Reversible Shape-Memory Effect in Cross-Linked Linear Poly(Îµ

Article pubs.acs.org/Macromolecules

Reversible Shape-Memory Effect in Cross-Linked Linear Poly(ε-caprolactone) under Stress and Stress-Free Conditions Oleksandr Dolynchuk,*,† Igor Kolesov,*,‡ Dieter Jehnichen,† Uta Reuter,† Hans-Joachim Radusch,§,∥ and Jens-Uwe Sommer†,⊥ †

Leibniz-Institut für Polymerforschung Dresden e.V., D-01069 Dresden, Germany Interdisciplinary Center for Transfer-oriented Research and §Center of Engineering Sciences, Martin Luther University Halle-Wittenberg, D-06099 Halle (Saale), Germany ∥ Polymer Service GmbH Merseburg, D-06217 Merseburg, Germany ⊥ Institute for Theoretical Physics, Technische Universität Dresden, D-01062 Dresden, Germany ‡

S Supporting Information *

ABSTRACT: The effect of cross-link density on the reversible shape-memory effect (SME) under constant load was systematically studied in cross-linked linear poly(ε-caprolactone) (PCL). A remarkable reversible SME under stress-free conditions was observed in PCL with the highest achieved cross-link density. Thermal properties as well as morphology, size, and orientation of the nanocrystalline structure formed in covalent networks of PCL under load were compared with those in PCL crystallized under stress-free conditions. As shown, the oriented growth of crystals is the origin of both the reversible SME under and without load. Furthermore, a significant rise of crystallinity and crystal thickness was detected in PCL crystallized under constant load. The fitting curves of the temperature-dependent strain as well as the quantities of crystallinity, type of crystalline structure, size, and orientation of the crystals got by modeling the reversible SME in PCL under stress well correspond to their values obtained experimentally.

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INTRODUCTION The ability of materials to move actively in response to external stimuli determines their outstanding potential in a huge variety of applications spanning from the space and aircraft engineering up to designing the “smart” structures acting on the nanoscale.1−20 The biomedical use of stimuli-responsive smart materials gets an indisputable significance due to the role they can play in maintaining human life and fighting against various threats like cancer, viruses, etc. Polymers and polymer-based materials revealing the reversible shape-memory effect (SME)14−20 have recently become very promising candidates for many biomedical applications, e.g., thermoresponsive sutures, catheters, etc., on account of widely adjustable properties and low manufacturing costs. Besides, the molecular mechanisms underlying the reversible actuation of polymers are of great fundamental interest, as they have not yet been fully understood. Historically, the first observation of the macroscopic reversible actuation of polymers in response to the temperature change was reported by Mandelkern et al.21,22 in 1958−1959 for fibrous polyethylene (PE) cross-linked by means of high energy ionizing radiation of different doses. As a result, highly cross-linked PE has demonstrated a remarkable elongation during crystallization and a sample contraction during melting amounted to roughly 25%, both detected in a free-standing state. On the contrary, un-cross-linked PE has shown a length © XXXX American Chemical Society

increase of about 5% in the vicinity of melting temperature (Tm) at heating, which can reliably be addressed to a volume expansion typical for polymers undergoing the first-order phase transition like melting. However, this challenging discovery was underestimated and got neither sufficient exploration nor further development up to the 2000s, when the interest to SME in polymers has been rising significantly. The renewed attempts to test cross-linked crystallizable polymer systems for the reversible thermally initiated shapememory (SM) behavior have started from the works of Mather and co-workers in 2008.23,24 The authors have reported a macroscopic crystallization-induced elongation at cooling followed by a contraction at heating in cross-linked poly(cyclooctene)/trans-polyoctenamer (PCO/TOR)23 and in double networks of PCL and polyhedral oligosilesquioxane (POSS);24 in both cases polymer materials were continuously loaded by a constant force during the thermomechanical cycles. The sketch in Figure 1 illustrates loading of a polymer network and emergence of the reversible actuation. This phenomenon known now as the two-way reversible SME under constant load was later observed and studied in other semicrystalline polymer Received: March 5, 2017 Revised: April 21, 2017

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DOI: 10.1021/acs.macromol.7b00481 Macromolecules XXXX, XXX, XXX−XXX

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and are crystallizable that serves as the necessary preconditions for observing the reversible SM actuation in polymer-based SM materials. Additionally, branched polymer systems have and may have a special focus in this connection.28,29,31,32,37−39 As seen from the brief review above, PCL attracts particular attention as a SM material and has a considerable potential in various biomedical applications not least because of its high cross-linkability, crystallinity, proximity of its crystallization and melting temperatures to the human body temperature as well as excellent biocompatible and biodegradable properties.40 However, a lack of knowledge on the manifestation and fundamentals of the reversible SM behavior of cross-linked linear PCL with different network chain length is obvious, as no one comprehensive study dealing with it is known up to now. Conventional chemical cross-linking of linear PCL can be used to create SM polymer networks with various cross-link densities in order to explore the manifestation and fundamentals of the reversible SME under load in this biodegradable polyester more deeply. Evidently, a comprehensive study of the two-way SME in crystallizable networks of PCL must necessarily include a thorough analysis of the mechanical and thermal properties of the materials along with insights into the nanocrystalline structure of PCL and its change before and after the reversible actuation. Besides, a suitable theoretical description and modeling both the thermomechanical behavior under load and the morphology of the crystals generated in PCL-based SM networks are of great importance as it may extend our understanding the basics of the reversible SME and may become a powerful tool for prediction and creating materials with predetermined properties. These tasks form the main goals of the present work, which in fact continues the experimental and theoretical investigation of the SME in cross-linked semicrystalline polymers started from linear high-density and short chain-branched PEs.28,29,31,37,41−44

Figure 1. Sketch of loading of covalent network (middle), which elongates during crystallization under constant force (right). Dashed red-blue arrow indicates reversibility of actuation at heating/cooling; the force is applied vertically.

networks on the basis of star-branched methacrylate- or hydroxyl-terminated PCL,25−27 linear high-density PE, shortchain branched ethylene−octene copolymers (EOCs),28,29 etc. Furthermore, the so-called two-way triple-shape memory effect, which appears as two successive strain steps in the course of nonisothermal crystallization and melting under external load, was found in cross-linked polymer composites prepared from crystallizable segments of PCL and poly(ω-pentadecalactone) (PPD)30 as well as in cross-linked binary polymer blends of TOR with high-density PE and EOCs and high-density PE with PCL.31 As concluded,25−31 the origin of the reversible actuation in all these cases is the oriented growth of the crystals in the predrawn network. Recently, Lendlein and co-workers have discovered the reversible bidirectional SME under stress-free conditions in cross-linked poly[ethylene-co-(vinyl acetate)] (cPEVA)32 and polymer networks based on PCL and PPD,33 oligo(εcaprolactone) and n-butyl acrylate.34 The reversible actuation under zero stress was also detected in the copolymers of PCL and poly(tetramethylene glycol) (PTMG)35 as well as in crosslinked poly(octylene adipate).36 The physical background of this phenomenon lies in the partial melting of the crystalline phase formed in the predeformed covalent network, which in turn causes the partial strain recovery. The recrystallization of such polymer networks manifests in the directed crystal growth leading to the actuation, which is reversible upon partial melting at heating, similarly to the reversible SME under constant load. Moreover, as an alternative to complex and expensive chemical synthesis, the blending of crystallizable polymers like PE and EOCs was successfully employed to create cross-linked polymer actuators acting under stress-free conditions.37 As shown, an additional advantage of the latter method is a possibility to adjust widely the mechanical and thermal properties of the produced SM blends. Thus, despite the network architectures and manufacturing strategies were obviously different all aforementioned polymer systems comprised a stable covalent network

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EXPERIMENTAL SECTION

Materials and Processing. PCL used in the present study is linear commercially available polyester with a trade name CAPA 6800 (Perstorp UK Ltd., Warrington, UK), which is derived from ε-caprolactone monomer via ring-opening polymerization using catalyst. 2,5-Dimethyl-2,5-di(tert-butylperoxy)hexane (DHBP) (United Initiators GmbH & Co. KG, Pullach, Germany) was used as a crosslinking agent. The PCL pellets containing DHBP with different mass ratio of components (Table 1) were prepared by melt mixing for 5 min at 140 °C in a laboratory kneader Brabender Plasti-Corder (Brabender GmbH & Co. KG, Duisburg, Germany). After mixing, the DHBPimpregnated PCL pellets were first compression molded to films with thickness of 0.5 mm by pressing under load at 140 °C and immediately afterward cross-linked at 190 °C. The schematic explanation of crosslinking mechanism is given in the Supporting Information (see pages S2 and S3). The cross-link density of the samples νc was evaluated by means of stress-relaxation experiments carried out for 3 h at a draw ratio of λ ≈ 2 and a temperature of 97 °C, which is much above the melting point of all PCL samples under study.

Table 1. Designations as Well as Some Physical and Molecular Parameters of Used PCL designation

melt-flow index (MFI)a (deg min−1) 190 °C/2.16 kg

T mb (°C)

density (kg m−3)

weight-average molar mass M̅ w (kg mol−1)

polydispersity M̅ w/M̅ n

DHBP content (phr)

cross-link density ν̅cc (mol m−3)

PCL

7.29

59

1145

120

1.74

1 2 3 4

26 126 174 212

a

Measured by test apparatus MI 21,6 (Göttfert). bMelting peak temperature obtained by DSC. cObtained by stress-relaxation tests. B


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Macromolecules The values of the terminal stress σ∞ (stress at the infinitely long time) were taken to calculate the cross-link density using the classical outcome of the Gaussian theory of rubber elasticity45 for affine network as follows:

νc =

σ∞ RT(λ − λ−2)

deformation amplitude of 0.1% enabled in situ measuring the temperature dependencies of both the macroscopic strain and the dynamical storage modulus E′. Thus, the thermal loops of storage modulus E′ as well as SM creep and recovery strain in the first cooling and second heating runs, respectively, were recorded simultaneously. Transmission Electron Microscopy (TEM). The electron microscopic study of two sets of PCL samplesslowly crystallized under stress-free conditions and drawn samples crystallized after cooling under load during the two-way SM experiment (see section Two-Way SM Behavior)was carried out on a LIBRA200 MC (Carl Zeiss Microscopy GmbH, Oberkochen, Germany). The samples were prepared using ultramicrotome EM UC6/FC6 (Leica Microsystems GmbH, Wetzlar, Germany). The initial stage of sample preparation was the manufacture of cross sections of both sets of specimens at −150 °C under a dry nitrogen atmosphere. Subsequently, the cross sections were stained in RuO4 vapor for 24 h and then degassed for further 24 h. Finally, the ultrathin sections were sliced at −150 °C under dry nitrogen atmosphere and transferred for image acquisition onto copper grids coated with carbon film. The place of sampling of PCL samples is schematically shown in Figure 2.

(1)

where R is the gas constant and T is absolute temperature. The received values of νc along with other relevant parameters of PCL are listed in Table 1. Two-Way SM Behavior. The SM study of PCL was carried out in tensile mode using a mechanical spectrometer measuring head Mark III (Rheometric Scientific Inc., Piscataway, NJ). The PCL samples with various cross-link densities shaped as shouldered test bars with cross-sectional area 2.0 × 0.5 mm2 were tested during the cyclic thermomechanical experiment at an initial clamps distance of 6 mm. The test bars were initially heated to the highest temperature of SM test Thigh = 100 °C, which is above the melting temperature of PCL, and then loaded by different constant nominal stress σ0N. The latter is defined as force divided by the initial cross-sectional area.46 The loading immediately entailed the initial strain εini. Subsequent cooling under constant load to the lowest temperature Tlow = 0 °C at an average rate of 2 K min−1 resulted in the nonisothermal creep εcr prior to crystallization as well as in the considerable crystallization-induced elongation of the samples by strain increment Δεini. Afterward, the loaded specimens were thermally equilibrated for 5 min at Tlow, reached the deformation εlow, and then were heated to the temperature Thigh at a rate of 2 K min−1. Because of melting of the crystalline phase, the specimens contracted by strain decrement Δεdec. After thermal equilibration for 5 min at Thigh the deformation amounted to εfin, and the thermomechanical cycle was repeated. The strain as a function of temperature was recorded during the thermomechanical experiment. As the temperature sensor is placed beside a sample, some difference between the measured temperature and its true value appears in each cooling/heating tensile test using a mechanical spectrometer. In order to overcome this inconsistency, the calibration experiments were carried out as described in our previous work.28,29 The received data were used to determine the true temperature of specimens during the SM investigation. Differential Scanning Calorimetry (DSC). Melting and crystallization behavior of PCL under study was explored by a power compensation DSC 7 equipped with the liquid nitrogen accessory CCA-7 (PerkinElmer LAS GmbH, Rodgau, Germany) for controlled cooling and heating at a rate of 10 and 20 K min−1, respectively. The DSC data were recorded for two sets of PCL samples, each set consisted of two PCL samples with the appreciably different cross-link density νc. The first set was taken from undeformed PCL films in their original permanent shape, whereas the second set was cut from the drawn PCL samples crystallized at slow cooling under nominal stress σ0N of 0.6 and 1 MPa. The drawn samples were sealed in 20 mL aluminum pans between two thin films of polytetrafluoroethylene (PTFE) in order to ensure free mobility of a specimen during heating and cooling. The undeformed PCL samples were sealed without PTFE films. The sample mass in each case was about 8 mg. The raw heat-flow rate data were corrected for the instrumental asymmetry and converted into the temperature dependencies of the apparent specific heat capacity cp(T). The measured cp(T) values along with the theoretical cp(T) for crystalline and amorphous PCL as well as the enthalpy of fusion ΔHf taken from the ATHAS database47 were used to calculate the enthalpy-based crystallinity as a function of temperature χc(T) on the basis of the two-phase model as described by Mathot et al.48 In Situ Dynamic Mechanical Thermal Analysis (DMTA) and SM Tests under Constant Load. DMTA of the PCL samples loaded with a static force was performed in tensile mode using a mechanical spectrometer with measuring head Mark III (Rheometric Scientific Inc., Piscataway, NJ). The specimens with the same shape and sizes as described in the section above were tested at the frequency of 1 Hz during the thermal loops at heating/cooling rate of 2 K min−1. The combination of a high static load of 0.6 and 1 MPa with low dynamical

Figure 2. Sketch of the exposure of the drawn PCL samples to X-ray beam during WAXS and SAXS as well as sampling for TEM images. Wide- and Small-Angle X-ray Scattering (WAXS and SAXS). WAXS and SAXS of the PCL samples crystallized under stressfree conditions and under load during the two-way SME (see section Two-Way SM Behavior) were carried out at the self-constructed 3-fold pinhole system with rotating anode (Rigaku Corporation, Tokyo, Japan) using Cu Kα radiation (λ0 = 0.15418 nm), monochromatized by primary confocal multilayer optic (Max-Flux Optics, now: Rigaku Corporation, Tokyo, Japan) and area detection system MarCCD (now: Rayonix, L.L.C., Evanston, USA). The samples were exposed to X-rays as shown in Figure 2. The discussion of the results (see Results and Discussion section) is based on 2D WAXS and 2D SAXS patterns. WAXS intensity plots I(q) and I(χ) (q = 2π/d is scattering vector, χ is azimuthal angle) were created by means of sectorial integration over ±5° to the orientation on equator and 360° integration of the most intensive reflections, respectively. The scattering curves of SAXS are displayed as I(q)q2 vs scattering vector q without background subtraction (not corrected) and I(χ) vs azimuthal angle χ.

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RESULTS AND DISCUSSION Two-Way SM Behavior. The development of the two-way SME in PCL with various cross-link densities νc as well as under different nominal stress σ0N during the first thermomechanical cycle is shown in Figure 3. Further cycles of the SM behavior of PCL under 1 MPa are presented in the Supporting Information (Figure S1). The strategy of evaluating the influence of νc on the two-way SM behavior of PCL was to apply such a stress σ0N that leads to almost the same initial elongation εini. As seen from Figure 3a, the values of εini amounted to about 90−100% for all samples. However, being initially stretched to C


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Figure 3. Effect of cross-link density (a) and loading (b) on the two-way SME in PCL.

Table 2. Parameters of the Two-Way SME in PCL with Various Cross-Link Densities νc and under Different Load σ0N designation

νc (mol m−3)

σ0N (MPa)

Tcsw (°C)

Thsw (°C)

εini (%)

εcr (%)

Δεinc (%)

Δεdec (%)

γeff

γrec (%)

Wsp (kJ m−3)

Rr

PCL-26 PCL-126 PCL-174 PCL-212

26 126 174 212

0.2 0.6 0.8 0.3 0.6 0.7 1.0

35.6 41.1 39.8 26.1 34.2 43.9 46.0

52.5 57.2 54.2 50.3 52.1 53.9 57.1

90.0 87.8 99.4 13.5 33.2 63.6 94.3

25.6 25.1 32.9 2.4 8.4 12.5 25.0

27.7 35.5 50.2 9.3 20.2 39.6 65.2

29.0 40.1 61.0 12.0 24.1 39.7 65.1

0.31 0.40 0.51 0.69 0.61 0.62 0.69

104.6 112.9 121.5 129.0 119.3 100.4 99.8

58 241 488 36 145 278 651

0.77 0.89 0.90 0.95 0.94 0.89 0.92

entire SM test in the present study, the strain recovery ratio Rr also reflects the energy storage efficiency under load, similarly to that defined by Heuwers et al.51 Finally, the aforementioned characteristics were calculated for all curves of the temperaturedependent strain in Figure 3 and are listed in Table 2. The received results allow concluding that switching temperatures Tcsw and Thsw strongly correlate with crystallization Tc and melting temperatures Tm (see sections Crystallization and Melting Behavior as well as Viscoelastic Behavior and SME under Constant Load) and rise in general for PCL samples with increasing cross-link density νc when they are stretched almost to the same εini. First, a strict correlation between Tcsw/Thsw and Tc/Tm was also found for linear and short-chain branched PEs28,29 as an evidence that crystallization and melting of predrawn covalent networks are mainly responsible for the two-way SME. As seen, the present findings support this conclusion for PCL as well. Second, in spite of the found tendency for increasing Tcsw and Thsw with the increase of νc, PCL-126 under 0.6 MPa stands out from the trend because of having higher Thsw as compared to PCL-174. Indeed, such behavior seems to be nontrivial, as it deals with two opposite influences. On the one hand, cross-linking entails spatial hindrances on the formation of crystals and therefore reduces the crystal thickness. According to the Gibbs−Thomson equation derived under assumption of unaltered entropy,46 thinner crystals melt at lower temperature. Thus, higher cross-link density of PCL samples obviously results in their lower Tc and Tm. On the other hand, as reported previously,52,53 application of a load and deformation of polymer chains causes the acceleration of nucleation and crystal thickening. Hence, the higher cross-link density is, i.e., the shorter network chain is, the higher effect has deformation on the chain alignment and, consequently, on the nucleation and crystal thickness, which in turn directly affects Tc and Tm. Taking into account all these factors, one may assume that PCL-126 experiences stronger influence of deformation on the

the same extent, the PCL specimens under study are characterized by different magnitudes of the nonisothermal entropy−elastic creep prior to crystallization εcr as well as strain increment Δεinc and strain decrement Δεdec during crystallization and melting, respectively, which are exemplified for PCL under 1 MPa in Figure 3b. The strain recovery ratio Rr = (εlow − εfin)/(εlow − εini) was evaluated as well. Furthermore, using the previously introduced approach,29 the two-way SM effectiveness factor γeff and the two-way SM recovery factor γrec were estimated as follows:

γeff =

Δεinc εini

γrec =

Δεdec × 100% Δεinc

(2a)

(2b)

(Tcsw)

Besides, the switching temperatures at cooling and heating (Thsw) were obtained as temperatures corresponding to the highest strain rate peaks on the basis of plots −(dε/dt) vs T (Figure S2). In addition, as the samples contract against the applied constant force F during melting the crystalline phase at heating, it is worth calculating the specific work done in the course of the SM recovery, which under assumption of sample incompressibility (constant volume V) yields Wsp =

F Δldec = σ0NΔεdec V

(3)

where Δldec is the sample contraction during melting. In contrast to other valuable approaches used to evaluate the specific work during the entire heating run of the constrained one-way SM test,49−51 Wsp defined by eq 3 estimates a pure contribution arising only from the melting-induced release of the entropy− elastic forces stored previously during the two-way SME. Nevertheless, since the applied force is constant during the D


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Macromolecules nucleation than PCL-26 but comparable with PCL-174, since the latter has almost the same Tcsw as PCL-126. However, PCL-126 has the chain length enough to form thicker crystals as compared to PCL-174 that results in the higher Thsw. Though PCL-212 has the shortest chain length, the initial deformation of about 95% apparently leads not only to the strongest nucleation effect among PCL under study, which is justified by the highest Tcsw, but also to the significant increase of the crystal thickness as it comes from Thsw value of PCL-212. The qualitative explanation presented here is also quantitatively confirmed by the study of the nanocrystalline structure of PCL discussed below (see sections WAXS and SAXS). As seen from Figure 3a and Table 2, both strain increment Δεinc and decrement Δεdec drastically increase with rising νc of PCL. It is remarkable that Δεinc and Δεdec of PCL-212 are more than 2 times higher as compared to those of PCL-26 although both samples were stretched to the same εini. Moreover, values of the two-way SM effectiveness factor γeff and strain recovery ratio Rr rise monotonically with increasing cross-link density as well. Thus, these facts clearly indicate the higher perfection as well as SM performance of covalent networks in PCL having higher νc magnitudes. Besides, the two-way SM recovery factor γrec of the curves in Figure 3a goes above 100% and increases with elevating crosslink density νc. The reason for such behavior may be assumed in forming the imperfect crystals of low stability during annealing the samples at the lowest temperature of the test Tlow. Obviously, the formation of these imperfect crystals is then directly associated with storing some additional part of deformation. Thereby, in the course of heating and melting of the crystalline phase the samples recovery exceeds the elongation stored in the range of crystallization temperature Tc. This effect becomes more pronounced with increasing νc or, in other words, with shortening the network chain length (Table 2). The quite similar behavior of γrec was also observed for short-chain branched PEs in our previous study.29 Namely, adding the short branches to PE shortens the effective length of its chains entering to the crystals that serves as the analogy of increasing cross-link density in PCL. Again, this qualitative interpretation is confirmed by the DSC traces at heating presented below (see section Crystallization and Melting Behavior). The specific work Wsp estimated by eq 3 considerably rises with increasing νc. The corresponding values of Wsp are plotted in Figure 4. As seen, though the applied stress changes quasilinearly vs cross-link density, the dependence of Wsp on νc seems to be of sigmoid type with a tendency to saturation. This actually indicates some limit of Wsp that can be done by crosslinked PCL during the two-way SM recovery when a sample is initially stretched to about 100%. Since the latter conclusion may be suggested a matter-of-course outcome, it is however important to have its evident experimental manifestation as well as quantitative estimation as depicted in Figure 4. Furthermore, Figure 3b and Table 2 illustrate the evolution of the two-way SM behavior of PCL-212 under increasing load σ0N spanned from 0.3 to 1 MPa. Both Tcsw and Thsw increase monotonically with elevating σ0N that clearly points out the strong influence of the initial deformation and, consequently, of the initial alignment of network chains on the nucleation and crystal growth in PCL, as discussed above. The same tendency for Thsw was also found in SM natural rubber by Heuwers et al.54 The considerable rise of the strain increment Δεinc and decrement Δεdec as well as of the specific work Wsp done by PCL-212 (Table 2) serves as an additional proof to this statement.

Figure 4. Dependencies of the applied stress and the specific work done during the two-way SM recovery on cross-link density of PCL.

Note that Wsp experienced a tremendous gain of almost 20 times when σ0N increased in roughly 3.3 times. At the same time, values of γeff fluctuate around 0.6−0.7, indicating apparently the limited productivity of the crystalline phase in PCL-212 to induce the elongation during crystallization. Moreover, the strain recovery ratio Rr decreases slightly with increasing σ0N, as the higher load obviously causes the higher plastic deformation of PCL-212. However, we would like to stress that (1) Rr is quite high for PCL-212 under all applied loads (Rr ≥ 0.89) and (2) the recoverability of deformation in PCL-212 increases with cycling as it is exemplified for PCL-212 under 1 MPa (see Figure S1). Finally, the factor γrec drops down from 129% to about 100% with rising values of σ0N. The reason for that evidently lies behind the fact that the sample elongation during crystallization occurs toward the σ0N, whereas the sample contraction during melting proceeds against the applied external force preventing the strain recovery. Thus, the higher stress is applied, the stronger it affects the sample contraction. Note that the same tendency was observed for cross-linked short-chain branched PEs.29 Crystallization and Melting Behavior. The apparent specific heat capacity vs temperature measured for undeformed PCL-126 and PCL-212 and for their stretched samples, which underwent nonisothermal crystallization under load during the two-way SM tests, are shown in Figures 5a,b to analyze the influence of deformation on the thermal properties of PCL. Figures 5c,d represent the curves of enthalpy-based crystallinity as a function of temperature calculated from the DSC traces. In addition, Table 3 lists Tc and Tm values estimated as peak temperatures of the corresponding DSC curves as well as crystallinity χc at room temperature (20 °C). First, the data in Figures 5a,b and Table 3 clearly prove that cross-linking of PCL results in lowering its Tc and Tm magnitudes, as discussed above. Second, application of deformation and initial orientation of network chains already in the amorphous phase inevitably affect crystallization in PCL under load so that both melting point Tm and crystallinity χc in the entire temperature region increase in the course of subsequent heating and melting. Indeed, as seen from Figures 5a,c and Table 3, Tm and χc at heating rise by about 7 °C and more than 9%, respectively, for both stretched PCL samples as compared to their undeformed counterparts. Note that cross-linked linear PE does not exhibit such a drastic increase of Tm and χc despite it is stretched to a much higher extent.28 Besides, the DSC traces of the preliminary stretched PCL-126 and PCL-212 in the first heating run demonstrate wide melting peaks of low intensity E


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Figure 5. Temperature-dependent apparent specific heat capacity (a, b) and enthalpy-based crystallinity (c, d) at heating and cooling (cooling runs preceded the second heating) of undeformed PCL-126 and PCL-212 and their stretched samples crystallized at cooling under load.

Table 3. DSC Peak Melting and Crystallization Temperatures as Well as Enthalpy-Based Crystallinity of Undeformed and Stretched Samples of PCL-126 and PCL-212 Crystallized under Load during Cooling Stage of Two-Way SM Test crystallinity χc at 20 °C (%)

melting temperature Tm (°C) samples

type

1st heating run

PCL-126

undeformed drawn undeformed

59.1

drawn

57

PCL-212

2nd heating run

crystallization temperature Tc (°C)

51.7

26.2 27.5 24.7 28.2 25.2 28.2

50.5

appearing at about 35−36 °C. Similarly to short-chain branched PEs,48,55−57 these low-temperature DSC peaks originate from melting of imperfect crystal population formed during annealing the samples at room temperature. In addition, the mechanical manifestation of the phenomenon has already been discussed above. Although preliminary stretched PCL samples were placed in-between two PTFE films, sample recovery during the first heating was probably incomplete, as the corresponding Tc magnitudes are slightly higher than Tc of undeformed PCL. Interestingly, in contrast to PCL-126, the DSC traces of PCL-212 at cooling are featured by the fractionated crystallization peaks (Figure 5b) that reflects multiple nucleation occurred at different temperatures. The reason for that may be assumed in a peculiar network topology, namely, in the anisotropy of distribution of the network chain length got by cross-linking under press. As soon as a polymer network is anisotropic, shorter

1st heating run

cooling run

2nd heating run

34.3 35.3 33.5

42.0

51.2

49.9

40.3

33.5

network chains are more oriented than the longer chains that accelerates formation of crystal precursors and, consequently, leads to crystallization at a higher Tc. Thus, higher degree of the chain orientation is achieved by the shortening chain length even without an application of any external force. However, the presented assumption needs a very thorough experimental verification, though it is consistent with the specific reversible SME in PCL-212 under stress-free conditions discussed below. Viscoelastic Behavior and SME under Constant Load. The temperature-dependent curves of storage modulus and macroscopic strain measured simultaneously under and without load for two PCL samples with different cross-link density are presented in Figure 6. As the sample shape changes drastically, in particular the cross-sectional area, during DMTA tests under constant load, the measured values of storage modulus E′meas were corrected under assumption of sample incompressibility ′ = Emeas ′ (1 + Δl/l0), where Δl and l0 are absolute as follows: Ecor F


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Figure 6. In situ measured temperature-dependent storage modulus Ecor ′ and strain ε of loaded and unconstrained PCL-126 (a) and PCL-212 (b). Strain curves of the unconstrained PCL were magnified 10 times for better resolution. Arrows indicate cooling and heating runs.

Figure 7. High-resolution TEM images of unconstrained PCL-126 and PCL-212 (a, c) as well as of their samples stretched during nonisothermal crystallization under load (b, d). White arrows in images (b, d) indicate the stretch direction.

deformation and initial length of a sample, respectively. The ′ vs temperature plots are thus corrected storage moduli Ecor shown in Figure 6. The received results allow concluding that the two-way SME under constant load in both PCL samples is accompanied by the characteristic stepwise change of their storage moduli. This serves as an explicit proof that both sample elongation and contraction during the two-way SM behavior of PCL originate

from crystallization and melting, respectively. Moreover, values of crystallization Tc and meting Tm temperatures as well as of ′ in entire temperature range increase storage modulus Ecor with application of load. A rise of Tc and Tm well correlate with the findings described in the previous section and support the explanations concerning nucleation and crystal growth presented there. Besides, the applied constant load and ′ values in the temperature range deformation increase the Ecor G


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Macromolecules before crystallization and after melting, i.e., in elastic state of PCL networks. Obviously, it happens due to the alignment of network chains upon deformation, as discussed above, which in turn increases the stiffness of the covalent network of PCL. In the meantime, the gain of E′cor in the crystallized PCL under load reflects the formation of crystalline texture (see section Crystal Morphology as well as WAXS and SAXS) and the influence of crystal orientation on the mechanical properties of PCL. While unconstrained PCL-126 demonstrates a typical sample contraction and expansion during crystallization and melting, respectively, PCL-212 reveals a remarkable two-way SM actuation characterized by Δεinc = 6.7% and Δεdec = 7.4% under stress-free conditions. Moreover, a similar reversible actuation has recently been observed by Anthamatten and co-workers58 in thiol−acrylated PCL prepolymers cross-linked in two stages: (1) initial partial cross-linking occurred at 60 °C with 4-(dimethylamino)pyridine as base catalyst; (2) second-stage cross-linking was carried out by UV irradiation of stretched PCL networks using dimethoxy-2-phenylacetophenone as photoinitiator. Thus, the results in Figure 6b evidence that the reversible actuation of free-standing PCL networks can also be achieved by less complex, cheaper, and faster cross-linking of linear PCL than it was used before.58 Note that the reversible SME of PCL-212 under stress-free conditions is accompanied by the stepwise change of its E′cor that unambiguously points out the predominant role of crystallization and melting during the phenomenon. Thereby, a thorough study of the nanocrystalline structure in PCL-212 is demanded and therefore appears below. It has to be emphasized that the anomalous elongation of PCL-212 under “zero” stress occurs in the vicinity of the human body temperature. Crystal Morphology. TEM images in Figure 7 allow discussing the morphology of crystals formed in unconstrained slowly crystallized PCL-126 and PCL-212 as well as in their specimens crystallized under 0.6 and 1 MPa, respectively, at cooling during the two-way SM tests (see section Two-Way SM Behavior). As seen, crystal morphology in both sets of the samples is represented by lamellae, i.e., by folded-chain crystals, with a conspicuous tendency to formation of lamellae stacks, especially in drawn PCL. Furthermore, the stretched PCL samples are featured by the distinct orientation of the crystalline texture; namely, the basal surface of lamellae in Figure 7b,d is on average perpendicular to the stretch direction. The lamellae thickness in drawn PCL-126 and PCL-212 is estimated to be roughly 7.5−8 nm. However, the crystalline texture in the unconstrained PCL specimens does not show any particular orientation to the applied force, although it is characterized by lamellae stacks (Figure 6a,c) comprising crystals parallel to each other within the stacks. Such a local orientation of crystal “domains” in unconstrained PCL may originate from the restricted chain mobility due to spatial hindrances imposed by the relatively high cross-link density of the PCL networks under study. Since TEM cannot shed light on the integral overall orientation of both crystalline texture and crystal unit cell in the PCL samples, X-ray diffraction study is then required. WAXS. 2D WAXS patterns in Figure 8 were recorded to evaluate the symmetry and orientation of the crystal unit cell in slowly crystallized unconstrained and stretched PCL-126 and PCL-212 crystallized under load during the two-way SM test. The equatorial scattering plots (see Figure S3) were acquired by radial sector integration over ±5° off the equator for patterns

Figure 8. WAXS patterns of unconstrained slowly crystallized (a, c) and stretched foils of PCL-126 and PCL-212 crystallized under load of 0.6 and 1 MPa (b, d), respectively. White arrows in images (b, d) indicate the stretch direction.

in Figure 8. According to the received scattering plots and crystallographic database,59 the symmetry of crystals in the PCL samples under study is orthorhombic. The peak position of the most intensive reflections from the (110), (111), and (200) crystallographic planes (Figure S3) were used to calculate the parameters a, b, and c of the unit cell on the basis of Bragg’s law, which are listed in Table 4 along with the orientation of the unit cell to the stretch direction Z (drawn PCL) or to the sample length (unconstrained PCL). As seen from Table 4, structural parameters of the unit cell do not change much with an application of load; only parameter c experiences a slight increase for the stretched PCL-126. The curves of intensity I vs azimuthal angle χ were obtained by 360° integration of WAXS images in Figure 8 over two small intervals of scattering angle 2θ: 20.9°−21.8° and 23.2°−24°. The first 2θ interval contains (110) and (004) reflections, whereas the second one includes the reflection (200). The received curves are presented in Figure 9. The values of full width at half-maximum (fwhm) were ascertained from the curves in Figure 9 and are listed in Table 4. As it follows from Figures 8b,d and 9b, the stretched PCL samples show very distinct orientation of their unit cell so that the c-axis, which lies along the crystal chain, is parallel to the stretch direction. Surprisingly, the unconstrained slowly crystallized PCL-126 and PCL-212 also possess a certain orientation of the unit cell with the c-axis being parallel to the sample length, i.e., to the largest sample dimension, but has much wider distribution of orientation as compared to drawn PCL what is seen from the corresponding fwhm values in Table 4. Note that the orientation of the c-axis in unconstrained PCL-212 coincides with the direction of shape changing observed in this sample (Figure 6b) during the specific reversible SME under stress-free conditions. However, the discussed orientation phenomenon is accompanied by a high amount of crystals oriented randomly, which make about 70% and 69% of the crystalline phase in unconstrained PCL-126 and PCL-212, respectively. The latter quantitative characteristics were obtained by integration of relative reflections in Figures 8a,c. H


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Macromolecules

Table 4. Structural Parameters a, b, and c of the Unit Cell and Its Orientation Relative to the Sample Length l (Unconstrained) or to the Stretch Direction Z (Stretched) As Determined for Unconstrained Slowly Crystallized and Stretched Crystallized under Load PCL-126 and PCL-212 samples

type

a (Å)

b (Å)

c (Å)

φa,Z/la (deg)

φb,Z/la (deg)

φc,Z/la (deg)

fwhm(110)b (deg)

PCL-126

unconstrained drawn at 0.6 MPa unconstrained drawn at 1 MPa

7.51 7.52 7.51 7.52

4.98 4.98 4.98 4.99

17.36 17.42 17.45 17.46

90 90 90 90

90 90 90 90

0 0 0 0

67.0 30.7 62.5 31.8

PCL-212 a

Determined from azimuthal position of peak intensity of (110), (111), and (200) reflections. bFull width at half-maximum of (110) reflection.

Figure 9. Azimuthal scattering curves of 2D WAXS patterns of unconstrained (a) and stretched (b) PCL in Figure 8 obtained by 360° integration over two intervals of angle 2θ: (1) 20.9°−21.8°, which contains reflections (110) and (004); (2) 23.2°−24° comprising the reflection (200).

So, only a minority of crystals in unconstrained PCL has preferred orientation. SAXS. 2D SAXS patterns are presented in Figure 10 to discuss size and orientation of the crystalline texture in unconstrained slowly crystallized PCL-126 and PCL-212 as well as in their stretched samples undergone crystallization under load during the two-way SM test. The graphs in Figure 11 illustrate the corresponding SAXS intensity Iq2 as a function of scattering vector q = 2π/d got by 360° integration and the intensity I vs azimuthal angle χ obtained at the maximum of the Iq2 curves by integration over a small Δq interval. The peak positions of the curves in Figure 11a allow calculating the long period L (repeating unit) of the alternating crystalline and amorphous layers in the covalent networks of PCL: L = La + Lc, where Lc is the thickness of the crystalline sublayer and La is the thickness of the amorphous sublayer. Under assumption that all crystalline and amorphous regions are set as sublayers in lamellar stacks, the crystal thickness Lc of PCL may be found as product of long period L and crystallinity χc determined from the corresponding DSC traces: Lc = Lχc. Indeed, the presented above morphological study (Figure 7) has revealed the strong tendency to formation of lamellae stacks, especially in the case of drawn PCL. Hence, the assumption on periodicity of amorphous and crystalline layers seems reasonable and experimentally justified. The received values of Lc are listed in Table 5. As well seen, the crystal thickness of both PCL-126 and PCL-212 experiences a drastic increase with an application of load and deformation. Although Lc values of PCL-126 are higher than that of PCL-212 on account of less cross-link points in PCL-126, stretching of PCL-212 results in about 1.3 nm rise of its Lc as compared to 1.1 nm increase of Lc for PCL-126. Thus, the initial deformation of about 90% stronger affects crystallization in PCL-212 than in PCL-126, as it was discussed in details in the section Two-Way SM Behavior.

Figure 10. 2D SAXS patterns of unconstrained slowly crystallized PCL-126 and PCL-212 (a, c) and their stretched foils crystallized under load of 0.6 and 1 MPa (b, d), respectively. White arrows in images (b, d) indicate the stretch direction.

The two-point reflections in Figure 10b,d along with the azimuthal curves in Figure 11b are evidence of a strong orientation of the alternating crystalline and amorphous layers in the stretched PCL with respect to the applied force. Thus, based on I


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Macromolecules

Figure 11. SAXS intensity Iq2 vs scattering vector q (a) as well as intensity I vs azimuthal angle χ (b) obtained for unconstrained slowly crystallized and stretched PCL-126 and PCL-212 crystallized under load of 0.6 and 1 MPa, respectively.

Table 5. Crystal Thickness as Well as Relative Orientation Degree of Crystalline Structure Formed in Unconstrained PCL-126 and PCL-212 and Their Stretched Foils Crystallized under Load during the Two-Way SM Test samples

type

long period L (nm)

crystallinity χc at 20 °C (DSC) (%)

crystal thickness Lc (nm)

fwhm (deg)

Θrel

PCL-126

unconstrained drawn at 0.6 MPa unconstrained drawn at 1 MPa

16.2 15.5 15.9 15.5

42.0 51.2 40.3 49.9

6.8 7.9 6.4 7.7

82.2 20.8 78.3 20.0

0.54 0.88 0.56 0.89

PCL-212

pure, though consistent, assumption with a strong intention to study the phenomenon thoroughly in the nearest future. Modeling the Two-Way SM Behavior. The theoretical description of the reversible SME in PCL networks under constant load follows in general the mean-filed approach derived previously and successfully applied to model the SM behavior of linear and branched PEs.28,29,43 Namely, the free energy change of a deformed polymer network undergoing crystallization is set to consist of the following terms: the free energy of transmitting statistical segments from the amorphous region to the crystals; the fold surface free energy of a crystal with f folds (or lateral surface free energy in the case of the extended-chain crystal, f = 0); the entropy change in the remaining amorphous segments. A schematic representation of loaded semicrystalline network chain is shown in Figure 12. The expression for the free energy change as a function of

the dependencies in Figure 11b and the morphological study presented in Figure 7b,d, it is concluded that lamellae in drawn PCL are normal to the direction of elongation. In addition, since the c-axis in both drawn PCL was found parallel to the stretch direction, as shown in the previous section, it is quite clear that the c-axis is orthogonal to the basal surfaces of lamellae in these samples as well. Thereby, no crystal stems obliquity or tilted arrangement of the crystal stems is observed in the stretched cross-linked PCL under study, in contrast to drawn linear PE as reported previously.28,29 Furthermore, the SAXS azimuthal dependencies acquired for unconstrained PCL (Figure 11b) also show the anisotropy of the crystalline texture orientation, namely their lamellae tend to be normal to the sample length. The fwhm values of the peak intensities in Figure 11b along with relative orientation degree given as Θrel = (180° − fwhm)/180° are enumerated in Table 5. As visible in Table 5 and Figure 11b, the unconstrained PCL-212 has a narrower distribution of the orientation of its crystalline texture in comparison with PCL-126. Thus, summarizing data on the mechanical (Figure 6b) and thermal behavior (Figure 5) as well as on the nanocrystalline structure of unconstrained PCL-212, it can be accurately concluded that the origin of the specific reversible SM behavior of PCL-212 under stress-free conditions is the oriented growth of crystals. The reason for the latter, as it was first assumed in the section Crystallization and Melting Behavior, might be the anisotropy of distribution of the network chain length got during cross-linking under press. Namely, we assume that pressing the melt of entangled PCL chains causes a primary deformation of such a physical network, which becomes anisotropic. The subsequent immediate peroxide cross-linking fixes the anisotropy of the network chain length in PCL. So, the higher cross-linking is the more anisotropic the covalent network becomes. Therefore, the reversible SME under stressfree conditions is observed in PCL with the highest cross-link density so far. At the same time, we do not want to speculate too much and would like to leave the presented explanation as a

Figure 12. Sketch of loaded semicrystalline network chain linking two neighboring cross-links (red points); white imagined point divides the chain into amorphous and crystallized subchains. J


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Macromolecules temperature T and draw ratio λ for various number of folds f is given as ⎧ ⎛ ⎪ Nχc Ul , (f = 0) T ⎞ ΔFf (T , λ) = ⎨ − Nχc ΔHμ⎜1 − 0 ⎟ ⎪ Tm ⎠ ⎝ ⎩ fUf , (f > 0) ⎛ 1 ⎤⎡ 2 1 ⎞⎟ RT ⎡ + λ ⎢⎣1 − ⎥⎢λ − 2φδ ⎝⎜1 + 2θ 20N ⎠ Nθ ⎦⎢⎣ +

⎤ 3φ4 2 3φ 2 3RT ⎡ 4 4 8 1 + − 3θ ⎥ + λ + λ + + ⎢ 3 2 3 3 λ2 λ N ⎦ 20Nθ ⎣ N

6φ 2 2 4φ 2 1 8 − φδ(λ 3 + 1) λ + 3 N N λ ⎞⎤ 4 3 2 2⎛ ⎟ − φ δλ −θ ⎜5 − ⎥ ⎝ N Nθ ⎠⎦ +

l = [f 2 a0 2 + βζ 2b0 2]1/2 ≜ φb0 ; ⎧ 0, f = 2k β=⎨ ⎩1, f = 2k + 1

ζ=

Figure 13. Free energy of crystallization ΔFf (T,λ) as a function of temperature and draw ratio calculated for PCL-212 in the case of different numbers of crystal folds f. Blue dashed arrow indicates the crystallization path at cooling.

(4a)

Nχc f+1

;

thermodynamically for lower deformation (1 < λ ≲ 1.25). Obviously, this conclusion falls outside the experimental findings about the orientation of the nanocrystalline structure in unconstrained PCL-212 (sections WAXS and SAXS). However, it should be noted that the theory was derived for monodisperse network chains and does not take into account a probable anisotropicity of distribution of the network chain length in unconstrained PCL-212. At the same time, as soon as the external force is applied, deformation of the network plays a more important role in orientation of chains than their poly dispersity. Thereby, the used analytical approach becomes reasonable and appropriate for the description of the nonisothermal crystallization of the polymer network under load. Based on the results in Figure 13 and on the equations derived in our previous publications,42,43 the temperature-dependent strain measured during the two-way SME in PCL-212 under load was fitted using the Levenberg−Marquardt algorithm and presented in Figure 14. Note that the final analytical expression

⎪

⎪

θ = (1 − χc );

δ=

⎛ 6 ⎞1/2 ⎜ ⎟ ⎝ πN ⎠

(4b)

(4c)

where N is the number of statistical segments, Uf the fold surface free energy per fold, Ul the lateral surface free energy per segment, ΔHμ the enthalpy of fusion per segment, T0m the equilibrium melting temperature,and b0 and a0 the length and thickness of a segment, respectively. It has to be emphasized that the theory assumes the crystal vector l (Figure 12) to be always toward the applied external force F. Therefore, the number of folds f also determines the orientation of the crystallite relative to the applied force as follows: (1) odd f results in perpendicular orientation of crystal stems to F; (2) even f (including the extended-chain morphology f = 0) implies that the crystal stems are parallel to F or make an acute angle with F, which is precisely governed by the numbers of statistical segments N and folds f (eq 4b). Thus, the free energy change ΔFf (T,λ) in Figure 13 was calculated for different number of crystal folds f formed in PCL-212 using the following values of material constants in eq 4a: N = 46, Ul = 1364.9 J mol−1,60 Uf = 8869.6 J mol−1,61,62 ΔHμ = 17 877.5 J mol−1,47,61,62 T0m = 372 K,61,62 a0 = 4.313 Å,60 b0 = 8.624 Å.63 The number N was set to correspond directly to the cross-link density of PCL-212 νc = 212 mol m−3. ΔFf (T,λ) values for f > 2 are higher than those in Figure 13, and therefore they are not displayed. As it comes from Figure 13, the extended-chain crystallites (f = 0) (stems are parallel to F) possess a lower free energy at the onset of crystallization at the given deformation range. Further cooling at deformation 1.25 ≲ λ < 3 entails the formation of the folded-chain crystals, i.e., lamellae, with f = 2 and stems oriented nearly parallel to the stretch direction. Thus, the crystal c-axis is oriented toward the applied force in the whole temperature range of crystallization. These analytical outcomes are completely consistent with the results of TEM, WAXS, and SAXS in loaded PCL-212 presented above. According to Figure 13, the formation of the crystallites having one fold (f = 1) with the c-axis normal to F is the most favored

Figure 14. Experimental (bold) and fitting (thin yellow) curves of the two-way SME in PCL-212 under various loads.

for strain as a function of temperature includes not only the influence of crystallizing covalent network, though it is the most significant one, but also the thermal contraction/expansion of a sample and the viscoelastic deformation of entangled macromolecules.43 The latter is characterized by such important fitting parameters as the density of entangled slipped macromolecules K


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Macromolecules

Table 6. Influence of Load on Relevant Fitting and Material Parameters Obtained by Modeling the Two-Way Reversible SME in PCL-212 σ0N (MPa)

νe (mol·m−3)

Vh (nm3)

Vm (102 nm3)

U (kJ mol−1)

f

χcf (%)

Lcal c (nm)

0.3 0.6 0.7 1.0

50.4 50.9 52.1 54.5

4.1 4.1 4.1 4.1

3.8 3.8 3.9 4.0

92.8 96.8 102.2 108.1

2 2 2 2

49.6 53.6 56.8 58.1

6.5 7.1 7.5 7.7

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νe, the activation volume of viscous flow Vh, the volume of the flowing molecule Vm,42,43,64 and the activation energy of the viscous flow U. These and other important material and fitting parameters are listed in Table 6. As visible in Figure 14, the present mean-field approach successfully describes the thermomechanical two-way SM behavior of PCL-212 under various static loads. In addition, the fitting values χcf, which represent a final degree of crystallinity when crystallization ceases,37 increase with rising load and, consequently, increasing deformation of the covalent network. This trend is in accordance with the DSC results (section Crystallization and Melting Behavior) showed that the application of load to PCL-212 causes a drastic gain of its crystallinity χc. However, the magnitude χcf in the case of PCL-212 under 1 MPa is somewhat higher than its experimentally determined counterpart χc (Table 3). Probably, this may point to the heterogeneous distribution of cross-link points in PCL-212 so that some local areas of the real network are less mobile and less active in forming the large lamellar structure on account of too short chains, whereas longer chains mainly form lamellae stacks and have higher crystallinity than the entire network on average. Moreover, the received fitting parameters allow calculating the crystal thickness or stem length as follows: Lccal =

Nχcf b. f+1 0

CONCLUSIONS AND OUTLOOK

The evolution of the two-way reversible SME as a function of cross-link density and load was thoroughly studied in the networks of linear PCL with a particular focus on the nanocrystalline structure formed during nonisothermal crystallization of predrawn networks under load. It was ascertained that the two-way SM performance of cross-linked PCL quantified by the two-way SM effectiveness factor γeff strongly depends on the cross-link density νc and monotonically increases with rising νc as γeff gains from about 0.3 to 0.7. Although PCL-212 demonstrates a drastic rise of the strain increment Δεinc and decrement Δεdec with elevating load σ0N, the factor γeff only fluctuates around 0.6−0.7 and indicates the restricted two-way SM performance of the crystalline phase in PCL-212. In addition, the specific work Wsp done by the PCL samples against the external force during sample contraction at melting increases tremendously with increasing both νc and σ0N and reaches a large value of 0.65 MJ m−3 for PCL-212 under 1 MPa. In situ measurements of storage modulus and macroscopic strain of PCL networks under load clearly proved that the twoway reversible SME in PCL arises from crystallization and melting of the covalent network. Moreover, PCL-212 revealed a striking reversible actuation under stress-free conditions featured by about 7% of the shape change during crystallization and melting. The thermal behavior of PCL disclosed a significant increase of both melting temperature Tm (≈7 °C) and crystallinity χc (≈9%) in the stretched PCL-126 and PCL-212 crystallized under load as compared to their unconstrained counterparts. Besides, crystallization of PCL-212 is characterized by the fractionated peaks of the specific heat capacity that points out to multiple nucleation at different temperatures. TEM investigation showed that the morphology of crystals formed in the stretched and unconstrained PCL-126 and PCL-212 is represented by lamellae stacks oriented strongly in the case of the stretched samples so that the basal surface of lamellae is normal to the stretch direction Z. The latter outcome was completely confirmed by SAXS and WAXS, which in addition allowed concluding the c-axis lying along Z in the drawn PCL. As found, the crystal thickness increases with the application of load and subsequent deformation by about 1.1 and 1.3 nm for PCL-126 and PCL-212, respectively. Furthermore, the nanocrystalline structure in the unconstrained PCL-212 were found to have qualitatively the same anisotropic orientation relative to the direction of the specific actuation as in the case of drawn PCL-212, though the amount of such oriented crystals is only about 30%. Thus, it is concluded that the origin of the specific reversible actuation of PCL-212 under stress-free conditions is the oriented growth of crystals. It is assumed that the physical background of such a nontrivial phenomenon might be the anisotropy of distribution of the network chain length. However, a thorough experimental study is to be carried out to either confirm or refute this raw assumption. Finally, the two-way SME in PCL-212 under various loads was modeled on the basis of the recently developed

As f = 2 for all fitting curves in Figure 14, Lcal c

magnitudes follow the same tendency to increase with rising load as in the case of χcf (Table 6). The Lcal c value calculated for PCL-212 under 1 MPa surprisingly well corresponds to its Lc obtained experimentally on the basis of DSC and SAXS. Lcal c of PCL-212 under 0.3 MPa is slightly higher than Lc of unconstrained PCL-212 as it is expected due to higher stretching, but both values are quite close. Finally, although Vh and Vm remain almost unaltered, the activation energy U gain with increasing external load. Apparently, the reason for that lies behind a drastic change of chain mobility during crystallization/ melting, which in turn affects the viscous flow of the entangled macromolecules. As Tc and Tm rise with increasing σ0N, slipping of the entangled macromolecules takes place at higher temperatures as well and requires higher activation energies. The same tendency was also found while modeling the reversible SM behavior of cross-linked high-density PE under various loads.29 Note that the material parameters used above along with Lcal c magnitudes listed in Table 6 can be applied to calculate the melting point Tm of PCL-212 under various loads based on the Gibbs−Thomson equation.62 Although the Gibbs−Thomson equation is a very rough approach to crystallization in a predrawn covalent network, it gives adequate evaluations for Tm as follows: Tm (0.3 MPa) = 49.9 °C, Tm (0.6 MPa) = 54 °C, Tm (0.7 MPa) = 56.4 °C, and Tm (1 MPa) = 57.5 °C. Thereby, one may also estimate the melting point of the thickest lamellar crystal (1-fold) formed in the PCL network with the chain length of 46 segments that yields Tm = 66.7 °C. The latter may indicate the highest possible Tm of lamellae formed in PCL-212, when higher load is applied. L


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(7) Leng, J.; Lu, H.; Liu, J.; Huang, W. M.; Du, S. Shape-memory polymers − A class of novel smart materials. MRS Bull. 2009, 34 (11), 848−855. (8) Leng, J.; Du, S. Shape Memory Polymers and Multifunctional Composites; CRC Press: Boca Raton, FL, 2010. (9) Lendlein, A.; Jiang, H. Y.; Junger, O.; Langer, R. Light-induced shape-memory polymers. Nature 2005, 434, 879−882. (10) Guskova, O.; Schünemann, C.; Eichhorn, K.-J.; Walzer, K.; Levichkova, M.; Grundmann, S.; Sommer, J.-U. Light absorption in organic thin films: The importance of oriented molecules. J. Phys. Chem. C 2013, 117 (33), 17285−17293. (11) Ivaneyko, D.; Toshchevikov, V.; Borin, D.; Saphiannikova, M.; Heinrich, G. Mechanical properties of magneto-sensitive elastomers in a homogeneous magnetic field: Theory and experiment. Macromol. Symp. 2014, 338 (1), 96−107. (12) Daniel, W. F. M.; Burdyńska, J.; Vatankhah-Varnoosfaderani, M.; Matyjaszewski, K.; Paturej, J.; Rubinstein, M.; Dobrynin, A. V.; Sheiko, S. S. Solvent-free, supersoft and superelastic bottlebrush melts and networks. Nat. Mater. 2015, 15, 183−189. (13) Paturej, J.; Sheiko, S. S.; Panyukov, S.; Rubinstein, M. Molecular structure of bottlebrush polymers in melts. Sci. Adv. 2016, 2 (11), e1601478. (14) Zhou, J.; Sheiko, S. S. Reversible shape-shifting in polymeric materials. J. Polym. Sci., Part B: Polym. Phys. 2016, 54 (14), 1365− 1380. (15) Li, Q.; Zhou, J.; Vatankhah-Varnoosfaderani, M.; Nykypanchuk, D.; Gang, O.; Sheiko, S. S. Advancing reversible shape memory by tuning the polymer network architecture. Macromolecules 2016, 49 (4), 1383−1391. (16) Pilate, F.; Toncheva, A.; Dubois, P.; Raquez, J.-M. Shapememory polymers for multiple applications in the materials world. Eur. Polym. J. 2016, 80, 268−294. (17) Wermter, H.; Finkelmann, H. Liquid crystalline elastomers as artificial muscles. e-Polym. 2001, 1 (1), 111−123. (18) Finkelmann, H.; Nishikawa, E.; Pereira, G. G.; Warner, M. A new opto-mechanical effect in solids. Phys. Rev. Lett. 2001, 87 (1), 015501. (19) Quitmann, D.; Gushterov, N.; Sadowski, G.; Katzenberg, F.; Tiller, J. C. Solvent-sensitive reversible stress-response of shape memory natural rubber. ACS Appl. Mater. Interfaces 2013, 5 (9), 3504−3507. (20) Quitmann, D.; Gushterov, N.; Sadowski, G.; Katzenberg, F.; Tiller, J. C. Environmental memory of polymer networks under stress. Adv. Mater. 2014, 26 (21), 3441−3444. (21) Mandelkern, L.; Roberts, D. E.; Diorio, A. F. Reversible contractile processes in fibrous macromolecules. J. Am. Chem. Soc. 1958, 80 (2), 500. (22) Mandelkern, L.; Roberts, D. E.; Diorio, A. F.; Posner, A. S. Dimensional Changes in Systems of Fibrous Macromolecules: Polyethylene. J. Am. Chem. Soc. 1959, 81 (16), 4148−4157. (23) Chung, T.; Romo-Uribe, A.; Mather, P. T. Two-way reversible shape-memory in a semicrystalline network. Macromolecules 2008, 41 (1), 184−192. (24) Lee, K. M.; Knight, P. T.; Chung, T.; Mather, P. T. Polycaprolactone−POSS chemical/physical double networks. Macromolecules 2008, 41 (13), 4730−4738. (25) Pandini, S.; Passera, S.; Messori, M.; Paderni, K.; Toselli, M.; Gianoncelli, A.; Bontempi, E.; Riccó, T. Two-way reversible shape memory behaviour of crosslinked poly(ε-caprolactone). Polymer 2012, 53 (9), 1915−1924. (26) Pandini, S.; Baldi, F.; Paderni, K.; Messori, M.; Toselli, M.; Pilati, F.; Gianoncelli, A.; Brisotto, M.; Bontempi, E.; Riccó, T. Oneway and two-way shape memory behaviour of semi-crystalline network based on sol-gel cross-linked poly(ε-caprolactone). Polymer 2013, 54 (16), 4253−4265. (27) Pandini, S.; Dioni, D.; Paderni, K.; Messori, M.; Toselli, M.; Bontempi, E.; Riccó, T. The two-way shape memory behaviour of crosslinked poly(ε-caprolactone) systems with largely varied network density. J. Intell. Mater. Syst. Struct. 2016, 27 (10), 1388−1403.

mean-field approach. The analytical results on the morphology and orientation of the nanocrystalline structure formed in the predrawn network of PCL-212 are completely in line with the corresponding experimental outcomes. The theory predicts that the reversible actuation of polymer network under constant load originates from the directed growth of crystals when crystal stems are parallel or nearly parallel to the applied force. The fitting curves of strain vs temperature, which were received by modeling the SM behavior of PCL-212 under load, show excellent coincidence with the experimentally obtained findings. Nevertheless, the theory in its current form is not able to explain and describe the peculiar reversible actuation of the unconstrained PCL-212. A consistent theoretical description of this phenomenon is an attractive and challenging task, which is to be completed in the future.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.7b00481. Pages S2−S3: schematic explanation of cross-linking mechanism of PCL using peroxide; Figure S1: cycling of the two-way SME in PCL-212 under 1 MPa; Figure S2: the effect of cross-link density and loading on the kinetics of the two-way SME in PCL; Figure S3: WAXS equatorial scattering curves of unconstrained slowly crystallized as well as of stretched PCL-126 and PCL-212 crystallized under load during the two-way SME (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail [email protected], [email protected] (O.D.). *E-mail [email protected] (I.K.). ORCID

Oleksandr Dolynchuk: 0000-0002-5336-5068 Jens-Uwe Sommer: 0000-0001-8239-3570 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors are grateful to Dr. A. Jurjiu (Babes-Bolyai University, Cluj-Napoca, Romania) as well as to Dr. P. Formanek, Dr. J. Paturej, M. Koch, A. Checkervarty, and especially to Dr. O. Guskova (Leibniz-Institut für Polymerforschung Dresden e.V.) for helpful discussions of the results and their valuable recommendations.

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