Well-Defined Shape-Memory Networks with High Elastic Energy

Article pubs.acs.org/Macromolecules

Well-Defined Shape-Memory Networks with High Elastic Energy Capacity Christopher L. Lewis, Yuan Meng, and Mitchell Anthamatten* Department of Chemical Engineering, University of Rochester, 206 Gavett Hall, Rochester, New York 14627-1066, United States S Supporting Information *

ABSTRACT: Controlling network architecture and chain connectivity is critical to understanding elastic energy storage and improving performance of shape-memory polymers. Acrylate-terminated poly(caprolactones) were converted into thermoset networks by three different reactions: conventional free radical polymerization, radical-induced coupling with multifunctional thiols, and base-catalyzed Michael addition with multifunctional thiols. The highly efficient thiol−acrylate coupling reaction ensures that the molecular weight between cross-links is uniform, resulting in tougher, more elastic materials with a high degree of crystallinity and outstanding shapememory properties. Elastomers can be cold drawn to achieve several hundred percent of strain, and upon heating, nearly complete shape recovery is observed. Shape fixity upon cold drawing is correlated to the degree of strain-induced crystallization which is influenced by the draw rate and stress treatment immediately following cold drawing. Slow unloading of samples drawn to 400% elongation indicates the material is capable of storing greater than 1.5 MJ/m3 of elastic energy.

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INTRODUCTION Shape-memory polymers (SMPs) form an exciting class of materials that can store and release elastic energy upon applying an external stimulus such as heat or light.1−6 A shapememory material can be deformed to a temporary shape and can return to its original shape after the application of an external stimulus, most commonly temperature. For example, a material heated above its shape-memory transition temperature, TSM, can be elastically deformed by subjecting it to external stresses and subsequently cooled, while under stress, beneath TSM. In the cooled state, external stresses can be removed, and the material retains its deformed shape. Upon heating above TSM, the material recovers its elastic strain energy and returns to its original shape. SMPs are noted for their ability to recover from extremely large strainsup to several hundred percent which are imposed by mechanical loading. The large-strain recovery observed in SMPs is a manifestation of entropy elasticity. SMPs are especially recognized for their potential to serve in biomedical devices such as vascular stents, clot-removal devices, catheters, programmable sutures, implants, and numerous others. Applications increasingly demand that shape-memory materials perform significant mechanical work against external loads; therefore, SMPs must exhibit high shape energy densities which, unfortunately, are seldom measured.7 Other commercialization requirements are diverse but may include: (i) a clearly defined shape recovery stimulus (e.g., heat, light, chemical), for heat stimulated materials, a tunable shape recovery temperature, TSM; (ii) ease of processability into different shapes; (iii) highly reproducible and robust shapememory behavior upon cycling; and (iv) low cost and straightforward scale-up. © XXXX American Chemical Society

Polymers that undergo strain-induced crystallization can simplify the shape programming step of a shape-memory cycle. Kraft and co-workers have demonstrated cold-drawable, aramid-containing polyamides with crystallizable poly(ε-caprolactone) (PCL) soft segments that undergo strain-induced crystallization at room temperature.8−11 Upon load removal, the material remains deformedas if plastic deformation has occurredhowever, upon heating above its soft segment melting temperature, the material reverts to its original shape. Similar observations have been reported for other materials, including polyurethanes and natural rubber.12−17 Morphological and crystal texture changes that occur during cold-drawing have been studied for PCL-based polyurethanes. As samples are strained, the amorphous PCL phase is initially oriented, followed by reorientation of both hard segment domains and the crystalline PCL phase, where the latter undergoes stress induced disaggregation and recrystallization in the direction of the applied load.15 Strain-induced crystallization of PCL has enabled two-way shape memory18 and recently stress-free shape actuation.19 The role of network architecture on strain-induced crystallization and on the performance of cold-drawn shapememory polymers remains unclear. It is well-known that network characteristics such as the molecular weight distribution between cross-links, cross-link functionality and spatial distribution, and the presence of dangling chains and primary loops have large effects on elastomeric properties.20,21 A Received: April 12, 2015 Revised: July 1, 2015

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Figure 1. Synthesis of semicrystalline network elastomers with different network structures from a three-arm poly(caprolactone) prepolymer. Prepolymer acrylate end-groups are cross-linked by free radical polymerization, thiol−acrylate coupling, or both. Cartoons depict the expected network structures: green circles denote trifunctional cross-links originating from the prepolymer; orange circles are multifunctional cross-links from radical polymerization of acrylate end-groups; and red symbols are bifunctional and tetrafunctional cross-links from thiol−acrylate coupling with multifunctional thiol reagents.

Table 1. Characteristics of Synthesized Poly(caprolactone) Networks sample PCL-FR PCL-2T-MA PCL-4T-MA PCL-2T-FR

gel fraction (%) 99.5 99.0 98.9 89.4

E′RTa [MPa] 7.5 334 187 121

e

E′HTb [MPa]

elongation at breakc (%)

ultimate strengthc [MPa]

d Meff [g/mol] c

8.70 1.45 3.84 2.06

130 1410 ± 40 119 ± 6 670 ± 250

2.2 22.7 ± 0.9 8.4 ± 0.62 13.8 ± 3.3

1100 6700 2500 4700

Room temperature storage modulus taken at 25 °C from DMA. bHigh temperature storage modulus taken at 70 °C from DMA. cObtained from a strain-to-break experiment performed under tension at 60 mm/min. dEffective molecular weight between cross-links; Meff c = 3ρRT/E′HT where ρ was taken as 1.14 g/cm3 and temperature was 70 °C. eThis value is significantly lower because it was obtained above the sample’s melting temperature. a

addition were compared to those cross-linked using freeradical-initiated strategies. Materials prepared via the highly selective and efficient thiol−acrylate Michael addition reaction exhibit outstanding mechanical and thermal characteristics, enabling good shape-memory behavior and high energy density.

common approach to controlling network architecture is to synthesize networks from polymer precursors with crosslinkable end-groups. For example, Messori et al. have prepared cross-linked, semicrystalline networks from linear, three-arm, and four-arm star PCL precursors bearing methacrylate endgroups.22 Resulting networks exhibited nearly full recovery, and the melting temperature could be tuned by adjusting the crosslink density. In a related study, mild sol−gel cross-linking was employed instead of a free radical process to better control network structure, and resulting networks exhibited two-way shape memory with nearly full recovery.23 While these studies have clarified the importance of establishing model networks, a shape-memory PCL network with high crystallinity, nearly monodisperse and fully reacted network strands, and welldispersed junctions of known functionality has not yet been accomplished. The current effort aims to further understand how molecular architecture influences cold-drawability and shape-memory properties of PCL-based covalent networks. The study demonstrates the importance of a priori knowledge and design of network strand characteristics including cross-link density and the cross-link functionality. Three-arm PCL prepolymers containing acrylate end-groups were synthesized and crosslinked, fixing the molecular weights between cross-links. Network formation was accomplished via multiple synthetic pathways. The mechanical and crystallization behavior of thermally cured networks prepared via thiol−acrylate Michael

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RESULTS AND DISCUSSION Cold-drawable poly(caprolactone) (PCL) networks were prepared with nearly the same chemical composition but deliberately variant network architectures. As shown in Figure 1, this was accomplished by cross-linking an acrylate terminated PCL prepolymer via different reaction pathways. This section will include a description of network synthesis, tensile characterization of as-prepared networks, differential scanning calorimetry of both undrawn and cold-drawn samples, and the shape-memory performance of a sample with the most ideal network characteristics. Network Synthesis. A three-arm PCL prepolymer was synthesized by ring-opening polymerization of caprolactone using glycerol as a trifunctional seed.24 Terminal hydroxyl groups were acrylated, and the molecular weight of the prepolymer was evaluated by 1H NMR end-group analysis and MALDI-TOF spectrometry to be about 6200 g/mol (see Supporting Information, Figure S1). This molecular weight is low enough to ensure that chain entanglements do not interfere with interpretation of results by acting as temporary cross-links. The same batch of PCL prepolymer was used to prepare four B

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thiols is greatly favored over propagation by radical addition to the acrylate moiety.27 Mechanical Properties. Dynamic mechanical analysis (DMA) temperature ramp experiments were performed on all four cured networks. Experiments were performed in tension at 1 Hz and 5 °C/min over the temperature range 25−180 °C. Figure 2 shows that all networks formed by thiol−acrylate

networks with different topologies, according to Figure 1. In the following, we will discuss the synthesis and properties of each network; key results from dynamic mechanical analysis and tensile stress-to-break experiments are summarized in Table 1. Thermal-induced free radical polymerization of the prepolymer’s terminal acrylate groups resulted in PCL-FR networks. FT-IR analysis confirmed that acrylate groups were nearly completely consumed, and resulting networks possess the highest gel fraction among networks studied (Supporting Information, Figure S2 and Table S1). The PCL-FR network is brittle and exhibits rather poor mechanical properties, and this is attributed to the high cross-link density and the network’s poorly defined topology. The network contains trifunctional branch points from the prepolymer as well as multifunctional cross-links from polymerization of acrylate end-groups. The number of branches emanating from each cross-link is undetermined and is equal to the degree of polymerization at that site. Consequently, these cross-link sites are believed to serve as topological constraints, greatly reducing ductility and, as will be discussed later, suppressing the crystallization temperature to below ambient temperature. These attributes render the PCL-FR unfit for use as an ambient temperature, strain programmable shape-memory material. To achieve a more uniform spatial distribution of net points with defined functionality, networks were prepared using basecatalyzed, thiol−acrylate coupling of prepolymers with multifunctional thiols. This reaction proceeds as the base catalyst abstracts a thiol proton, forming a thiolate anion, followed by addition of the anion to an electron-deficient double bond. The resulting carbon-centered anion finally rearranges, taking a proton from the conjugate acid to form a thiol−ether linkage.25 The use of thiol−acrylate chemistry to form networks also offers processing advantages including reduced oxygen sensitivity, delayed gelation, and the possibility of inductive curing.26 The reaction curing kinetics were studied using oscillatory shear rheometry, and results are provided as Supporting Information (Figure S3). Both PCL-4T-MA and PCL-2T-MA networks exhibit high gel fractions (∼99%) following solvent swelling and drying, although FT-IR analysis (see Figure S2) showed small amounts of unreacted acrylate groups, indicating some dangling ends may be present. Both networks contain trifunctional branch points from the prepolymer. PCL-4T-MA also contains tetrafunctional net points, whereas PCL-2T-MA contains linear linkages, effectively doubling the chain length between trifunctional branch points. Because of its higher cross-link density, PCL-4T-MA displays a lower elongation at break and is much stiffer than PCL-2T-MA in the rubbery state (T > Tm). At room temperature the opposite is true: PCL-2T-MA exhibits the highest modulus (E′ = 334 MPa) which is attributed to its higher crystallinity, arising from fewer branch points. Radical-induced, thiol−acrylate coupling of stoichiometric mixtures was also conducted to achieve PCL-2T-FR networks. The formed network’s mechanical properties were intermediate to those of PCL-FR and PCL-2T-MA. This may be expected because generated free radicals can facilitate thiol−acrylate coupling, but they can also initiate acrylate polymerization, resulting in an ill-defined distribution of cross-links. In an ideal free-radical thiol−acrylate reaction, initiated free radicals result in thiyl radicals which add to acrylate moieties to create radical intermediates [R−S−CC•−R′]. Homopolymerization is only avoided if transfer of radical intermediates to unreacted

Figure 2. Storage modulus of PCL networks from dynamic mechanical analysis versus temperature.

Michael addition become much softer near the PCL melting temperature. This transition is not observed in the network formed by conventional radical polymerization (PCL-FR) because its melting transition is beneath the experimental range. At temperatures above the PCL melting transition, the modulus can be used to estimate the effective molecular weight between cross-links by the relationship

Mceff = 3ρRT /E′HT where ρ is density and E′HT is the storage modulus. The estimated values of Meff c for the networks using base-catalyzed Michael addition (PCL-2T-MA and PCL-4T-MA) are 6700 and 2500 Da, respectively. These values roughly agree with ideal values of 4132 and 2066 Da obtained if prepolymers (6200 Da), each containing three identical arms, were fully reacted with stoichiometrically balanced, multifunctional thiols. This agreement is further evidence that the reaction of network prepolymers proceeded to high conversion. Note that PCL-FR exhibits an Meff c ∼ 1100 Da, further suggesting it is a more highly cross-linked network. Tensile tests were performed on PCL networks prepared by free radical polymerization (PCL-FR), thermal-induced free radical thiol−acrylate coupling (PCL-2T-FR), and basecatalyzed thiol−acrylate Michael addition (PCL-2T-MA). All three samples were prepared from the same batch of prepolymer. Representative nominal stress−strain curves are shown in Figure 3. Networks formed by thiol−acrylate Michael addition, with more uniform bond topology, are mechanically superior compared to their radically grown analogues. For example, PCL-2T-MA was able to be stretched to over 1000% strain prior to failure whereas PCL-2T-FR and PCL-FR exhibited average elongations to failure of 670 and 130%, respectively. The PCL-2T-MA sample showed sawtooth oscillations in stress at high strain that are attributed to a small amount of cold drawing beyond the gauge region of the dogbone specimen at high elongation. C

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Figure 3. Representative stress−strain curves of poly(caprolactone) networks. The inset is the same data rescaled to lower stress−strain values.

Tensile test results can be explained by considering how stress is distributed on network subchains. As suggested by Figure 4, when PCL-FR is stretched, the applied stress is

Figure 5. DSC cooling and second heating scans (5 °C/min) of the triarm prepolymer and poly(caprolactone) networks formed by free radical polymerization of acrylate end-groups (PCL-FR), radical induced thiol−acrylate coupling (PCL-2T-FR), and thiol−acrylate Michael addition (PCL-2T-MA and PCL-4T-MA). Exothermic heat flow (q) corresponds to the positive ordinate direction. The reported enthalpy changes are averages of the integrated peak intensities corresponding to melting and crystallization.

Figure 4. Possible failure mode of (a) ill-defined networks made by radical polymerization of end-groups compared to (b) well-defined networks made by stoichiometric reaction of ends with multifunctional thiols. Cartoon symbols are described in Figure 1.

temperatures are greatly reduced such that the network is amorphous at room temperature. These trends are attributed to PCL-FR’s topologically frustrated network where strands emanate from cross-links sites and are covalently bound to the network. Strands are highly entangled and are unable to easily sample their configurational space. In comparison, PCL2T-MA, prepared using thiol−acrylate Michael addition, exhibits much higher crystallinity (X = 0.36). Moreover, all networks formed by the highly selective Michael Addition reaction are semicrystalline at room temperature and display sharp, unimodal thermal transitions which are consistent with more uniform network topographies. Cold-drawing and shapememory studies discussed later will focus on PCL-2T-MA because it displayed such exemplary crystallization and melting behavior and mechanical properties. The calorimetry scan of PCL-2T-FR, with the possibility of concurrent free radical polymerization, displays a low-temperature shoulder (∼17 °C). The shoulder at lower temperatures coincides with PCL-2T-MA’s endotherm, suggesting some chains form the same type of well-defined thiol−acrylate substructures. However, the more intense part of the peak appeared at a slightly higher temperature (21 °C), indicating that many chains are less constrained and can crystallize more easily. Some end-groups may have experienced free radical polymerization, causing a stoichiometric excess of thiol reactive groups and resulting in more easily crystallized dangling ends.

concentrated onto a small fraction of subchains. These chains must bear the entire applied load and thus are more vulnerable to breakage. On the other hand, PCL-2T-MA offers a more well-defined topology, and the applied stress is distributed more uniformly among subchains, greatly lowering the forces on individual chains. Consequently, PCL-2T-MA exhibits a high breakage strain (>1000%) and hence has a wide shape-memory working strain range. Sakai and co-workers have applied similar reasoning to design high-strength hydrogels.28 Crystallinity of Undrawn Networks. Network crystallization is highly sensitive to the type of network structure, as indicated by the data in Figure 5. The prepolymer exhibits sharp crystallization and melting transitions. A bimodal melting endotherm was observed upon melting of the prepolymer, and both peaks were integrated when calculating the enthalpy of fusion. Bimodal peaks have been observed in other studies involving PCL materials.18,29 The origin of bimodal melting is not entirely clear; however, An and co-workers have reasoned that microphase segregation plays an important role,29 and we suspect that the prepolymer acrylate chain-ends are microphase segregated from PCL-rich phases influencing their melting temperature. Based on the reported enthalpy of melting of fully crystalline PCL (ΔHm = 135 J/g),30 the prepolymer’s degree of crystallinity is X = 0.53. When acrylate end-groups are radically polymerized (PCL-FR), the degree of crystallinity is much lower (X = 0.21), and the melting and crystallization transition D

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Macromolecules Crystallinity of Drawn Networks. Strain-induced crystallization of PCL-2T-MA was studied by cold-drawing at different rates. Samples were drawn to 400% strain at 3, 30, and 300 mm/min. After drawing, the load was immediately removed, except for one test where the strain was maintained for an additional 10 min. The strain fixities εf following cold drawing were all between 66 and 72% for samples immediately unloaded, and the sample held at maximum strain for extra time showed a somewhat higher fixity of 84%. Our results are in agreement with a study by Wong and co-workers on un-crosslinked biodegradable shape-memory polymers.31 They observed the highest fixities when lower stress was applied for a longer duration. The DSC heating scans displayed in Figure 6 were acquired within 4 h of cold-drawing to 400% strain. Although cold

drawing does not significantly change the degree of crystallinity, it shifts a part of the melting transition to higher temperatures. This effect is most pronounced at the slowest draw rate (3 mm/min) because the cold-drawing process take longer, and this allows more time for strain-induced recrystallization to occur. This reasoning also explains why the cold-drawn sample held at 400% strain for an additional 10 min shows the most elevated melting temperature and the higher level of shape fixity. Mechanistically, cold drawing a PCL-rich material just beneath its melting point reorients and disrupts small crystallites and significantly stretches network strands, removing configurational entropy.32 The amount of entropy lost upon forming a crystal (ΔSm) is lessened, causing an increase in the melting point, Tm = ΔHm/ΔSm, and consequently, new crystals form. PCL is known to crystallize in an orthorhombic unit cell with polymer chains oriented parallel to the c-axis.33,34 The orientation of crystal unit cells with respect to the draw direction was probed using wide-angle X-ray scattering (WAXS). Figure 7 shows orientation arcs for the two most prominent planes: the (110) and (200) planes. There is considerable variation in intensity along the azimuthal angle ϕ, i.e., the angle between the cold-draw direction and the diffraction planes. The azimuthal position of both the (110) and the (200) arcs at the equator indicates that the c-axis of oriented unit cells is parallel to the draw direction. The degree of orientation can be further analyzed by considering the azimuthal variation of diffracted X-rays from the (110) plane near 2θ = 21.5°, also shown in Figure 7. Unfortunately, diffraction from the (110) planes convolutes with scattering from amorphous material. To separate these contributions, we assume that amorphous scattering is about the same at slightly lower angles (14.7° < 2θ < 19.8°), where strong Bragg peaks are absent. Then, the corrected azimuthal variation of intensity arising from the (110) planes can be analyzed with Hermans’ orientation function.35 The resulting orientation factor can then be calculated as

Figure 6. Influence of cold-draw rate on melting for PCL-2T-MA: DSC heating scans (5 °C/min) and melting enthalpies for specimens cold-drawn at 3, 30, and 300 mm/min. The dashed curves refer to the sample prior to cold drawing. The bold solid line in the middle plot was obtained after holding the cold-drawn sample for 10 min at 400% strain.

f110 =

3⟨cos ϕ⟩2 − 1 2

(1)

where

Figure 7. Determination of unit cell orientation and the degree of order following cold-drawing of PCL-2T-MA: two-dimensional WAXS data of (a) unoriented samples and (b) samples cold drawn to 400% strain. (c) Azimuthal variation of scattered intensity from the (110) plane. The arrows on the WAXS images indicate the cold-draw direction. E

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Figure 8. Conventional, stress-free shape memory performance of PCL-2T-MA networks: (a) successive shape memory cycles and (b) shape fixity and recovery ratios plotted against cycle number.

Figure 9. Measurement of energy storage capacity for PCL-2T-MA: (a) slow-unload shape memory cycles and (b) a plot of stored elastic energy density and efficiency plotted against cycle number.

explains why the stress required for cold drawing in the first cycle is distinctly higher than subsequent cycles because the sample initially had a high degree of crystallinity. Subsequent cycles were performed after a 10 min waiting time at room temperature, and longer hold times resulted in higher colddraw stresses. All cycles demonstrate high shape fixity and nearly complete shape recovery. The high shape fixity is attributed to enhanced crystallization of the uniform network strands when they are elongated. The outstanding shape recovery is due to the permanent, covalent network with a unimodal chain-length distribution that allows all chains to more equally deform at high strain. Slow-unload shape-memory cycles were performed to assess the capacity of the fixed shape to store elastic energy. Results are shown in Figure 9. Samples were strained at 30 mm/min to 400% of the gauge length and held there for 5 min. The load was removed, and the strain was allowed to elastically recover to the sample’s fixed shape. Next, the strain was held constant while the sample was heated to 60 °C, causing a stress to develop. After 70 s of melting the sample was then unloaded at a rate of 0.5 MPa/min, allowing the sample to perform measurable work as it retracted. The melting of 70 s was chosen as it yielded the maximum amount of stored elastic energy. Upon reaching a zero-stress condition, the crosshead was returned to its original zero-strain value, and the sample was allowed to fully recover its permanent shape. As before, the sample was allowed to crystallize at room temperature for 10 min before starting the subsequent cycle.

π

⟨cos ϕ⟩ =

∫0 I110(ϕ) cos2 ϕ sin ϕ dϕ π

∫0 I110(ϕ) sin ϕ dϕ

(2)

This analysis shows that f110 ∼ −0.30 for the cold-drawn sample, indicating that the normal vectors to the (110) planes are preferentially oriented perpendicular to the draw direction; this is consistent with the polymer main chains being preferentially oriented along the draw direction, as discussed earlier. Shape-Memory and Elastic Energy Storage Capacity. PCL-2T-MA, with its well-defined network formed by basecatalyzed Michael addition, exhibits exemplary shape-memory properties. SMP behavior was characterized in two distinct ways: (i) a conventional, stress-free SMP test program where the material is allowed to recover without an external load and (ii) a slow-unloading program that slowly reduces the load at constant rate, thus allowing for the assessment of the elastic energy storage capacity. Results from a typical, conventional shape-memory experiment are shown in Figure 8. A sample was drawn at 30 mm/min to 400% strain at room temperature, held for 5 min to allow for crystal re-formation, and the load was released. The sample was allowed sufficient time (30 s) to equilibrate to its temporary shape. It was then heated, in the absence of a load, to a temperature above PCL’s melting temperature, 60 °C, to allow for stress-free shape recovery to its permanent shape. The stress−strain behavior observed during the cold-drawing step varied according to the amount of time that the sample was allowed to crystallize between cycles. This F

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equiv), glycerol (807 mg, 1 equiv), and SnOct2 (88.8 mg, 1/40 equiv) were added. The reaction was carried out in bulk at 120 °C under continuous N2 purge for 24 h. The resultant PCL polymer was purified through precipitation into methanol and vacuum-dried at 60 °C overnight. Yield 98.6%. Molecular weight of the obtained prepolymer was determined to be 6100 g/mol (NMR end-group analysis), and the molar mass dispersity was 1.18 (GPC). PCL-Prepolymer. The hydroxy-terminated prepolymer was acrylated by nucleophilic substitution with acryloyl chloride to form a three-arm prepolymer. To PCL-triol (30.0 g, 1 equiv) under N2 at 0 °C was added freshly distilled toluene and triethylamine (2.273, 1.5 equiv). After 30 min of degassing with N2, acryloyl chloride (1.82 mL, 2.035 g, 1.5 equiv) was added dropwise over 2 h. The reaction was warmed to 80 °C and allowed to stir for 48 h. The mixture was filtered, and the liquid fraction was precipitated into methanol to afford a white powder. The product was vacuum-dried at 60 °C overnight. Yield 96.0%. NMR indicated a molecular weight of 6200 g/mol and complete end-group conversion while GPC showed a molar mass dispersity of 1.20. MALDI-TOF spectrometry, provided as Supporting Information (Figure S1), was applied to compare PCL-triol to the prepolymer and showed peak offsets upon acrylation of M/Z = 162, indicating complete transformation of hydroxyl to acrylate end groups. PCL-FR. At 60 °C, PCL-triacrylate (1.0 g, 1 equiv) was mixed with thermal initiator benzoyl peroxide (BPO) (10 mg, 1 wt %). The mixture was immediately sandwiched between two glass slides separated by a 200 μm thick Teflon spacer. The sample was allowed to cure for 48 h at 60 °C, and IR spectroscopy (8000S, Shimadzu) was employed to assess end-group conversion. PCL-2T-MA. At 60 °C, a stoichiometrically balanced mixture was formulated with PCL-triacrylate (1.0 g, 1 equiv), 2,2′-(ethylenedioxy)diethanethiol, a dithiol (43.5 mg, 1.5 equiv), and a catalytic amount of 4-(dimethylamino)pyridine (DMAP) (10 mg, 1 wt %). The mixture was immediately sandwiched between two glass slides separated by a 200 μm thick Teflon spacer. The sample was allowed to cure for 48 h at 60 °C, and IR spectroscopy (8000S, Shimadzu) was employed to assess end-group conversion. PCL-4T-MA. At 60 °C, a stoichiometrically balanced mixture was formulated with PCL-triacrylate (1.0 g, 1 equiv), pentaerythritol tetrakis(3-mercaptopropionate) (PETMP), a tetrathiol (58.3 mg, 1.5 equiv), and a catalytic amount of DMAP (10 mg, 1 wt %). The mixture was immediately sandwiched between two glass slides separated by a 200 μm thick Teflon spacer. The sample was allowed to cure for 48 h at 60 °C, and IR spectroscopy (8000S, Shimadzu) was employed to assess end-group conversion. PCL-2T-FR. At 60 °C, a stoichiometrically balanced mixture was formulated with PCL-triacrylate (1.0 g, 1 equiv), dithiol (43.5 mg, 1.5 equiv), and BPO (10 mg, 1 wt %). The mixture was immediately sandwiched between two glass slides separated by a 200 μm thick Teflon spacer. The sample was allowed to cure for 48 h at 60 °C, and IR spectroscopy (8000S, Shimadzu) was employed to assess endgroup conversion. Characterization. Differential scanning calorimetry was performed using a TA Instruments, Q2000 DSC. Five to six mg samples were placed in a hermetically sealed pan and subjected to heating and cooling at 5 °C/min over the temperature range of −50 to 100 °C. Dynamic mechanical analysis (DMA) temperature ramp experiments were performed in tension at 1 Hz and 5 °C/min over the temperature range 25−180 °C using a Rheometrics RSAII solids analyzer. Data were analyzed using the commercially available TA Orchestrator software program. Two-dimensional wide-angle X-ray scattering (WAXS) were performed on Brüker general area detector diffraction system (GADDS), with a highly collimated 1 mm diameter beam and Brüker HI-STAR area detector. All stress−strain, shape-memory, and cold-drawing experiments were performed using an MTS tensile testing apparatus (Q Test/5). All measurements were made on film samples (0.2 mm thickness) that were cut using a trim die according to ASTM D638-Type 5 (gauge length = 7.62 mm). Unless otherwise specified, tensile tests were performed at a strain rate of 30 mm/min. Shape memory experiments were performed using a custom-made heating chamber. Following

The elastic energy storage capacities of each cycle (Figure 9b) were calculated by integrating force versus extension curves during the shape-recovery phase and dividing the result by the sample volume. The energy efficiency of each cycle could be calculated based on the ratio of the work recovered to the work input during cold drawing. The measured energy storage capacities of PCL-2T-MA exceeded 1.5 MJ/m3, and this corresponds to an efficiency of about 13%. These values are among the highest reported for shape memory polymer that are fixed at around 300% strain.7 The measured energy storage capacity was found to be sensitive to the following parameters: (i) the time the sample was held in the elongated state as shown in the Supporting Information (Figure S4), (ii) the time the sample was heated while maintaining constant elongation, and (iii) the rate that the sample was unloaded. It is important to hold these parameters constant to allow for valid comparisons between samples.

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CONCLUDING REMARKS Well-defined shape memory networks were formed by the basecatalyzed reaction of acrylate-terminated poly(caprolactone)s with multifunctional thiols. Resulting networks exhibit desirable thermal and mechanical properties that are associated with appealing shape-memory properties. This was confirmed by evaluation of shape-memory properties, including cold drawability, high shape fixity, complete shape recovery, and a high capacity to store elastic energy. These outstanding shapememory properties are attributed to the network’s high density of crystallizable chains, its homogeneous topography, and its uniform chain-length distribution. In contrast, networks prepared by free radical polymerization of acrylate end-groups were less ideal and did not show such outstanding mechanical and thermal behavior characteristic of an ideal SMP. During cold drawing, PCL networks re-form crystals that exhibit heightened thermal stability, and the polymer chains within the crystalline unit cells are oriented predominantly along the cold-draw direction. The measured energy storage capacity is dependent on its elastic stress−strain history. This study encourages additional investigations to understand mechanically assisted crystallization and melting under external loads to further advance these and other shape-memory networks.

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METHODS

Polymer and Network Synthesis. ε-Caprolactone (CL) monomer, triethylamine (TEA), and glycerol were purchased from Sigma-Aldrich and were distilled before use. Prior to distillation, CL was dried over calcium hydride, and TEA was dried over KOH. All other chemicals were used as received from Sigma-Aldrich. NMR spectra were acquired on a Brü ker AVANCE-III 400 NMR spectrometer system operating at 400 MHz for 1H observation. Attenuated total-reflection Fourier transform infrared (FT-IR) spectroscopy (Shimadzu 8000S) was used to assess the presence of acrylate end-groups near the sample surface. FT-IR spectra of prepared networks are provided as Supporting Information (Figure S2). Molecular weight and molar mass dispersity were measured by size exclusion chromatography (PolyAnalytik PAS103-L and PAS104-L GPC columns and Viscotek TPA301 detector) using THF as an eluent and polystyrene standards. Matrix-assisted laser deionization/ionization mass spectroscopy (Brüker Autoflex III MALDI-TOF) was applied to determine the molar mass distribution. PCL-Triol. The hydroxy-terminated, three-arm poly(caprolactone) was synthesized by ring-opening polymerization of caprolactone (CL) in the presence of SnOct2 as catalyst and glycerol as trifunctional initiator. To a dried, silanized 100 mL flask, distilled CL (50.0 g, 50 G

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cold-drawing, most of the sample strain was confined within the gauge length. Following each shape-memory cycle, the cross-head was returned to the original zero-strain setting, and the temperature was reduced to room temperature to allow the sample to crystallize at room temperature for 10 min before starting the subsequent cycle. For cold-drawing, DSC, and WAXS experiments, samples were trained by performing three shape memory cycles with stress-free shape recovery prior to cold drawing.

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ASSOCIATED CONTENT

* Supporting Information S

(i) MALDI-TOF characterization of PCL network precursor; (ii) FT-IR spectroscopy of PCL networks; (iii) gel fraction measurements; (iv) rheological observation of network formation; (v) effect of hold-time on shape memory cycle. The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.5b00763.

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AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected]; tel (585) 273-5526, fax (585) 273-1348 (M.A.). Notes

The authors declare no competing financial interest.

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REFERENCES

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DOI: 10.1021/acs.macromol.5b00763 Macromolecules XXXX, XXX, XXX−XXX

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DOI: 10.1021/acs.macromol.5b00763 Macromolecules XXXX, XXX, XXX−XXX