 6% samples, respec-tively. The 2% sample exhibits a 24% increase in strengthwhile the higher silicate loadings results in decrease instrength. The enhancement and strength for the 2% samplemay be due to a more homogeneous dispersion of the orga-nophilic silicate in the nanocomposite, coupled with itsmuch lower viscosity, relative to the higher silicate load-ings. This lower viscosity may lead to better wetting ofthe fibers and dispersion of the nanoparticles. Scanningelectron microscopy (SEM) of the failure surface inFig. 8 indicates that there is some fragmentation of fiberbundles, evidence of extensive fiber debonding, delamina-tion and extensive matrix cracking. The flat fiber fracturesurface suggests the dominance of brittle-type fracture

test failure surfaces.

Table 5Interlaminar shear test results for fiber-reinforced nanocomposites

Specimen Shear strength(lbf/in.2)

Standarddeviation

% Variation

Neat 5674 973 N/A2% 5057 212 �114% 4652 266 �186% 4885 255 �14

D. Dean et al. / Composites Science and Technology 66 (2006) 2135–2142 2141

behavior at higher magnification. These features were typ-ical of all the samples. It is also noted that the fact that nonanoparticles are observed is suggestive that no filtering ofthe nanoparticles by the fibers has occurred.

3.5. Short beam interlaminar shear test

Table 5 summarizes the data from the interlaminarshear tests for the fiber-reinforced samples. A decrease of11%, 18% and 14% is observed for the 2%, 4% and 6% sam-ples, respectively. We reported an increase in interlaminarshear strength for glass reinforced epoxy and vinyl ester/sil-icate nanocomposites and attributed it to enhancedmatrix–fiber adhesion [31,32]. Rice et al. reported marginalincreases in the room temperature interlaminar shearstrength for a carbon fiber composite in which the matrixwas an epoxy resin dispersed with carbon nanofibers. Theysuggested that better nanofiber–matrix adhesion wouldresult in better property enhancements [39]. The decreasesthat we observe are presumably due to the dissimilaritybetween the silicate nanoparticle and the reinforcing car-bon fiber. A follow up study using chemically modified car-bon nanofibers will be conducted to test both hypotheses.

4. Conclusion

The processing-property relationships of multiscalefiber-reinforced epoxy nanocomposites with a hierarchalstructure ranging from nanoscale particles to micron sizefibers have been studied. The effect of silicate loading onthe flow and isothermal cure behavior of the epoxy resinwas also studied. The addition of silicate results in anincrease in resin viscosity by a factor of 10, although it isstill within the range of processability by VARIM. A slightdecrease in resin gel times is also observed. Fiber-rein-forced systems were prepared using a combination of vac-uum assisted resin infusion and compression molding. Theresulting laminates exhibited a maximum modulus andstrength enhancement of 31%, and 19%, respectively, forthe system containing 2% silicate, while these propertiesdecreased for the higher silicate loadings.

The enhancement in modulus and strength for the 2%sample may be due to a more homogeneous dispersion ofthe organophilic silicate in the resin, coupled with its muchlower viscosity. This lower viscosity presumably leads tobetter wetting of the fibers, and reduces the possibility offiltering of the nanoparticles by the fibers. While fiber dom-inated properties of fiber-reinforced composites can benefit

from the incorporation of nanoparticles, the resin domi-nated properties may benefit the most. Thus, future studieswill investigate the effect of the hierarchical structures pres-ent in these multiscale composites on properties such astoughness, thermal expansion coefficient and thermalstability.

Acknowledgment

This work was funded by the Air Force Office of Scien-tific Research Grant No. F49620-00-1-0348.

References

[1] Mallick PK. Fiber reinforced composites. 2nd ed. New York (NY):Marcel Dekker; 1993.

[2] Kornmann X, Berglund LA, Sterte J, Giannelis EP. Polym Eng Sci1998;38:1351–8.

[3] Kornmann X, Lindberg H, Berglund LA. Polymer 2001;42:4493–9.[4] Lan T, Kaviratna PD, Pinnavaia TJ. J Phys Chem Solids 1996;57(6–

):1005–10.[5] Messersmith PB, Giannelis EP. Chem Mater 1994;6:1719–25.[6] Giannelis EP. Adv Mater 1996;8:29–35.[7] Le Baron P, Wang Z, Pinnavai T. Appl Silic Sci 1999;15:11–29.[8] Okada A, Kojima Y, Kawasumi M, Fukushima Y, Kurauchi T,

Kamigaito O. J Mater Res 1993;8:1179.[9] Kornman X, Lindberg H, Berglund LA, Sterte J, Giannelis EP.

Polym Sci Eng 1998;38:1351–8.[10] Xu W-B, Bao S-P, Shen S-J, Hang G-P, He PS. J Appl Polym Sci

2003;88:2932–41.[11] Xu W-B, Bao S-P, He PS. J Appl Polym Sci 2002;84:842–9.[12] Tolle TB, Anderson DP. J Appl Polym Sci 2004;91:89–100.[13] Nigam V, Setua DK, Mathur GN, Kar K. J Appl Polym Sci

2004;93:2201–10.[14] Abdalla M, Dean D, Campbell S. Polymer 2002;43:5887–93.[15] Abdalla MO, Ganguli S, Abdalla MA, Dean D, Campbell S. High

Perform Polym 2005;17:239–50.[16] Wang J, Lan T, Pinnavaia JT. Chem Mater 1996;8(9):2200–24.[17] Lesser A, Zerda AS. Mechanical properties of intercalated epoxy-

silicate nanocomposites. Mater Resour Symp Proc 2000.[18] Ganguli S, Dean D, Jordan K, Price G, Vaia R. Polymer

2003;44:1315–9.[19] Vaia R. APMTIAC Quart 2002;6(1).[20] Tolle TB, Anderson A. Compos Sci Technol 2002;62:1033–41.[21] Theodore T, Dean D, Obore A, Nyairo E. Polym Prepr

2004;45(1):863.[22] Zhang K, Wang L, Wang F, Wang G, Li Z. J Appl Polym Sci

2004;91:2649–52.[23] Halley PJ, Mackay ME. Polym Eng Sci 1996;36:5.[24] Ganguli S, Dean D, Jordan J, Price G, Vaia R. Polymer

2003;44:6901–11.[25] Daniel J. Compos Sci Technol 2002;63(11):1607–16.[26] Salahuddian N, Moet A, Hiltner A, Baer E. Eur Polym J

2002;38:1477–82.[27] Chen JS, Ober CK, Zhang Y, Ulrich W, Giannelis E. Polymer

2002;43:4895–904.[28] Chin TIJ, Ho-Cheol K, Russell TP, Wang J. Polymer

2001;41:5947–52.[29] Rice BP, Chen C, Cloos L, Curliss D. SAMPE J 2001;37:7–9.[30] Chen C, Curliss D. SAMPE J 2001;37:11–8.[31] Haque A, Shamsuzzoha M, Hussain F, Dean D. J Compos Mater

2003;37(20):1821–38.[32] Hussain F, Dean D, Haque A, Shamsuzzoha M. J Adv Mater

2005;37(1):16–25.[33] Krishnamoorti R, Giannelis EP. Macromolecules 1997;30:4097.

2142 D. Dean et al. / Composites Science and Technology 66 (2006) 2135–2142

[34] Krishnamoorti R, Ren J, Silver AS. J Chem Phys 2001;108:7175.[35] Krishnamoorti R, Vaia RA, Giannelis EP. Chem Mater 1996;8:1728.[36] Krishnamoorti R, Giannelis EP. Langmuir 2001;17:1448.[37] Hsieh AJ, Moy P, Beyer FL, Madison P, Napadensky E, Ren J, et al.

Polym Eng Sci 2004;44(5):825–37.

[38] Dean D, Walker R, Theodore M, Hampton E, Nyairo E. Polymer2005;46(9):3014–21.

[39] Rice BP, Gibson T, Lafbi K. Development of multifunctionaladvanced composites using a VGNF enhanced matrix. Int SAMPETech Conf 2004:2223–35.