Preparation and Properties of a Main-Chain Smectic Liquid-Crystalline

Article pubs.acs.org/Macromolecules

Preparation and Properties of a Main-Chain Smectic LiquidCrystalline Elastomer with Shape-Memory Ability Aránzazu Martínez-Gómez,† Juan P. Fernández-Blázquez,‡ Antonio Bello,† and Ernesto Pérez*,† †

Instituto de Ciencia y Tecnología de Polímeros (ICTP-CSIC), Juan de la Cierva 3, 28006 Madrid, Spain IMDEA Materials Institute, Eric Kandel 2, 28906 Getafe (Madrid), Spain

‡

ABSTRACT: This article presents the preparation of a smectic liquid-crystalline elastomer from a main-chain polybibenzoate with cross-linkable methacrylate groups incorporated at the ends of the polymeric chains. The advantage of this functionalization is that it can be thermally cross-linked in the bulk state without initiator. Rheological experiments have established that temperatures around 20− 30 °C above the isotropization point were found to be the most appropriate for the preparation of the network, leading to excellent sheets with a superior crosslinking yield, and prepared in relatively short curing times. The DSC and X-ray diffraction experiments indicate that the network retains the liquid crystalline ability (with a slightly smaller order than that of the linear polymer). The shapememory behavior of the network has been determined from the well-known cyclic thermomechanical experiments. It follows that the investigated network shows a rather remarkable shape-memory behavior, with considerably good shape recovery and shape fixity values.

1. INTRODUCTION In recent years, the interest in liquid crystalline polymers has been renewed as precursors of liquid crystalline elastomers, LCEs, which exhibit remarkable mechanical and optical properties due to the combination of polymer network elasticity with the anisotropic structure of the liquid crystalline state.1,2 The most interesting property of LCEs is their ability to change their shape reversibly at the mesophase−isotropic phase transition, which brings potential applications of these materials as actuators.3−5 Most of the research about LCEs has focused on side-chain LCEs, where the mesogenic groups are attached to a flexible polymer backbone as lateral side groups.6,7 Polyacrylates, polymethacrylates, polysiloxanes, and polyurethanes have been widely used as polymeric backbones. However, elastomers based on main-chain LC polymers have received increasing attention recently8−11 due to the direct coupling between the liquid crystalline order and the conformation of the polymer chains. As a consequence, large changes in the sample dimensions should occur at the mesophase−isotropic phase transition, directly expressed in the thermoelastic behavior. Apart from the coupling between orientational order of mesogens and polymer backbone, the structure of the mesophase should also influence the elongation behavior. Smectic elastomers are expected to exhibit the strongest deviations from ordinary rubber elasticity behavior due to the conformational constraints imposed on the polymer backbone by the positional order correlations. Semiflexible polybibenzoates built by an alternating arrangement of flexible spacers and biphenyl mesogenic units represent a well-known family of smectic LC polyesters.12−17 It has been © XXXX American Chemical Society

demonstrated that it is possible with simple chemical and structural modifications in the flexible spacer, and/or with an optimum comonomer content, to tailor their thermotropic behavior: transition temperatures, mesophase structure, and mesophase stability. Following such strategies, a great diversity of polybibenzoates has been prepared, showing thermal transitions in a wide temperature range and different smectic mesophases with varying degrees of order. When the chemical structure is suitably chosen, the polymer shows a stable mesophase at easily accessible temperatures. In some cases, transition temperatures near room temperature are observed. Thus, semiflexible polybibenzoates might be interesting mainchain polymers for the preparation of smectic elastomers with tailored thermal transitions. Only a few works about networks based on polybibenzoates polymers are found in the literature,18−20 which were prepared by incorporation of trifunctional cross-linkers or photo-crosslinkable comonomers into the polybibenzoate chains. It has been shown that the isotropic state of these networks undergoes large and reversible rubberlike deformation and transforms into the liquid crystalline phase when it is highly elongated.18 This LC state returns to the isotropic phase with decreasing strain. Moreover, these studies mainly focused on structural changes occurring during uniaxial elongation of the smectic state and molecular reorientation in the course of spontaneous shape change in a cyclic heating−cooling process.19 Received: June 1, 2016 Revised: July 21, 2016

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DOI: 10.1021/acs.macromol.6b01166 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules

every day, until non-cross-linked polymer was extracted from the network. The absence of polymer in the chloroform was checked by precipitation of aliquots in methanol. Then, the network was deswollen by slow addition of methanol (a poor solvent) to the chloroform over a period of a few days. Finally, it was dried successively at room temperature and then at 50 °C and at 80 °C. After drying, the network was weighted again, revealing a weight loss of approximately 25%. 2.5. Techniques. Both 1H and 13C NMR spectra were recorded on a Bruker spectrometer operating at 300 MHz, in solutions of deuterated chloroform. Differential scanning calorimetry (DSC) measurements were performed under a nitrogen atmosphere in a TA Instruments Q100 calorimeter, provided with a cooling system. About 6−8 mg of the sample was encapsulated in an aluminum pan, and a scanning rate of 20 °C/min was used. The glass transition temperature, Tg, was taken as the temperature where the specific heat increment is half of the total one at the transition. Other transition temperatures were taken as the peak minimum or maximum in the calorimetric curve. The molecular weights were determined by size exclusion chromatography (SEC) in a Waters chromatography system with Styragel columns (300 × 7.8 mm, 5 μm nominal particle size). THF with 4% triethylamine was used as the eluent. The measurements were carried out at 35 °C at a flow rate of 1.0 mL/min using a RI detector. The polymer concentration was 3 mg/mL. The calibration was performed with monodispersed PMMA standards. The rheological behavior of the polymers was measured under oscillatory mode using a parallel plate rheometer (AR200EX, TA Instruments) with disposable aluminum plates. Samples of 25 mm in diameter and around 1 mm of thickness were placed between the plates of the rheometer and subjected to an oscillatory shear strain of constant amplitude and different frequencies at determined temperature. The dynamic viscoelatic response of the samples is given by the storage and the loss moduli, G′ and G″, respectively, which change with time as a result of cross-linking reactions. Tan δ expresses the ratio between the storage and loss moduli, tan δ = G′/G″. The evolution of the complex viscosity η* (modulus of real and imaginary parts) can be determined as

In this work, we describe the preparation of a smectic elastomer from a main-chain polybibenzoate, poly(triethylene glycol p,p′-bibenzoate) (PTEB), in which cross-linkable methacrylate groups were incorporated at the ends of the polymeric chains. The advantage of this methacrylate endfunctionalized PTEB is that it can be thermally cross-linked without an initiator to give a network in good yield. The aim of the present work is to study the properties and the thermal actuation of this new smectic main-chain elastomer in order to better understand the behavior of these exciting materials.

2. EXPERIMENTAL SECTION 2.1. Materials. Triethylene glycol (Fluka) was purified by distillation at reduced pressure. Dimethyl 4,4′-biphenyldicarboxylate (Aldrich) was purified by recrystallization in chloroform. Titanium(IV) isopropoxide, methacrylic acid, dicyclohexylcarbodiimide (DCC), and 4-(N,N-dimethylamino)pyridine (DMAP) were supplied by Aldrich and were used without further purification. 2.2. Synthesis of Poly(triethylene glycol p,p′-bibenzoate). Poly(triethylene glycol p,p′-bibenzoate) was synthesized by a two-step melt polycondensation of dimethyl-4,4′-biphenyldicarboxylate with triethylene glycol in the presence of titanium(IV) isopropoxide as catalyst. The first-step reaction was the transesterification of the diester with the glycol at around 200 °C for 24 h under a nitrogen atmosphere. In the second step, the polycondensation was performed at 250 °C under vacuum to provide a polymer of relatively high molecular mass. The reaction mixture was allowed to cool to room temperature and was dissolved in chloroform. This solution was poured into methanol to precipitate the polymer. Finally, the polymer was purified twice by precipitation from chloroform solutions into methanol and dried under vacuum. Molecular weight and molecular weight distribution of PTEB were determined as Mn = 21 600 and Mw/Mn = 1.8, respectively, from SEC (see below). The intrinsic viscosity, measured in chloroform at 25 °C using an Ubbelohde viscometer, was found to be 0.40 dL/g. δ (1H NMR, CDCl3) = 3.74 (s, 4H, OCH2CH2O), 3.86 (t, 4H, COOCH2CH2O), 4.49 (t, 4H, COOCH2), 7.62 and 8.10 (two doublets AB system, 8H, Har), ppm. δ (13C NMR, CDCl3) = 64.3 (COOCH2), 69.4 (COOCH2CH2O), 70.9 (OCH 2 CH 2 O), 127.3 (OCCCHCHC), 129.7 (O CCCHCHC), 130.4 (OCCCHCHC), 144.5 (OCCCHCHC), 166.3 (OCCCHCHC) ppm. 2.3. Preparation of Poly(triethylene glycol p,p′-bibenzoate) with Methacrylate End Groups. The introduction of methacryloyl groups at the chain ends of PTEB chains was carried out by esterification reaction of the OH end groups with methacrylic acid using the N,N-dicyclohexylcarbodiimide/aminopyridine method.21 Procedure. PTEB-OH polymer (5.20 g, 0.5 mequiv) was dissolved in 20 mL of anhydrous chloroform with amylenes as stabilizer. Then, methacrylic acid (0.17 g, 2.0 mequiv), DMAP (0.10 g, 0.8 mequiv), and DCC (0.41 g, 2.0 mequiv) were added to the polymer solution. The reaction mixture was left stirring at room temperature for 2 days. Then, the N,N-dicyclohexylurea precipitate was filtered, and the filtrate was poured into methanol. The precipitated polymer was purified by repeated precipitations from chloroform solutions into methanol until no impurities of reactants were detected in the NMR spectra of the polymer. Finally, PTEB-MMA was dried at room temperature under vacuum for several days. 2.4. Preparation of Poly(triethylene glycol p,p′-bibenzoate) Network. PTEB network was prepared in a Collin press by the following procedure. First a sheet of methacrylate ended PTEB polymer (PTEB-MMA), void-free and of suitable thickness, was prepared by compression molding. PTEB-MMA was melted between hot plates at 120 °C for 4 min, and then a low pressure of around 10 bar was applied during 1 min. Next, pressure was removed, and the polymer was cross-linked by heating at 120 °C for 60 min. After cooling to room temperature, the network was weighted, and it was placed in a glass Petri plate filled with chloroform. The resulting highly swollen network was kept there for several days, replacing the solvent

η* =

|G′ + G″| ω

where ω is the frequency of the imposed oscillatory strain. The function δ is a frequency-dependent variable representing the angle between the viscous stress and the shear stress. Isothermal experiments at frequencies of 20, 15, and 10 Hz with oscillatory shear strain amplitude of 1% were performed to study the curing process in LCEs. The thermomechanical properties have been studied by DMA in a TA Q800 dynamic mechanical analyzer, in a tensile and force controlled mode. The PTEB network was initially heated to 120 °C at 5 °C/min and then stretched at that temperature until a strain of 40% with force increasing at 0.05 N/min. Subsequently, it was cooled under constant strain down to 25 °C with a cooling rate of 5 °C/min. Then, the temporary shape was attained by removing the load at a rate of −0.05 N/min (until a minimum force of 0.001 N), followed by shape recovery when heating again to 120 °C at 5 °C/min. This thermomechanical cycle was repeated four times with the same parameters. The X-ray diffraction experiments were performed in beamline BL11-NCD at ALBA synchrotron (Cerdanyola del Vallés, Barcelona, Spain), at a fixed wavelength of 0.1 nm. The WAXS profiles were acquired with a Rayonix LX255-HS detector, placed at about 19 cm from sample and a tilt angle of around 30°. The calibration of spacings was obtained by means of silver behenate and Cr2O3 standards. The temperature control unit was a Linkam hot stage. The initial 2D X-ray pictures were converted into 1D diffractograms as a function of the inverse scattering vector, s = 1/d = 2 sin θ/λ. B


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Figure 1. General procedure for the synthesis of the methacrylate end-functionalized PTEB polymer (PTEB-MMA).

Figure 2. 13C NMR (75 MHz, CDCl3) spectra in the range δ = 60.0−74.0 ppm and peak assignment for (a) hydroxy-terminated PTEB polymer (PTEB-OH) and (b) methacrylate end-functionalized PTEB polymer (PTEB-MMA).

The first step of the polymerization, transesterification of the diester with the glycol at 200 °C under a nitrogen atmosphere, was prolonged to 24 h in order to ensure that all the ester groups COOCH3 were transesterified with the glycol and that a hydroxy-terminated polymer was finally obtained. As can be observed in Figure 2, the 13C NMR spectrum shows, besides the signals assigned to the triethylene glycol spacers in the polymer backbone, three signals (labeled as i, g, and h) of very low intensity that can be associated with the chain end ∼CH2OCH2CH2OH. On the other hand, no signal at around δ

3. RESULTS Hydroxy-terminated PTEB (PTEB-OH) has been prepared by a two-step melt polycondensation of dimethyl-4,4′-biphenyldicarboxylate with triethylene glycol using titanium(IV) isopropoxide as catalyst. Afterward, PTEB-OH polymer has been converted into a methacrylate end-functionalized polymer, named PTEB-MMA, by esterification reaction of the OH end groups with methacrylic acid using the N,N-dicyclohexylcarbodiimide/aminopyridine method. The general procedure is shown in Figure 1. C


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Figure 3. 1H NMR (300 MHz, CDCl3) spectra and peak assignment for (a) hydroxy-terminated PTEB polymer (PTEB-OH) and (b) methacrylate end-functionalized PTEB polymer (PTEB-MMA). The magnified spectrum around δ = 5.8 ppm in the inset shows the resonances of the methacrylate end groups.

In order to establish the conditions, cure temperature and cure time, for the preparation of the PTEB network from PTEB-MMA polymers, parallel plate rheology isothermal experiments were carried out at different temperatures (110, 120, and 130 °C). Figure 4 shows the evolution of storage moduli (G′), loss moduli (G″), and loss tangent (tan δ) during the thermal curing of PTEB-MMA. The polymer at these temperatures is molten, and for that reason G″ is higher than G′ at the beginning of the experiments and then tan δ > 1 in all cases. However, both moduli diminish with the temperature increase, as expected. After a period of time (induction time), quite short in all cases, both G′ and G″ rise rapidly, mainly G′. Therefore, tan δ values fall, indicating a restriction in molecular mobility as a result of the formation of a chemically cross-linked network. Finally, the tan δ value reaches a plateau that depends on the cure temperature. In the case of 130 °C this value is >1 because during the cure process G″ is always higher than G′, so that the viscous properties are dominant over elastic properties. This fact can be explained by two reasons: First, the actual temperature is always higher than the isotropization temperature, therefore promoting the viscous properties. Second, the cross-link density in this system (essential to enhance the elastic properties) cannot be very high because only the chain ends can be cross-linked. Nevertheless, values of tan δ < 1 are reached at lower temperatures, 120 and 110 °C, mainly due to the lower molecular mobility of the polymer at those temperatures, rather than an important variation of the degree of curing, as is discussed below.

= 53 ppm indicates the absence of terminal COOCH3 groups at both ends of the polymer chain. It is to be noted that the end groups ∼CH2OH cannot be quantified by integration of the signals in the 1H NMR spectrum because the corresponding methylene protons appear overlapped with the methylenes CH2O of the backbone (see Figure 3a). Thus, the equivalents of OH groups in PTEB-OH were estimated based on the SEC Mn value. Regarding the functionalization of PTEB-OH to give PTEBMMA polymer, when the reaction was carried out following the standard stoichiometric conditions, a partial esterification of the OH end groups was only achieved. However, the complete esterification was finally attained by employing a large excess of reactants. The successful and quantitative functionalization of PTEB-OH can be easily detected by NMR. The signals i, g, and h, associated with the chain ends ∼CH2OCH2CH2OH, are not detected in the 13C NMR spectrum (see Figure 2b). Furthermore, two new signals of low intensity are distinguished at δ = 18.5 and 126.8 ppm (not shown in the amplified Figure 2), which can be associated with the methyl methacrylate end groups (carbons k and j, respectively). The presence of methyl methacrylate end groups in PTEBMMA polymer was definitively confirmed by 1H NMR. In the olefinic range at 5.40−6.30 ppm the acrylate protons j are clearly observed (see inset in Figure 3b). Additionally, the protons of the methyl k appear as a singlet at δ = 1.92 ppm, and a triplet corresponding to the methylene neighboring the methyl methacrylate group is also observed at δ = 4.30 ppm. D


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the angular frequency. Since the gel point is an inherent material constant, it should not depend on experimental conditions, so that the crossover between G′ and G″ is not a good criterion to determine the gel point. Chambon and Winter23 first revealed that the G′ and G″ curves, in logarithmic scales, are parallel in a wide range of angular frequencies at the gel point, and therefore tan δ is independent of the angular frequency at that point. Consequently, a “frequency crossover point” appears in a determined time when monitoring tan δ at different frequencies, being this point the best criteria to determine the gel point. Figure 4 shows that the gel point for cure at 130 °C is at around 6 min and a little bit longer when curing at 120 °C (about 7 min). On the contrary, the frequency crossover point is not observed when curing at 110 °C. It seems that this temperature, approaching that of the isotropization temperature, reduces the chain mobility and makes difficult the crosslink reaction; thus, probably the cross-link density was lower at this temperature and hinders the presence of the frequency crossover point. The complex viscosities, shown in Figure 5,

Figure 5. Evolution of the complex viscosity as a function of the crosslinking reaction time at different temperatures (frequency = 10 Hz).

corroborate this fact, since the viscosity drop from 110 to 120 °C is much more important than that from 120 to 130 °C. Later, as expected, the increase of the viscosity follows a similar behavior than G′ and G″, reaching higher viscosity values at lower temperature but slower curing rate. These rheological results must be considered for the preparation of the PETB network from the precursor PTEBMMA film. In fact, the cross-linking of PTEB-MMA at 130 °C is quite fast, showing a very short induction time (less than 1 min). Thus, in order to prevent cross-linking during the previous compression molding of the precursor PTEB-MMA polymer to form a sheet, the temperature of 120 °C and short time (2 min) were finally chosen. Lower temperatures seem to be favorable in principle since the beginning of the cross-link reaction would be delayed, but the resulting higher viscosity will hinder the void elimination. Therefore, temperatures of 120 or 130 °C were found to be the most appropriate for the preparation of the PTEB network, leading to excellent sheets with a very good cross-linking yield, and prepared in relatively short curing times. The liquid crystalline behavior of PTEB has already been analyzed by our group in several works.13,24−26 On cooling

Figure 4. Evolution of storage moduli (G′), loss moduli (G″), and loss tangent (tan δ) with time during the thermal curing of PTEB-MMA at different temperatures.

One of the main transformations taking place during the formation of a polymer network is the gelation. This process corresponds to the incipient development of an infinite network, characterized by the divergence of the mass average, Mw, and the radius of gyration and by formation of an insoluble gel. Rheological monitoring of network formation has been largely applied in the evaluation of the gel point. The G′−G″ crossover (tan δ = 1) was first identified as the gel point,22 since the elastic properties (predominant in solid state) are higher than viscous properties (predominant in liquid state). The problem is that the gel point defined in such way increases with E


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polymer, which suggests a smaller degree of liquid crystallinity and/or less ordered smectic structure in the network. It is also interesting to notice, as observed in the inset in Figure 6c, that the DSC endotherm characterizing the transition SmC−SmA seems to be also present in the network, although now with a much lower intensity. Real-time variabletemperature diffraction experiments (similar to those already reported26 for PTEB) are needed in the network to ascertain this aspect. The room-temperature X-ray diffractograms for samples of PTEB and its network both annealed 4 months at room temperature (4mRT) and freshly cooled from the isotropic melt are presented in Figure 7. It is well evident that annealed

from the isotropic melt, PTEB develops an orthogonal SmA mesophase which is later transformed into a tilted SmC mesophase. The SmA phase is stable in a rather wide temperature interval, and the transformation into the SmC phase occurs at temperatures close to the glass transition, so that not very high tilting angles are attained.26 A reverse phase sequence (SmC−SmA−I) is observed in the subsequent heating of the polymer. It is also known that the liquid crystalline state of PTEB is stable for a considerable time although the transformation into a more ordered phase is produced when the polymer is annealed at temperatures above its glass transition.13,25,27 Therefore, it is pertinent to investigate the ability of crosslinked PTEB to crystallize. The DSC curves for the heating of PTEB and its network, after annealing at room temperature for 4 months, are shown in Figure 6a. The curve corresponding to

Figure 7. Room temperature X-ray diffractograms for samples of PTEB and its network both annealed 4 months (4mRT) and 4 days (4dRT) at room temperature and freshly cooled from the isotropic melt. Figure 6. DSC curves (scanning rate 20 °C/min) for PTEB and its network: (a) first melting (f1) of samples annealed 4 months at room temperature; (b) cooling from the isotropic melt; (c) second melting (f2), showing in the inset the amplified region for the transition SmC− SmA.

linear PTEB shows well-defined crystalline diffractions on top a relatively important amorphous-like component, indicating a considerable, but not very high, crystallinity. The annealed sample of the network also shows crystalline diffractions, although now the degree of crystallinity appears to be rather small. We have also performed several experiments of annealing at shorter times. For linear PTEB, nothing is observed below 2 days of annealing, but after 4 days a considerable crystallization occurred, as observed in the second from the top diffractogram in Figure 7. On the contrary, the corresponding diffractogram for the network shows no clear indication of crystallization even after 4 days of annealing. Regarding the freshly cooled specimens, the two diffractograms are much more similar, with an amorphous-like wide peak at higher angles (higher s values) and a narrow peak, and its second order, at lower angles, indicating a layered structure. These features are characteristic of a low-ordered smectic mesophase, with a well-defined smectic layer periodicity but with absence of lateral order within the layers. It is evident, however, that the layer peak is less intense in the network. These diffraction results confirm the conclusions extracted above from the DSC experiments: the network retains the liquid crystalline ability (with a slightly smaller order than that of the linear polymer), and it is even able to crystallize in a rather small degree after long-time annealing.

PTEB shows a typical melting−recrystallization−melting behavior, indicating, most probably that two different crystal polymorphs are involved. The total enthalpy is relatively high: 44.0 J/g. Regarding the network, a similar behavior seems to be obtained, except for the facts that the recrystallization is not observed and the total enthalpy involved is considerably smaller: only 14.4 J/g. It seems, therefore, that the network is also able to crystallize (although in a much lower degree than the linear polymer). The X-ray results will confirm this aspect (see below). Figures 6b and 6c show the DSC curves for PTEB and its network on cooling from the isotropic melt and the subsequent second melting, respectively. It can be observed, first, that the cross-linking leads to a rise of the glass transition temperature of 5 °C: the network shows Tg = 16 °C, while for the linear polymer Tg = 11 °C. On the other hand, the temperatures of mesophase formation and of mesophase isotropization of the network (TLC = 78 °C and Ti = 98 °C) are slightly lower than those observed for the linear polymer (82 and 105 °C, respectively). The corresponding enthalpies are also somewhat lower, 8.0 J/g for the network vs 10.8 J/g for the linear F


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Macromolecules The pure crystal profiles of the annealed samples can be easily obtained by subtracting the corresponding noncrystalline component taken as the diffractograms for the freshly cooled samples. Those profiles are shown in Figure 8 and indicate that

Figure 9. 3D stress−strain−temperature cyclic thermomechanical experiments for the PTEB network, showing a characteristic shapememory effect with the four typical stages: (a) initial deformation at T > Ti; (b) cooling to T < Ti at constant deformation; (c) contraction at constant temperature down to zero deformation in order to fix the transient shape; (d) recovering of the permanent shape by heating to T > Ti; and then cycling repeat three more times.

The plot of strain/temperature versus time corresponding to these four shape-memory cycles is shown in Figure 10.

Figure 8. Pure crystalline profiles for samples of PTEB and its network annealed 4 months (4mRT) and for PTEB annealed 4 days (4dRT) at room temperature.

the apparent degrees of crystallinity after 4 months of annealing are 40 ± 5 and 15 ± 5 for annealed linear PTEB and its network, respectively. The pure crystal profile for linear PTEB annealed for 4 days is also shown in Figure 8, in this case leading to a degree of crystallinity of 34 ± 5. The considerably lower crystallinity is perfectly expected when considering the three-dimensional network obtained after cross-linking, hampering the possibility of well-ordered structures. Anyway, it has to be also taken into account that as mentioned above, the glass transition of the network has been raised to a value of 16 °C, which is very close to room temperature, so that the transport term for crystallization may play an important role for retarding the ordering process. Anyway, the most relevant issue is that the liquid-crystallinity is perfectly retained in the particular network of PTEB here studied, since it means that we are dealing with a liquidcrystalline elastomer, and probably the most interesting feature of an LCE is the possibility of displaying a shape-memory effect. In these systems, temperature is the external stimulus, so that when the isotropization temperature, Ti, of the mesophase is exceeded, the LCE shrinks, owing to its elastomeric structure, but then when the temperature is brought below Ti, the systems expands again in a reversible way, resulting on a shape-memory effect. This behavior is usually quantified by the well-known cyclic thermomechanical experiments.4,28−30 The present network of PTEB has been subjected to such thermomechanical experiments, and the results, expressed by a 3D stress−strain−temperature plot, are shown in Figure 9. A well characteristic shape-memory effect is observed, with the four typical stages: (a) initial deformation at T > Ti; (b) cooling to T < Ti at constant deformation; (c) contraction at constant temperature down to zero deformation in order to fix the transient shape; (d) recovering of the permanent shape by heating to T > Ti; and then cycling repeat three more times.

Figure 10. Strain/temperature versus time corresponding to the four shape-memory cycles represented in Figure 9 for the PTEB network.

It can be observed in Figures 9 and 10 that a perfect reproducibility is obtained after the first cycle (cycles 2, 3, and 4 are practically coincident). On the contrary, the first cycle appears to be slightly different. Evidently, this may a problem of reduced shape recovery. However, it has been suggested31 that the first cycle may be rather used for clearing the thermomechanical history of the specimen, and the shapememory characteristics are better obtained from cycles ≥2. Additionally, the discrepancy in the first cycle may be also due to a certain slippage of the sample from the clamp during the first DMA cycle experiment. The quantification of the shape-memory effect is typically made trough the determination of the two more characteristic shape-memory quantities: shape-fixity, Rf, and shape-recovery, Rr, ratios,29,31 which can be determined from the stress−strain− temperature plots. The corresponding values extracted from Figures 9 and 10 indicate that the shape fixity is between 99.1 and 99.3% in the four cycles, while the shape recovery decreases from 96 to 94% from cycle 1 to 4. On the other hand, if the first cycle is taken as preconditioning and the shape-memory quantities are recalculated, it has a negligible effect on the values of Rf, keeping G


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Macromolecules

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around 99.2%. On the contrary, the values of Rr increase considerably, changing now from 99.2% for cycle 2 to 98.1 for cycle 4. Anyway, it follows that the investigated network of PTEB shows a rather remarkable shape-memory behavior, with considerably good shape-memory quantities.

4. CONCLUSIONS A smectic LC elastomer has been prepared from a main-chain polybibenzoate, poly(triethylene glycol p,p′-bibenzoate) (PTEB), in which cross-linkable methacrylate groups were incorporated at the ends of the polymeric chains. The advantage of this methacrylate end-functionalized PTEB is that it can be thermally cross-linked in the bulk state without an initiator to provide a network in good yield. The curing temperatures have been optimized by means of rheological experiments, so that excellent sheets were obtained, with a superior cross-linking yield and prepared in relatively short curing times. From well-established cyclic thermomechanical experiments, the shape-memory behavior of the PTEB network has been determined, showing rather good shape recovery and shape fixity values. As mentioned in the Introduction, by simple variations in the spacer it is possible to prepare polybibenzoates exhibiting smectic mesophases with varying degrees of order and with thermal transitions in a wide temperature range, which may be even above 300 °C when using short methylenic spacers.13 Additionally, those parameters can be also tailored by copolymerization17,32−34 (as it has been shown also in other polymer systems35). For instance, transitions near room temperature are obtained34 by using triethylene glycol as spacer and different proportions of bibenzoate and naphthalate units (the latter ones being nonmesogenic). Therefore, semiflexible polybibenzoates might be interesting main-chain polymers for the preparation of smectic elastomers with tailored thermal transitions. And all these systems are amenable of generating shape-memory LCEs via incorporation of cross-linkable methacrylate groups, as it has been shown here in PTEB. In fact, we have already prepared some other networks based on different polybibenzoates (homo- and copolymers), and the preliminary results are rather encouraging.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (E.P.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We acknowledge the financial support of MINECO-Spain (Project MAT2013-47972-C2-1-P). J.P.F.-B. thanks the “Marie Curie” Amarout Europe Programme. The synchrotron experiments were performed at beamline BL11-NCD at ALBA Synchrotron Light Facility with the collaboration of ALBA staff.

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REFERENCES

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Article

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Preparation and Properties of a Main-Chain Smectic Liquid-Crystalline

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