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Letter Cite This: ACS Macro Lett. 2017, 6, 1290-1295
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Synthesis of Elastomeric Liquid Crystalline Polymer Networks via Chain Transfer Nicholas P. Godman,† Benjamin A. Kowalski,†,‡ Anesia D. Auguste,† Hilmar Koerner,† and Timothy J. White*,† †
Air Force Research Laboratory, Materials and Manufacturing Directorate, Wright-Patterson Air Force Base, Dayton, Ohio 45433-7750, United States ‡ Azimuth Corporation, 4027 Colonel Glenn Highway, Beavercreek, Ohio 45431, United States S Supporting Information *
ABSTRACT: Materials capable of complex shape changes have broad reaching applications spanning biomimetic devices, componentless actuators, artificial muscles, and haptic displays. Liquid crystal elastomers (LCE) are a class of shape programmable materials which display anisotropic mechanical deformations in response external stimuli. This work details a synthetic strategy to quickly and efficiently prepare LCEs through the usage of chain transfer agents (CTA). The polyacrylate materials described herein exhibit large, reversible shape changes with strains greater 475%, rivalling properties observed in polysiloxane-based networks. The approach reported here is distinguished in that the materials chemistry is readily amenable to surface alignment techniques. The facile nature of the materials chemistry and the compatibility of these materials with directed self-assembly methods could further enable paradigm shifting end uses as designer substrates for flexible electronics or as actuating surfaces.
P
liquid crystalline polymer networks (LCNs). Glassy LCNs retain the optical properties of the LCM in robust films enabling end use as light control films in displays. LCMs are readily amenable to surface alignment techniques including photoalignment. Numerous papers have reported on the stimuli-response of these materials and the ability to induce shape transformation within LCN.14−16 The thermomechanical properties of LCNs can be influenced by copolymerizing difunctional LCMs with monofunctional LCMs to reduce the molecular weight between cross-links. LCE-like materials have been prepared and characterized, exhibiting upward of 20−30% strain and semisoft elasticity.17 However, this approach inherently dilutes the main-chain character of the LCN/LCE and known to limit the optimum association of orientation and elasticity or actuation. Recently, we have reported on the preparation of main-chain type LCEs by oligomerizing LCMs via an aza-Michael addition reaction.18−22 This approach is amenable to spatially patterned surface-derived alignment.23−25 However, while this method has enabled initial studies of the directed self-assembly of LCEs, it is slow (>24 h per sample) and limited in the range of accessible cross-link densities. Here, we report on a distinct method to prepare LCEs via chain-transfer processes between thiols and acrylates.26−29 As illustrated in Figure 1b, thiol−acrylate reactions can proceed either through acrylate chain growth or thiol chain transfer
reprogramming complex mechanical responses into soft materials is a topic of considerable interest within the scientific community. Spatial organization of monomeric (oligomeric) precursors to form anisotropic polymeric materials have been discussed as potential paradigm shifts in stimuli-responsive drug delivery,1 sensing,2 and soft robotics3,4 (the focus of this work). While many approaches have been explored in the literature,5 three key aspects are quintessential to the design and implementation of novel, multifunctional material platforms: facile synthetic strategies, processability, and local regulation and control of the mechanical properties. Liquid crystalline elastomers (LCEs) have garnered significant attention in the past decade due to their nonlinear mechanical deformation and dramatic stimuli-induced deformation.4,6−9 The primary approach to prepare LCEs is via hydrosilylation of olefinic liquid crystalline monomers to produce polysiloxane elastomers.10 Polysiloxane elastomers are widely recognized for their high extensibilities and low moduli, largely attributable to Si−O−Si bond within the polymer backbone.11 However, alignment of polysiloxane LCEs has been limited to mechanical loading either during or after12,13 preparation to align the mesogens to the strain direction. Liquid crystalline monomers (LCMs) capable of free-radical photopolymerization (such as the commercially available diacrylates illustrated in Figure 1a) have been investigated for more than 30 years due to their straightforward fabrication and compelling optical, thermal, and mechanical properties. Prior examination of these materials have primarily focused on homopolymerization of difunctional LCMs to produce glassy © XXXX American Chemical Society
Received: October 17, 2017 Accepted: October 27, 2017
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The kinetics of radical initiated copolymerization of multithiol/di(meth)acrylate systems have been extensively detailed in the literature.29−33 The two competing reaction pathways of thiol-mediated chain transfer and acrylate homopolymerization (chain growth) can be investigated by using real time infrared (RTIR) spectroscopy to monitor the conversion of acrylate (CC) and thiol (S−H) bonds. The CnM-0.5B formulations showed near complete conversion of the acrylate groups within 60 s of UV irradiation, but the conversion of thiol groups was much lower (40−55%, Figure S1). Hence, the chain growth is occurring at a faster rate than the chain transfer which is consistent with similar nonmesogenic systems.30 As illustrated in Figure 1c, the ratio of acrylate to thiol functional groups converted with time showed similar trends regardless of the LCM carbon spacer employed. Therefore, any differences in materials properties arise from changes in cross-link density due to the spacer length as opposed to changes in acrylate reactivity. Monodomain samples were prepared in surface-aligned cells, enforcing planar alignment of the nematic director. Upon photopolymerization, all samples were optically transparent (Figure 1d) and demonstrated anisotropic mechanical properties (Figure 1e). The addition of 0.5 mol equiv of BDMT to either C3M, C6M, or C11M (9.5−12.6 wt %) strongly influenced the mechanical properties of the LCEs. When the films were stretched parallel to the director, only slight increases in strains to failure (εf) were observed, and all films were fractured before 40% strain with a strain rate of 5% min−1. Notably, the Young’s modulus (E) decreased over an order of magnitude from 1250 ± 80 MPa for C3M-0.5 to 92.9 ± 7.5 MPa for C11M-0.5. This decrease in E is related to the increase in the molecular weight between cross-links attributable to the longer spacer length of C11M. Uniaxial extension perpendicular to the principal axis of the film demonstrated three distinct features as the spacer length was increased: an exponential increase in εf (51.3 ± 10.1% to 445.1 ± 13.3%), a decrease in E of almost two decades (641 ± 34 to 9.1 ± 1.0 MPa), and a sizable decrease in the onset of the soft elastic plateau (15.2 ± 0.6 MPa to 2.8 ± 0.2 MPa). Complete listings of mechanical data and film characterization are given in the Supporting Information (Tables S1 and S2). Simply adding a thiol-based CTA into the commercially available LCMs examined here proves to be a straightforward method to quickly and effectively generate robust LCEs with elastomeric properties that rival polysiloxane LCEs. The synthetic strategy detailed here does not employ timely oligomerization reactions that may not be amenable to industrial scaleup. Wide angle X-ray scattering (WAXS) spectroscopy confirms the retention of the characteristic nematic diffraction patterns for C3M-0.5B and C6M-0.5B as well as evidence of a cybotactic nematic phase in C11M-0.5B (Figure S2). LCE films examined here did not exhibit crystallization over time as has recently been reported for similar material compositions.21 Encouraged by the performance of C11M-0.5B, the ratio of BDMT was varied to further explore the utility of CTAs in the fabrication of LCEs. Films were prepared with 0, 0.25, 0.5, and 0.75 mol equiv of BDMT (C11-XB). The polymerization of C11M with an equimolar amount of BDMT (C11M-1B) did not result in a cohesive, free-standing film. RTIR showed an increase in thiol incorporation within the LCE as the BDMT loading was increased (Figure S3), and all films fractured before 45% strain (Figure S4) upon elongation parallel to the director.
Figure 1. (a) Chemical structure of common liquid crystalline diacrylates (C3M, C6M, and C11M) and 1,4-benzenedimethanethiol (BDMT). (b) Potential reaction pathways are illustrated for the copolymerization of acrylates and thiol. (c) The conversion of acrylate and thiol functional groups for CnM-0.5B mixtures was monitored with real-time infrared spectroscopy (RTIR). (d) A representative C11-0.5B film demonstrating the excellent optical transparency. (e) Representative uniaxial tensile testing data for CnM-0.5B films. Strain rate = 5%/min.
(chain termination). In this study, we employ classical LCMs (C3M, C6M, and C11M) and react them with dithiol monomers to maximize the number of reaction sites for the nonliquid crystalline additive at low concentration. Initial compatibility screenings confirmed that alkyl dithiols generally had poor miscibility with C3M, C6M, and C11M. Hence, 1,4-benzenedimethanethiol (BDMT, Figure 1a) was used in this examination. Unless otherwise noted, all chemical formulations and the resulting films are designated as CnM-XB, where n is the number of carbons between the mesogenic core and the acrylate moiety and X is the molar equivalent of BDMT. 1291
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Figure 2. (a) Stress−strain curves for C11M-XB films aligned such that the nematic director was orthogonal to the stretch direction. Strain rate = 5% min−1. (b) The synthetic approach reported here produces robust, high quality films. Micrographs of a typical C11M-0.5B film, illustrating the birefringence. (c) WAXS patterns for C11M-XB films confirming the prevalence of the nematic phase. Note: the higher ordered “eyebrow” pattern indicate the presence of cybotactic packing.
The influence of the CTA concentration on the resulting mechanical response of C11M-XB films is illustrated in Figure 2a, which plots the strain measured in response to an applied stress with the nematic director oriented perpendicular to stretch direction. Increasing the CTA concentration dramatically affects the soft elasticity of the LCEs, both in the threshold stress and range for which soft elasticity is observed. For comparison, neat C11M films exhibit a linear elastic response until ε = 5% before yielding. The nonlinear response after the yield point (evident in Figure 2a) is likely due to mesogens attempting to rotate, but the high cross-link prevents complete reorientation. The addition of 0.25 mol of BDMT (5.0 wt %) increased εf by 300% and C11M-0.25B maintains a semisoft elastic plateau, unlike C11M. The semisoft elastic response of C11M-0.25B indicates incomplete mesogen reorientation. Increasing the concentration of BDMT within the network further resolves the soft-elastic plateau and increases the fracture strain to 486 ± 20% (C11M-0.75B). Increasing the concentration of BDMT also lowers the modulus from 260 ± 11 to 4.0 ± 0.4 MPa for neat C11M and C11M-0.75B, respectively. Glassy liquid crystalline networks prepared from diacrylate LCMs produce films with high optical clarity and birefringence. Figure 2b shows a typical C11M-0.5B film and demonstrates the birefringence associated with the fabricated LCEs. Maintaining a concentration of BDMT below ∼15 wt % yield materials with optical properties rivalling LCN films prepared from the neat LCM and elastic deformation that rival polysiloxane LCEs. The WAXS spectra for C11M-XB films all showed anisotropic diffraction patterns with distinctive diffuse, four-point “eyebrow” patterns that are characteristic of a cybotactic packing of the nematic phase (Figures 2c and S5).34−36 C11M-XB films also exhibited a decrease in the Tg and gel fraction with the corresponding decrease in cross-link density (Table S2). The materials chemistries reported here are also readily amenable to surface alignment. This is illustrated in Figure 3a, where a film is prepared with director orientation varied by 22.5° across the length. The effect of director orientation on the local elastic properties of the film can be observed in the birefringence changes of each region when the film is placed under cross polarizers (Figure 3b-i). The deformation of the material and its influence on the local orientation within each
Figure 3. (a) Illustration of an LCE film prepared with sequential variation in the orientation of the nematic director to the principal axes of the film. Local control of the orientation was directed by photoinduced surface alignment. (b) Localized deformation of the LCE film C11M-0.5B as a function of global strain, viewed between crossed polarizers (i−v). (c) Local strain within the LCE film C11M0.5B is quantified within each segment of the film and compared to the global strain of the film.
region are apparent in the images presented in Figure 3b-ii−v. The mesogens aligned parallel to the force direction exhibit linear elastic behavior under uniaxial tensile strain while the 1292
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Figure 4. (a) C11-0.5B film coated with TiO2 and corresponding 3D images at increasing temperatures. (b) Cross-sectional slices of the amplitude of the deformation of the C11-0.5B films as a function of temperature. (c) 3D images of C6M-C11M-1B films at increasing temperature. (d) Crosssectional slices of the amplitude of the deformation of the C6-C11-1B films as a function of temperature. Both films were 10 mm discs and all 3D profiles have the same overlaid height map.
tremendous formulation flexibility inherent to this synthetic approach, we prepare a new LCE formulation (C6M-C11M1B) that is nematic at room temperature (25 °C) and investigate the mechanical properties and shape-morphing capabilities. Surprisingly, planar, monodomain C6M-C11M-1B films display lower E and soft-elastic plateaus than C11M-0.5B, but also exhibit comparable total elongation (Table S1, Figure S7a) even though the material is composed of LCMs with various spacer lengths. When photopolymerized in a +1 azimuthal defect director profile, the resulting LCE is virtually flat, with an almost imperceptible anticone curvature at slightly below room temperature (20 °C, Figure 4c). The film completely flattens out when heated to 40 °C, and further heating to 200 °C results in a cone with a “circus tent” shape with a quantifiable diameter contraction (Figure 4d). This formulation is absent the cybotactic nematic phase seen in the WAXS diffraction (Figure S7b) for seen in C11M-XB samples. In conclusion, new LCE formulations have been developed by taking advantage of the competing chain growth and chain transfer reactions that occur between acrylates and thiol. Crucially, this fabrication method keeps the weight percent of thiol low and produces films with excellent optical clarity and birefringence that are typically seen in pure acrylate LCNs. The mechanical properties of the films can be tuned by varying the alkyl spacers in the backbone of the polymer and adjusting the molar ratio of dithiol employed. The materials are conducive to photoalignment techniques, which allows the local director profile to be manipulated. This was demonstrated by localizing soft-elasticity in a gradient pattern and in the production of shape memory materials corresponding to a +1 azimuthal defect. Polymerizing the materials at elevated temperatures resulted in a film with anticone geometry at low temperatures that morphed into a cone upon heating. This resulted in temperature-induced switching between positive and negative Gaussian curvatures in the same material. The residual thermal strain in the material could be suppressed by modifying the composition to broaden the nematic window and enable roomtemperature photopolymerization, while still achieving comparable mechanical properties. Ongoing work in our laboratory
other regions exhibit increasing degrees of soft elastic behavior. C11M-0.5B films were globally elongated to 125%, and Figure 3c quantifies the localization of strain within each region of the LCE film. For each variation of 22.5° in orientation, there is a corresponding increase in the local strain. This is associated with the amount that the mesogens within each region are able to reorient. The maximum contrast ratio between perpendicular and parallel director profiles is about 19. A topic of considerable recent research are shape-shifting materials, in which materials systems can be programmed to localize mechanical forces in response to stimulus. Photoalignment of liquid crystals23−25 has proven to be a valuable technique to direct the self-assembly of LCEs and enable dramatic shape transformation.18,37,38 Here, we organize the C11M-0.5B LCE into a director field described as a + 1 azimuthal defect (Figure S6) which will uniformly deform into a conical shape upon heating. Removal of the LCE film from the alignment cell and cooling to room temperature caused the film to spontaneously deform into an anticone (saddle) geometry (Figure 4a). This phenomenon was explicitly predicted in 201139 and is starkly evident in these experiments. In this instance, the presence of the anticone is the result of the elevated polymerization temperature used to prepare the C11M-0.5B film (65 °C). Cooling the materials to room temperature (25 °C) can result in a slight increase in the nematic order, inducing local average length changes within the network which manifests as residual strain. An optical profilometer was used to quantify the 3D shape transformation of C11M-0.5B (Figure 4a). Heating the film reduces the nematic order and allows the sample to return to a flat state. Increasing the temperature further produces the expected conical shape. Figure 4b plots a cross-sectional cutout of the profile of the film to illustrate the shape transformation with the corresponding increase in temperature. The LCE film prepared from C11M-0.5B transforms from negative Gaussian curvature to flat and ultimately to positive Gaussian curvature by simply changing the temperature. In certain applications, however, it may be preferred that the room temperature form be flat and not curved. Leveraging the 1293
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is already leveraging this approach to prepare auxetic materials, tunable diffraction gratings, and color changing elastomers.
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MATERIALS AND METHODS
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ASSOCIATED CONTENT
The authors declare no competing financial interest.
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C3M (1,4-bis[4-(3-acryloyloxybutyloxy)benzoyloxy]-2-methylbenzene), C6M (1,4-bis-[4-(6-acryloyloxyhexyloxy)benzoyloxy]-2-methylbenzene), and C11M (1,4-bis[4-(11-acryloyloxyundecyloxy)benzoyloxy]-2-methylbenzene) were purchased from Synthon Chemicals. BDMT (1,4-benzenedimethanethiol) was purchased from TCI America. Irgacure 651 (2,2-dimethoxy-2-phenylacetophenone) was donated by BASF. Elvamide-8023R was donated by DuPont. PAAD22 was purchased from Beam Co. C3M, C6M, and C11M were recrystallized from methanol prior to usage. All other materials were used as received unless otherwise noted. Liquid crystal cells were prepared using methods described elsewhere.18 Briefly, for cells patterned using rubbed surfaces, Elvamide was dissolved in methanol at 0.125 wt %. This solution was then spin-coated on plasma-cleaned glass and then rubbed with a felt cloth to introduce unixail alignment. For photoaligned cells, PAAD-22 in dimethylformamide (0.33 wt %) was spin-coated on plasma-cleaned glass. The glass was then baked at 100 °C for 10 min. For either cell type, two pieces of glass were glued together using Norland Optical Adhesive 65 mixed with 30 μm glass spheres to set the cell thickness and then UV cured for 5 min. After the cell was fabricated, photoalignment was carried out using exposure to 100 mW cm−2 of 445 nm laser light. The s-polarized output of a DPSS laser was spatially filtered, apodized, and then passed through a twisted-nematic spatial light modulator (HoloEye LC-2002) to rotate the angle of linear polarization appropriately in each region. A more detailed description of the photopatterning process employed is reported elsewhere.38 All formulations consisted of mesogen C3M, C6M, or C11M blended with BDMT. Irgacure 651 was used as a photoinitiator in concentrations of 0.5 wt %, and all formulations were prepared under red light. Monomer mixtures were prepared in a vial and melted at about 150 °C while vortexing repeatedly. The phase behavior of the monomer solution was investigated using polarizing optical microscopy with a heating stage. Melted formulations in the isotropic state were filled into a liquid crystal cell by capillary action at 100 °C. The cell was then cooled to 10 °C below TNI of the monomer and allowed to rest for 5 min. This allows the nematic defects to relax and for the monomer to take the order dictated by the surface. Samples were then polymerized using 365 nm UV light (ca. 150 mW cm−2) at 10 °C below TNI. Polymerization was carried out for 10 min, flipping the cell after 15 s and 5 min.
S Supporting Information *
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsmacrolett.7b00822. RTIR, WAXS, and additional tensile test data (PDF).
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REFERENCES
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AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. ORCID
Timothy J. White: 0000-0001-8006-7173 Funding
This work was enabled by financial support from the Materials and Manufacturing Directorate of the Air Force Research Laboratory and the Air Force Office of Scientific Research. A.A. acknowledges support from the Air Force Office of Scientific Research through a National Research Council (NRC) Postdoctoral Research Associateship. 1294
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