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Anisotropic Shape-Memory Elastomeric Composites: Fabrication and Testing

Erika D. Rodriguez , Derek C. Weed , Patrick T. Mather *

Anisotropic shape-memory elastomeric composites (A-SMEC) are fabricated by infi ltrating an aligned electrospun fi brous web with an elastomer, and curing. Mechanical testing proves strong anisotropy in Young’s modulus ( E Y ), strain-to-failure ( ε f ), and yield stress ( σ y ) observed with averages of E Y = 48.4 and 4.0 MPa, ε f = 198 and 351%, and σ y = 4.6 and 0.9 MPa for longitudinal and transverse loading directions, with respect to the fi brous web’s orientation. Shape-memory (SM) characterization reveals modest anisotropy in the shape fi xing ratios ( R f ), yet seemingly isotropic shape recovery ( R r ) nearing completeness. Such soft, anisotropic materials may be used in laminated composite elastomers.

1. Introduction

Shape-memory polymers (SMPs) have a unique ability to be temporarily fi xed in a non-equilibrium shape with internal stresses stored by immobilizing constituent net-work chains of the SMP through either vitrifi cation or crystallization upon cooling below the glass transition temperature ( T g ) or melting temperature ( T m ), respec-tively. Subsequent external stimulation remobilizes the network chains, whether electrically, photothermally, or thermally, to restore its original geometry. [ 1–3 ] While extensive studies have examined such shape-memory (SM) phenomena, comparatively little attention has been given to the construction of SMP systems that are soft in nature and exhibit anisotropic mechanical behavior. [ 4 , 5 ] Although extensive research and development has been conducted for rigid anisotropic materials, forming the basis for

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Dr. E. D. RodriguezMechanical and Aerospace Engineering, Syracuse Biomaterials Institute, Syracuse University, Syracuse, NY 13244, USA D. C. Weed, Prof. P. T. MatherBiomedical and Chemical Engineering, Syracuse Biomaterials Institute, Syracuse University, Syracuse, NY 13244, USAE-mail: [email protected]

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Macromol. Chem. Phys. 2013, DOI: 10.1002/macp.201300086© 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

modern laminated composites used in the automotive and aerospace industries, we are motivated to develop soft ani-sotropic composite materials and SMPs that mimic some characteristics of biological materials. [ 6 ] If successful, such anisotropic elastomers could form a basis for soft lami-nated structures analogous to their rigid counterparts but with benefi ts of softness required for some applications. [ 7 ]

One example of a biological material we are motivated to mimic is the skin membrane of a bat wing, as it has the distinctive properties of an anisotropic fi brous network. The bat’s wing is stiff in the airfoil (chord) direction, yet compliant and collapsible in the span direction for animal perching. [ 8 ] To accommodate these requirements, the skin membrane is both elastic and anisotropic, with mechan-ical properties that are dictated by a composite structure bearing internal, closely spaced fi bers. [ 8 ] In particular, the mechanical anisotropy is due to the nature of the bat’s skin membrane itself, involving a complex arrange-ment of both collagen and elastin fi ber bundles, that are mechanically distinct materials, and their respective ori-entations that provide specifi c functions to the skin. [ 8 ] More specifi cally, the collagen fi bers are used for the membrane’s structural stability for aerodynamic fl ight while the elastin serves as the elastic component for wing retractability. It has been deduced that these constituent

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fi bers also play a critical role in the transmittance of aerodynamic forces from the surface of the wing to the upper limb skeleton. [ 9 ] A synthetic elastomer with sim-ilar mechanical anisotropy would be useful for micro air vehicle systems featuring fl apping wings, an area of recent interest to aerodynamics researchers funded by the U.S. Air Force. [ 10 ] Such biomimicry would improve the adaptability, agility, and speed in unmanned aerody-namic vehicles. [ 10 ]

As mentioned above, several studies of SMPs have been reported. Osada and co-workers constructed a SM hydrogel where water-swollen gels were prepared by co-polymer-izing n -stearyl acrylate, acrylic acid, and a crosslinker. [ 11 , 12 ] Below 25 ° C, the water-swollen hydrogel is rigid, due to substantial crystallization of pendant n -stearyl acrylate. This crystallization was further found to enable shape fi xing of the hydrogels, where the SM was afforded by crosslinking. Above 50 ° C, the hydrogel behaves like a soft, compliant gel capable of relatively large ( > 100%), reversible deformations. Thus, stretching the compliant hydrogel at elevated temperature, followed by cooling to crystallize the side chains allowed for good shape fi xing, while subsequent heating to melt the side changes cause shape recovery to the equilibrium shape. Relative to the mentioned goal, softness was possible through hydration, though no mechanical anisotropy was reported.

Soft SMPs are also possible, with no need for hydra-tion, by crosslinking low- T g ( T g < room temperature (RT)) polymers to an elastomeric state and further incorpo-rating some mechanism of strain fi xing other than vit-rifi cation. As a good example, covalent cross-linking of semicrystalline trans -polyisoprene (TIP) yielded an SM network featuring an RT tensile modulus of 100 MPa, suffi cient crystallinity for shape fi xing, and superelastic mechanical properties with good shape recovery above the melting transition, T m = 67 ° C. [ 13 ] In other cases from our own group, main-chain liquid crystalline (smectic) elastomers were found to feature good SM, while POSS-functionalized and crosslinked poly(hydroxyl alkanoate) prepared by bacterial synthesis offered yet another soft SM approach. [ 4 , 5 , 14 ] In both of these cases, the softness of the SMPs was exploited for the purpose of reversible embossing of imprinted, temporary topographical fea-tures, while macroscopic anisotropy was not investigated.

Some progress has been made in the preparation of anisotropic elastomers, though without SM. For example, Li et al. [ 15 ] constructed a highly anisotropic thermoplastic elastomer composite by melt-blending poly(styrene- block -butadiene- co -butylene- block -styrene) (SBBS) and a liquid crystal polymer (LCP) and injection molding the blend. The LCP uniquely orients in fi ber form along the injection direction in the SBBS matrix where the micro-phase-separated polystyrene (PS) phase in the matrix conforms along the injection direction yielding high

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anisotropy. This composite showed high stiffness in the melt injection direction, yet high compliance in the transverse direction. Indeed, the modulus in the injec-tion direction was 47-fold greater than in the transverse direction owing to the anisotropic arrangement. Elonga-tion-at-break, ε Y , was reported to be 56-fold smaller in the injection direction when compared to the transverse direc-tion. Despite these desirable properties, such an approach is characterized by high sensitivity of morphology and macro scopic properties to molding conditions.

Cohen et al. demonstrated another unique anisotropic elastomer utilizing orientation of a triblock copolymer. Here, the lamellar phase of a triblock polystryene-poly-butadiene-polystrene (PS/PB/PS) copolymer was oriented in fi lm form via roll-casting. [ 16 ] (Numerous other reports of roll-casting triblock copolymers have appeared and the interested readers are referred to those papers without further discussion. [ 17–22 ] ) The roll casting process yielded highly oriented lamellar block copolymer fi lms featuring strongly anisotropic mechanical properties. Samples with different lamellar orientations were uniaxially strained to understand the deformation mechanism and show the difference in mechanical properties. When uniaxially loaded in directions parallel, perpendicular, and diagonal to the lamellar structure, the measured Young’s moduli were 180 ± 10, 65 ± 3, and 43 ± 4 MPa, respectively. [ 16 ] The higher moduli were evident in samples that had lamellae in the parallel orientation and, interestingly, lower in the diagonal orientation than the perpendicular direction. Loading parallel to the lamellar structure led to a plastic-to-rubber transition, while perpendicular loading led to layer folding similar to that of chevron morphology where kink boundaries were generated at higher strains. [ 16 ]

Still another soft and anisotropic material, liquid single crystal elastomers (LSCE) offer a unique approach to ani-sotropy. Finkelmann’s group has published extensively on nematic and smectic LSCEs. [ 23–25 ] One study in particular revealed the mechanical anisotropy for the LSCE smectic A elastomer, which was examined for its temperature dependence and associated with network chain entropy elasticity, specifi cally perpendicular to the normal of the smectic layers. [ 25 ] In the transverse direction, the mechanical response was concluded to be controlled by an enthalpic elastic response wherein the smectic layers contributed a modulus of two orders of magnitude greater than that of the soft spacers. [ 25 ] The modulus has also shown to decrease with increasing temperature in the transverse direction.

Our group recently introduced a composite approach involving electrospinning of thermoplastic fi bers to enable the creation of soft SMPs. In this approach, an isotropic shape-memory elastomeric composite (SMEC) was constructed through a three-step fabrication pro-cess using an electrospun nonwoven poly( ε -caprolactone)

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(PCL) semicrystalline fi ber mat with Sylgard-184 silicone (hereafter “Sylgard”) infi ltrating the mat, serving as the elastomeric matrix. [ 26 ] In such a composite, the separate roles of strain fi xing and strain recovery were distributed between the two interpenetrating phases of the com-posite: strain fi xing with the thermoplastic fi ber phase (PCL) and recovery by the elastomeric silicone phase. In this case, the SMEC was macroscopically isotropic and we reasoned that orientation of the fi ber phase could lead to the desired anisotropy in a straightforward manner.

Thus, for the present report, we aimed to construct an anisotropic shape-memory elastomeric composite (A-SMEC) by electrospinning aligned fi bers of poly(vinyl acetate) (PVAc), serving as the composite reinforcement and SM fi xing contribution to the system, and infi ltrating with Sylgard, to serve as the elastomeric matrix while also assisting in the composites shape recovery capability. We selected PVAc over previously studied PCL for the fi ber phase due to the much higher stiffness of glassy PVAc that was expected to yield high mechanical anisotropic in fi ber-aligned composites. We will show that anisotropy in


Figure 1 . Three-step fabrication of anisotropic shape-memory elasorientation angles ranging from 0 ° (longitudinal direction) to 90 ° (


Macromol. Chem. Phys. 2013, D© 2013 WILEY-VCH Verlag Gm

Young’s modulus ( E Y ), strain-to-failure ( ε f ), and yield stress ( σ y ) is achieved. We further show the SM effect is tailored as a function of fi ber orientation angle, χ , where shape fi xing depended on the orientation angle of the electro-spun microfi bers incorporated in the matrix.

2. Experimental Section

2.1. Materials and Sample Fabrication

The newly prepared A-SMEC materials are composed of two commercially available polymers: crosslinked poly(dimethylsiloxane) (Sylgard-184 from Dow Corning) and PVAc ( M w = 249 kg mol − 1 from Aldrich). Fabrication proceeded using a three-step process adapted from prior work and shown in Figure 1 . [ 26 ] PVAc was electrospun from a relatively concentrated (16.7 wt/vol%) mixed solvent solution in MeOH:DMF (volumetric ratio = 7:3). In particular, 2.00 g of PVAc was dissolved in 10.0 mL of the solvent mixture and stirred continuously until a transparent, clear, and homogenous solution was achieved. Using a custom electrospinning apparatus, fi bers from a charged needle

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tomeric composites (A-SMEC). Samples were punched out with fi ber transverse direction).


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( + 13.5 kV) were electrospun onto a rotating drum maintained at slightly negative potential ( − 0.5 kV) using a large angular velocity of 3000 rpm. The drum diameter was 5.08 cm, the tip-collector distance was 7.0 cm, and the fl ow rate was 0.7 mL h − 1 . Electrospinning ran for 10.75 h to yield robust, “handleable” fi brous webs, ca. 0.13 mm in thickness, that were easily removed from the rotating drum. A high degree of fi ber orientation for all samples was afforded by incorporation of a charged ( + 13.5 kV) square steel plate (length ( l ) = 15.0 cm and width ( w ) = 12.5 cm) mounted to the syringe hub with the plate normal oriented parallel to the fl ow direction (Figure 1 ). Here, the idea is to decrease the divergence of electrical fi eld lines between the drum and the plate that are characteristic of the point/plane or point/cylinder arrangement. Qualitatively, we observed dramatically higher fi ber orientation for such confi guration, all other conditions remaining the same, and so it was adopted for the preparation of all samples. The average fi ber diameter for the fi brous webs thus prepared was measured by analyzing scanning electron microscopy (SEM) micrographs in ImageJ (Version 1.45s) software where a total of 90 measurements were used to obtain the average.

To prepare A-SMEC samples from the oriented fi brous webs, the webs were infi ltrated by a two-part mixture of Sylgard (base: curing agent ratio = 10:1), assisted by vacuum application (ca. 760 mm Hg) for 20 min to ensure complete infi ltration of Sylgard into the PVAc fi brous web. After carefully removing the excess Sylgard from the fi ber mat with a spatula, the Sylgard was cured at 30 ° C for 48 h (Figure 1 ). Unexpectedly, the infi ltration process dilated the samples to a thickness of ca. 0.31 mm (compared with 0.13 mm for the fi brous webs), effectively decreasing the fi ber weight- and volume-fractions as we will show. Dogbone-shaped specimens were then punched at various angles from the A-SMEC sheet so as to yield a range of fi ber orientation angles, χ , relative to the loading direction: χ = 0 ° (longitudinal), 22.5 ° , 45 ° , 67.5 ° , and 90 ° (transverse), as shown in Figure 1 . For clarity, we will refer to A-SMEC samples with varying orientation angles, χ , listed above using the nomenclature, A-SMEC- χ . For example, the specimen with fi ber orientation angle inclined 45 ° from the loading direction will be referred to as A-SMEC-45.

2.2. Microstructural Characterization

The as-processed fi brous webs were analyzed using SEM and image analysis software, ImageJ to, obtain the average fi ber diameter. 2D fast Fourier transform (FFT) analysis was applied to SEM images of the same fi brous webs to obtain the degree of fi ber orientation, expressed as the angular breadth from an azimuthal sweep of the 2D-FFT streak using the Gywddion soft-ware’s Hahn model (Figure S1, Supporting Information). Utilizing an ImageJ plugin, “Oval_Profi le” radial sums of the magnitude/intensity were taken and the data were analyzed to study degree of fi ber alignment. Degrees of alignment were then determined using the full width at half max (FWHM) approach. The weight fraction of fi bers in the A-SMECs was then measured in two ways: i) gravimetrically using weights before and after Sylgard infi ltration, and ii) calorimetrically using the heat capacity step at T g for the PVAc, where X f = �CpASMEC

/�CpPV Ac . From these

quantities and known densities of each constituent, the associ-ated volume fractions were estimated, as well.

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2.3. Dynamic Mechanical Analysis

Dynamic mechanical properties for all samples were studied using a TA Q800 dynamic mechanical analyzer (DMA). Each sample was cut at the specifi ed angles mentioned in Section 2.1 from an A-SMEC sheet to obtain a rectangular sample with average dimensions of l = 7.46 ± 1.04 mm, w = 2.47 ± 0.33 mm, and t = 0.38 ± 0.04 mm. We note that the uncertainties above (one standard deviation) indicate dispersion in dimensions among all samples, not uncertainty in the values for individual samples, which were measured with 0.01 mm precision. For each sample, the linear viscoelastic properties were measured in a tensile mode at a single frequency and during temperature sweep, revealing thermomechanical transitions. Specifi cally, oscillatory tensile deformation was used with a small displace-ment amplitude of 15 μ m (tensile strain < 0.4%) with a 1 Hz frequency, a preload of 0.001 N, and force track (ratio of static to dynamic forces) of 110%. First, samples were cooled to 0 ° C and held at that temperature for 5 min. The temperature was then ramped to 100 ° C at 3 ° C min − 1 . This process was repeated, with tensile storage modulus ( E ′ ) values and loss tangent (tan δ ) values measured during the second heating. This procedure was adopted for all the A-SMEC samples, revealing depend-ence on χ . Neat Sylgard was also tested using the same dis-placement, frequency, and force track parameters stated above, though with a distinct thermal history. In particular, Sylgard samples were fi rst heated to 100 ° C at 3 ° C min − 1 , held at this temperature for 1 min, cooled to 0 ° C at 3 ° C min − 1 , held isother-mally for 1 min, then cooled further to − 150 ° C held for 5 min. Finally, the samples were heated to 120 ° C at 3 ° C min − 1 . The process was repeated and property measurements were made during the second heating. This particular thermal history was applied to Sylgard samples to allow observation of the T g .

2.4. Mechanical Testing

Mechanical properties of both the PVAc fi brous web and A-SMEC state were studied using a Linkam TST 350 tensile stress testing system, providing E Y , ε f , and σ y as functions of fi ber orientation angle, χ . For this purpose, each sample was punched into a dog-bone geometry using an ASTM Standard D638-03 Type IV (scaled down by a factor of 4, TestResources, Inc., Shakopee, MN). The dogbones were punched at angles ( χ ) of 0 ° (longitudinal), 22.5 ° , 45 ° , 67.5 ° , and 90 ° (transverse) relative to the nominal fi ber ori-entation direction and elongated at constant velocity according to ASTM standard. The average thickness of the fi brous web dog-bone samples was 0.13 ± 0.03 mm, and the average thickness of A-SMEC dogbone samples was 0.31 ± 0.03 mm. These thicknesses were used to construct the stress–strain curves (engineering stress versus engineering strain) from the raw load-displacement data.

Samples were held isothermally during each test at RT, the strain rate was 20% min − 1 (20.8 μ m s − 1 ), and a sensitive load cell with 20 N capacity was used. Plots of engineering stress vs. engineering strain were then constructed from the raw data. The apparent tensile moduli for the fi brous web specimens were calculated from the slope of the elastic region of the stress–strain curve, for which the cross-sectional area of each porous samples was used in estimating the macroscopic stress.


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Figure 2 . SEM micrographs of: a) PVAc fi bers, b) topography of A-SMEC, c) A-SMEC with longitudinal fi ber orientation, and d) composite with transverse fi ber orientation.

In this manner, only the apparent Young’s modulus ( E A ) could be determined. While ε f measurements were directly observable for A-SMEC samples using the point of sample fracture, PVAc fi brous web specimens exhibited a broad strain range over which failure occurred. For such cases, ε f values were extrapolated from the infl ection point of the ductile failure region of the stress–strain curve to the strain axis (zero stress), reporting that strain value as the ε f . Young’s moduli values for A-SMEC samples were obtained by fi tting a linear trendline in the elastic region of the stress–strain curve, while the ε f data are simply the strains at the point of fracture, where the tensile stress dropped to zero. The σ y for both the fi brous web and the A-SMEC samples was calculated using a 2% offset method whereby the elastic region of the stress–strain curve was fi tted by a linear trendline, horizontally offset in strain by 2%, and plotted to intersect the stress–strain data at the yield stress.

2.5. Shape-Memory Analysis

The SM characteristics were quantifi ed for all A-SMEC samples with the fi ve fi ber orientation angles mentioned previously. A well studied four-step SM cycle referred to as the one-way shape-memory (1WSM) method was performed on all A-SMECs. [ 1 , 2 , 27 ] Samples were tested using a TA Q800 DMA with a preload force of 0.001 N. The samples were heated to 80 ° C and held isother-mally for 5 min. With the strain set to 0% prior to stretching the sample, the force was ramped at 0.5 N min − 1 until 100% strain was achieved, and this force was held for 2 min ( Step 1 ). The tem-perature was then decreased from 80 ° C to 0 ° C at 3 ° C min − 1 and held at 0 ° C for 5 min for fi xing of the temporary strain ( Step 2 ). The force was then released to 0.001 N at 0.5 N min − 1 and held for 2 min to observe elastic recovery and sample fi xing ( Step 3 ). Finally, the temperature was increased from 0 ° C to 80 ° C at 3 ° C min − 1 and then held isothermally for 5 min to allow com-pletion of shape recovery ( Step 4 ). This four-step process was repeated two more times on the same sample to observe SM reproducibility. The fi xing ratio ( R f ) and recovery ratio ( R r ) were calculated using the following equations:

Rf (%) =εu (N)εm (N)

. 100 (1)

and:

Rr (%) =εu (N) − εp (N)

εu (N) − εp (N − 1). 100 (2)

where ε m , ε u , ε p and N are the strain before unloading, the strain after unloading, permanent (unrecovered) strain after shape recovery, and the SM cycle number, respectively. R f and R r aver-aged from the second and third cycles are reported.




3. Results and Discussion

3.1. Sample Fabrication

A three-step process (Figure 1 ) was conducted to fabri-cate A-SMEC materials whose anisotropy derives from the constituent oriented PVAc fi brous webs e-spun using the parameters specifi ed in the Section 2. Such fi brous webs were infi ltrated with Sylgard-184 and “cured” (crosslinked) to yield A-SMEC materials. Microstructurally, the fi brous webs and A-SMEC samples exhibited clear anisotropy. Figure 2 shows representative SEM micrographs of: a) aligned fi bers, b) topography of the A-SMEC, and cross-sections of the A-SMEC when the fi bers are aligned in the longitudinal (c) and transverse (d) directions. The average PVAc fi ber diameters were 0.79 ± 0.20 μ m and 0.77 ± 0.18 μ m for two separately prepared fi brous webs using 90 meas-urements on each web. SEM images of the fi brous webs were analyzed with a 2D FFT to determine the degree of alignment for the fi bers. The two fi brous webs proved to have signifi cant alignment, which amounted to FWHM of the 2D FFT spectra of 30 ° and 34 ° . Further details asso-ciated with such alignment measurements can be found in the Figure 1 in Supporting Information, which contains fi brous web SEM micrographs, 2D FFT images, and the resultant orientation plots.

A-SMEC samples showed a smooth layer of Sylgard on the surface, which differs from the report of Luo and Mather on isotropic SMEC, where material’s surfaces were reported to be somewhat porous. [ 26 ] This observed differ-ence is attributed to the comparatively dense arrangement

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Figure 4 . Scatter plot showing tensile storage modulus at 25 ° C ( � ), 50 ° C ( � ), and 80 ° C ( � ) as functions of fi ber orientation angle, along with Sylgard as a dashed reference line.

Figure 3 . Tensile storage modulus ( E ′ ) (a) and tan δ (b) versus tem-perature for A-SMEC samples as a function of fi ber orientation angles: 0 ° (i), 22.5 ° (ii), 45 ° (iii), 67.5 ° (iv), 90 ° (v) where Sylgard (vi) is plotted as a reference.

of oriented fi bers where only small voids were present between them. Nevertheless, SEM images of the cross sec-tions clearly showed complete Sylgard infi ltration along the width and length of the A-SMEC while also showing the fi bers in the A-SMEC were preserved in the Sylgard matrix (Figure 2 ). Using the mass of the fi brous webs, mass of the A-SMEC, and the densities of the PVAc ( ρ = 0.934 g mL − 1 ) and Sylgard ( ρ = 0.965 g mL − 1 ), the overall mass and volume fractions of the A-SMEC were estimated to be 17.1% and 17.6%, respectively. For comparison, the fi ber weight fraction determined from heat capacity steps at T g was calculated to be 19.7% where the Δ C p PVAc-fi bers = 0.502 J g − 1 ° C − 1 and Δ C p A-SMEC = 0.099 J g − 1 ° C − 1 . This yields an estimated volume fraction of 20.2%. The higher values measured using this DSC method may be attributed to localized testing compared with testing the entire fi brous webs as the method based on mass does. The DSC traces of the PVAc pellets, PVAc fi brous web, and A-SMECs are shown in Figure S2 in the Supporting Information.

3.2. Dynamic Mechanical Analysis of Composites

To obtain angle-dependent thermo-mechanical proper-ties using DMA, all A-SMEC samples were fi rst cut into a rectangular geometry whose long axis was oriented at a prescribed angle relative to the oriented fi ber direction (Figure 1 ). This testing yielded the temperature depend-ence of linear viscoelastic properties (tensile storage modulus and loss tangent) of each A-SMEC sample, as shown in Figure 3 . For such A-SMEC samples, the highest E' ( ≈ 100 MPa) at RT was obtained for fi bers oriented lon-gitudinally (A-SMEC-0) whereas the lowest values (below 10 MPa) at RT were obtained for samples with fi bers ori-ented in the transverse direction (A-SMEC-90). The E' then decreased signifi cantly with increasing temperature, par-ticularly as the PVAc T g was exceeded and regardless of fi ber orientation; observations of this phenomenon were seen to start at T = 55 ° C (Figure 3 ). The T g inferred from DMA data for all samples tested ranged from 50 to 55 ° C, values ≈ 10 ° C higher than those from DSC analysis (Figure S2, Supporting Information), which is commonly observed. We note that the Sylgard matrix remained in its rubbery state through the entire temperature sweep (vi in Figure 3 ) as expected, given that the T g of Sylgard is − 115 ° C. [ 26 ]

From the data of Figure 3 , the orientation depend-ence of E ' could be extracted and such trends are plotted in Figure 4 for several temperatures. Not surprisingly, we observed that the E ' was higher for temperatures below the PVAc T g for all orientations. Further, signifi cant mechanical anisotropy, refl ected in the slope of modulus versus orientation angle, was observed only below T g . (Data from this fi gure are presented in the Supporting Information). Importantly, the moduli for A-SMEC sam-ples are lower than that of neat Sylgard at 80 ° C, a fi nding



consistent with a prior report on isotropic SMEC. [ 26 ] Above T g , the A-SMEC samples are elastomers with fl uid inclu-sions (the PVAc fi bers), thus the samples features a lower elastic modulus than dense Sylgard.





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Figure 5 . Representative stress–strain curves for: a) PVAc fi ber and b) A-SMEC dogbone samples as a function of fi ber orientation angles of: 0 ° (i), 22.5 ° (ii), 45 ° (iii), 67.5 ° (iv), 90 ° (v) where Sylgard (vi) is plotted as a reference.

Figure 6 . Tensile modulus as a function of fi ber orientation angle at RT for PVAc fi brous web and A-SMEC samples with Sylgard as a reference: fi ber apparent modulus ( � ) A-SMEC Young's modulus ( � ), and composite storage modulus ( � ). Sylgard storage modulus is shown as a dotted reference line.

3.3. Mechanical Testing Characterization

The fi brous web and A-SMEC dogbone samples were sub-jected to tensile testing wherein punched dogbones were strained to the point of failure, enabling determination of the E A (fi brous web samples), E Y , ε f , and σ y as functions of fi ber orientation. As Figure 5 shows, the fi brous webs and A-SMEC composites exhibited mechanical properties that were robust, overall, and that depended strongly on orienta-tion angle relative to the loading direction. In particular, the fi brous web and A-SMEC samples alike featured dramatic anisotropy in all mechanical properties measured ( E A , E ′ , E Y , ε f , and σ y ). In comparing to neat Sylgard (Figure 5 ), the fi brous web and A-SMEC materials were stiffer and quali-tatively distinct in character, with appearance of sample yielding at strains well below the ε f values. The complete set of stress–strain curves for each fi brous web and A-SMEC sample tested is provided as Figure S3 and S4 in the Sup-porting Information. The tensile-energy-to-break (J cm − 3 ) values, determined from integration of the stress–strain curves, were comparable among fi brous web, A-SMEC, and Sylgard samples, with values spanning 5–15 J cm − 3 and no clear trend with orientation angle being evident. The plots of tensile-energy-to-break versus orientation angle, χ , are provided as Figure S5 in the Supporting Information.




The orientation dependences of E ′ (from DMA), E A , and E Y (from the stress–strain curves at RT for both the fi brous webs and A-SMEC specimens) are plotted in Figure 6 , revealing that the fi brous web samples feature a higher modulus than the A-SMEC samples at any given orienta-tion. Further, it is observed that the small strain E ′ was higher than E Y , as is commonly reported. To explain our observation that A-SMEC samples have lower tensile mod-ulus than fi brous webs of the same orientation angle, we consider our observation that the process of infi ltrating the webs featured the unintended consequence of dila-tion, the fi brous web thickness increasing approximately threefold, ostensibly increasing the distance between fi bers and decreasing weight and volume fraction of the fi bers. Given that PVAc is approximately 1000-fold stiffer than Sylgard, a decrease in volume fraction will directly decrease the composite modulus. We observed that A-SMEC-0 (longitudinal orientation) featured E Y = 48.4 ± 9.7 MPa. As χ increased, E Y decreased dramatically where A-SMEC-90 features E Y = 4.0 ± 0.4 MPa ( E A , E Y , and ε f data for both the fi brous web and A-SMEC samples are shown in Table S2, S3, S4, and S5, respectively in the Supporting Information). In the same manner as E' and E Y , E A also show strongly decreasing values with increasing χ .

The strong mechanical anisotropy evidenced for A-SMEC samples in Figure 6 is not observed above the fi ber T g , as anticipated due to a loss of the fi ber’s rein-forcement capability. This is seen clearly in Figure 7 , where E Y ( χ ) is shown for both 25 ° C and 80 ° C. Inter-esting, we observed that the model for angular depend-ence of elastic modulus for orthotropic lamina com-monly applied to structural composites works well for our data. [ 28 ] Accordingly, the solid line in Figure 7 is the

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Figure 7 . Young’s modulus as a function of fi ber orientation angle for A-SMEC samples at RT (symbols) along with the fi t from ortho-tropic laminate theory (solid line). The dashed line represents a trendline in Young’s modulus for bat wing membranes re-plotted from Swartz et al. [ 8 ]

Figure 9 . Yield stress versus orientation angle for PVAc fi brous web ( � ) and A-SMEC ( � ) samples.

Figure 8 . Strain-to-failure for PVAc fi brous web ( � ) and A-SMEC ( � ) samples, plotting average strain-to-failure as a function of fi ber orientation angle with average (solid line) and standard error (dashed line) Sylgard reference lines.

best fi t to our data using the orthotropic lamina formula Equation 3 :

1

Ex= cos4 x

ET+ sin4 x

EL+ 1

4(1

GLT

− 2vLTEL

) sin2 2x (3)

Where E x is the angle-dependent elastic modulus, E L is the longitudinal modulus (measured), E T is the trans-verse modulus (measured), G LT is the shear modulus of the matrix (unknown), and ν LT is Poisson’s ratio of the matrix (unknown). For fi tting purposes, ν LT was held fi xed at 0.3, and G LT determined by regression, yielding a reasonable value of 3.1 MPa. We note that the shapes of curves for this model are relatively insensitive to ν LT .

Given our motivation to achieve anisotropic elastomers that mimic the properties of bat wings, we compared our results with E Y ( χ ) data reported for bat wings for a number of bat species by Swartz et al. [ 8 ] Those results are shown as the dashed line in Figure 7 , revealing surpris-ingly similar values as the A-SMEC system for E Y , though slightly less mechanical anisotropy. It is evident, then, that A-SMEC mimics well the anisotropic mechanical properties of bat wings.

The angular dependence of ε f was examined and compared among the fi brous web and A-SMEC systems. As shown, in Figure 8 , ε f for the fi brous web samples were lower than those of the A-SMEC samples at any given χ , a fi nding attributed to the elasticity of Sylgard in the A-SMEC state. Strain-to-failure for both fi brous webs and A-SMEC composites signifi cantly increased as sample loading direction shifted to perpendicular to the fi ber direction ( χ increased). For example, ε f increased from 198 ± 62% for A-SMEC-0 to 351 ± 52%



for A-SMEC-90. (Table S5, Supporting Information pro-vides the complete set of results in tabular format). For comparison, Sylgard was found to feature a large strain-to-failure of ε f = 445 ± 49%, and is shown as a reference line in Figure 8 .

In the A-SMEC system, the Sylgard serves as the elas-tomeric matrix that, in part, increases ε f while also pro-viding shape recovery, discussed below. In comparison to our results, bat wing skin membranes featured lower failure strains spanning 13% < ε f < 232%, depending on fi ber bundle orientation angle. [ 8 ]

As was the case for E Y , the measured σ y for both fi brous webs and A-SMEC composites decreased monotonically with increasing orientation angle, χ (Figure 9 ). Also sim-ilar to our observations for E Y was the observation that the σ y values for fi brous webs were larger than those of the A-SMEC composites, regardless of orientation angle.





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Figure 10 . Shape-memory cycles for an A-SMEC where the fi bers are in the longitudinal direction (A-SMEC-0). The backdrop is a strain–temperature curve and the side face is a stress–strain curve of cycle 1.

This fi nding is attributed to the signifi cant decrease in fi ber volume fraction that occurs during A-SMEC fabrication.


Figure 11 . Shape-memory cycles for all samples tested as a function 45; d) A-SMEC-67.5; e) A-SMEC-90, and f) Sylgard.



3.4. Shape-Memory Characterization

All of the A-SMEC samples were tested for their SM behavior and it was found that shape fi xing ( R f ) decreased with increasing orientation angle, χ . Figure 10 shows the results from use of a four-step SM thermo-mechanical cycle method applied to the A-SMEC-0 composite. During Step 1 , the material was deformed at 80 ° C, a temperature where the PVAc fi bers are well above T g ( T − T g = 25 ° C) and the material behaves as an elastomer with fl uid-fi lled inclusions. Specifi cally, the tensile stress was ramped from a small pre-load value (for sample tensioning) to approximately 1 MPa, the sample responding with ten-sile elongation of ca. 100%. Next, during Step 2 , the mate-rial was cooled 0 ° C, well below the PVAc T g , while the tensile stress from Step 1 was maintained. During this step, little change in strain was observed. During Step 3 , the stress was released while the temperature was held at 0 ° C, allowing determination of the degree of strain fi xing, which is observed in Figure 10 to be exceed 95% for A-SMEC-0 for cycles 2 and 3. This indicates that the glassy fi brous web, whose prepared orientation was aligned with the strain axis of the SM test, is able to resist the elastic recovery stress applied by the Sylgard matrix. Finally, during Step 4 , the SM cycle was completed by heating to

9

of fi ber orientation angle for: a) A-SMEC-0; b) A-SMEC-22.5; c) A-SMEC-



10


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Figure 12 . Shape fi xing ratio ( R f ), black, and shape recovery ratio ( R r ), white, for A-SMEC samples as functions of fi ber orientation angle. The Sylgard (dashed) reference line for R f is shown.

80 ° C under negligible tensile stress, allowing observa-tion of the shape recovery. For A-SMCE-0, this recovery was large: R r ≈ 98%. Figure 10 also shows stress–strain and strain–temperature curves as dashed lines for cycle 1. Repetition of the SM cycle for this sample showed good repeatability for A-SMEC-0, with particularly close overlap of data for cycles 2 and 3.

SM testing of the type shown above for A-SMEC-0 was performed for neat Sylgard and all A-SMEC samples, revealing that R f decreased with increasing orientation angle, χ , while shape recovery was largely unchanged. Inspection of Figure 11 shows that all A-SMEC samples exhibited elastic deformation at 80 ° C and varying levels of fi xing, evident during Step 3 (unloading at 0 ° C) where surprisingly modest anisotropy was witnessed. While near-perfect fi xing was observed for A-SMEC-0, this fi xing level gradually decreased with increasing χ . Near-com-plete recovery was observed among all A-SMEC samples upon heating to 80 ° C, as expected considering the incor-poration of Sylgard as the composite matrix.

Trends in R f and R r , computed using Equation 1 and Equation 2 , are shown in Figure 12 . Good shape fi xing, R f = 95.6%, and recovery, R r = 99.3% were observed for A-SMEC-0. The composite with transversely oriented fi bers, A-SMEC-90, featured surprisingly high fi xing, R f = 86.2%, and expectedly good recovery, R r = 99.0%. Inspec-tion of the trend in shape fi xing with orientation angle showed a minimum for χ = 67.5 ° ( R f = 72.1%). Shape fi xing is good for all composite orientations and not the smallest for the transversely oriented sample (A-SMEC-90), which suggests the possibility that the strain imposed during Step 1 of the SM cycle alters the reinforcement micro-structure in a complex manner (i.e., the average orienta-tion angle, among other characteristics not characterized here, of the A-SMEC samples may be altered by the strain itself), impacting shape fi xing. This possibility is of great



interest to us for future study. R f and R r values are shown in Table S8 in the Supporting Information.

4. Conclusion

We have fabricated and characterized A-SMEC by using aligned PVAc fi bers to reinforce a Sylgard elastic matrix, demonstrating anisotropic mechanical properties. Tem-perature-dependent, anisotropic mechanical proper-ties for A-SMECs with varying orientation angle, relative to the loading axis were measured. Anisotropy of the A-SMEC was quantifi ed through tensile testing, specifi -cally obtaining the E A , E Y , ε f , and σ y . Anisotropy was fur-ther quantifi ed through SM testing to show the fi xing and recovery behavior as a function of fi ber orientation angle, χ . A-SMEC composites were found to feature decreasing fi xing capacity, with increasing χ , while shape recovery was invariably high. We anticipate application develop-ment of A-SMEC materials that exploits the unique combi-nation of elastomeric behavior, anisotropy, and SM.

Supporting Information

Supporting Information is available from the Wiley Online Library or from the author.

Acknowledgements: P.T.M. acknowledges generous fi nancial support for this work from AFOSR/NA, contract #: FA9550-09-1-0195. E.D.R. also acknowledges Maria Mintskovsky, high school student, for her hard work on fi ber-diameter analysis.

Received: February 7, 2013 ; Revised: March 5, 2013; Published online: ; DOI: 10.1002/macp.201300086

Keywords: elastomers; mechanical anisotropy; shape-memory polymers; silicone

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