Mechanical Properties of Isotactic 1, 2-Poly (E-3-methyl-1, 3

Article Cite This: Macromolecules XXXX, XXX, XXX−XXX

Mechanical Properties of Isotactic 1,2-Poly(E‑3-methyl-1,3pentadiene): An Example of Rubbery Elasticity below Glass Transition Temperature Claudio De Rosa,*,† Miriam Scoti,† Finizia Auriemma,† Anna Malafronte,† Chiara Santillo,† Giorgia Zanchin,‡,§ Ivana Pierro,‡,† Giuseppe Leone,‡ and Giovanni Ricci‡ †

Dipartimento di Scienze Chimiche, Università di Napoli Federico II, Complesso Monte S.Angelo, Via Cintia, I-80126 Napoli, Italy CNR-Istituto per lo Studio delle Macromolecole (ISMAC), Via A. Corti 12, I-20133 Milano, Italy § Dipartimento di Chimica, Università degli Studi di Milano, via C. Golgi 19, I-20133 Milano, Italy ‡

S Supporting Information *

ABSTRACT: A study of the mechanical behavior of isotactic 1,2-poly(E-3methyl-1,3-pentadiene) (iP3MPD) and of the structural and morphological transformations that take place upon stretching is presented. Amorphous samples do not show viscous flow but show high ductility, strain hardening, and elastomeric behavior. An unusual double yielding appears in the stress− strain curves that may be correlated to occurrence of events typical of yielding at low deformation followed by stress-induced crystallization of a mesomorphic form at higher deformation (200−300%). The yielding stress and an endothermic hysteresis effect at the glass transition simultaneously increase upon aging of the amorphous glass due to a local densification or formation of local order in the amorphous glass upon physical aging. The elasticity is accompanied by the crystallization of the mesophase during deformation and its melting upon relaxation. During successive cycles of tensile deformation and elastic recovery the yielding peak is still observed during every stretching step. The elasticity of iP3MPD is an unusual case of a polymer that shows rubbery elastomeric properties below the glass-transition temperature.

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INTRODUCTION Elastomers are generally amorphous polymers made by long and highly flexible chains with high molecular mass and glasstransition temperature much below room temperature.1 The flexible segments of the random coil conformation in the unstretched state assume extended conformation upon stretching, and the entropy-driven recoiling upon releasing the stress is responsible for the elastic recovery of size and shape of the whole material in the undeformed state.1,2 The existence of network knots prevents the material from flowing during application of the tensile stress. These amorphous rubbers may partially crystallize under stretching.1 The small number of small crystals acts as knots of the elastomeric network. When the applied tensile stress is removed, the crystals melt, providing a positive enthalpic contribution to the Gibbs free energy change involved during the elastic recovery of the material; therefore, the elasticity in these materials is merely of entropic nature due to the conformational changes experienced by the amorphous “tie chains”.1 In general, a high level of crystallinity may strongly reduce the elastic performances of these materials. However, many semicrystalline polymers show elastic properties in spite they present high crystallinity and high mechanical strength.3−19 Phase transitions may play a key role in the elasticity of these systems, resulting in materials where elasticity is not merely © XXXX American Chemical Society

entropic as in conventional elastomers but similar to superelasticity of shape-memory alloys that undergo martensitic phase transitions.6−19 An example of this type of rubber is the class of “crystalline elastomers” based on syndiotactic polypropylene (sPP) and its copolymers.13,14,16 The unusual elastic behavior of sPP is associated with a reversible crystal− crystal martensitic-like phase transition between the metastable form III with chains in the trans-planar conformation, which develops upon stretching, and the more stable form II, with chains in helical conformation, which develops upon releasing the tension.6−9,13−15 The study of these materials has introduced new concepts in thermoplastic elastomers and a definition of unconventional elastomer with high crystallinity and large modulus,8,13,14 with strength, modulus, and elasticity that can be tailored through a balance of enthalpic and entropic contributions to the elastic recovery by modification of the chemical structure and degree of crystallinity.13,14,16−19 1,4-Poly(1,3-diene)s are typical commercial elastomers (polybutadiene and butadiene copolymers, polyisoprene, polychloroprene, etc.) having glass-transition temperatures well below room temperature and containing double bonds Received: August 2, 2017 Revised: December 13, 2017

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DOI: 10.1021/acs.macromol.7b01661 Macromolecules XXXX, XXX, XXX−XXX

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SAXS patterns and equatorial sections covering an angular sector of ±20° around the horizontal (equator) for the WAXS patterns. The WAXS profiles integrated over the equator were reported as a function of the in-plane component qxy of the scattering vector q, whereas the SAXS intensity integrated over the equator or meridian was reported as a function of the in-plane component qxy or of the vertical component qz of the scattering vector q for the equatorial or meridional profiles, respectively. Mechanical tests were performed at 25 °C on amorphous compression-molded films of iP3MPD aged at 20 °C for different aging time as described in the Supporting Information. The elastomeric properties were evaluated measuring the residual deformation (tension set tb or ts(ε)) and elastic recovery (rb or r(ε)) after break or after a given deformation ε = [(L − L0)/L0] × 100 with L and L0 the final and initial lengths of the specimen, respectively, as tb or ts(ε) = [(Lr − L0)/L0] × 100 and rb or r(ε) = [(L − Lr)/Lr] × 100, respectively, with Lr the final length of the relaxed specimens, whereas the percentage of deformation which is recovered after the break was evaluated as Rb = [(L − Lr)/[(L − L0)] × 100 = 100 × (εb − tb)/εb.

along the main chain that afford light cross-linking to create the elastomeric network. 1,2-Poly(1,3-diene)s contain double bonds in the pendant groups and, generally, present glasstransition temperatures higher than those of the 1,4-polymers and, correspondingly, are not in the class of synthetic rubbers. Isotactic 1,2-poly((E)-3-methyl-1,3-pentadiene) (iP3MPD) is a new crystalline polymer that could not be synthesized with the classic stereospecific Ziegler−Natta catalysts but can be synthesized with catalysts based on complexes of cobalt with various phosphines in combination with methylaluminoxane (MAO).20−23 In our previous paper we have reported the crystalline structure of iP3MPD and an exploratory analysis of the mechanical properties.24 Here we analyze the mechanical behavior of the amorphous iP3MPD. Our results have indicated that iP3MPD, in spite of the (relatively) high glass-transition temperature (≈35 °C), shows rubbery elastomeric behavior at 25 °C correlated with the formation (crystallization) during tensile deformation of a mesophase. The elasticity of iP3MPD is an unusual case of a polymer that shows rubbery elastomeric properties below the glass-transition temperature. However, this new polymer could not be used as a synthetic rubber because of the high glass transition temperature, but as other 1,2-poly(1,3-diene)s, it could find application as a thermoplastic material with improved toughness.

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RESULTS AND DISCUSSION In our previous paper we showed that iP3MPD is crystalline, but it is not able to crystallize from the melt.24 The amorphous slowly crystallizes keeping the sample at 25 °C for a long time.24 This behavior is summarized in Figure 1. The large

EXPERIMENTAL SECTION

A highly isotactic sample of iP3MPD with Mw = 81 000 g/mol, polydispersity of 1.2, and isotacticity [mm] > 90% has been synthesized as reported in ref 20. Wide- and small-angle diffraction experiments with synchrotron radiation were conducted at the Dutch-Belgian beamline BM26Dubble of ESRF (Grenoble, France). The wavelength of incident Xrays was λ = 0.10402 nm, and the sample-to-detector distance was 0.27 m for wide-angle (WAXS) and 4.58 m for small-angle X-ray scattering (SAXS) measurements. The covered range for the scattering vector q was 6−30 nm−1 for WAXS and 0.6−1.8 nm−1 for SAXS data, where q = 4π sin θ/λ, and 2θ is the scattering angle. WAXS and SAXS patterns were recorded during tensile deformation of the sample. Strips of initial gauge length 15 mm and width of 5 mm have been cut from a compression-molded film having thickness of 0.41 mm and mounted in a Linkam tensile stage TST350. Sample jaws move in opposite directions, and scattering data have been collected in transmission, using a highly collimated incident beam hitting the sample at the center of the gauge length with a size of 200 μm × 200 μm. Scattering data have been collected at a rate of 1 frame/2 s, while deforming the sample at a rate of 10 mm/min. Two consecutive deformation step cycles have been performed, consisting in stretching the sample up to ≈150% strain (path 1) and then releasing the tension up to measure force zero, in the first step, and then stretching again the sample up to ≈210% strain (path 2) and then releasing the tension up to measure force zero, in the second step. The SAXS and WAXS bidimensional patterns display a cylindrically symmetric distribution of intensity as a function of two photographic coordinates z and r. The z-axis is vertical (meridian) and is parallel to the stretching direction, whereas r is horizontal (equator) and is perpendicular to z. The z coordinate is related to qz, the z-component of the scattering vector q in the reciprocal space, whereas the r coordinate is related to qxy, the projection of the scattering vector q in the plane normal to qz and placed at the origin of reciprocal space. Because of the cylindrical distribution of intensity around qz, qxy = (qx2 + qy2)1/2 with qx and qy the x- and y-components of the scattering vector q. The scattered intensities I(q) were then subtracted for the empty Linkam cell and processed with the programs FIT2D25,26 and Bubble27 to extract meridional and equatorial profiles, in angular sectors covering a range of ±30° around the horizontal (equator) and vertical (meridian) axes for the

Figure 1. X-ray diffraction curves of a compression-molded sample of iP3MPD (a) and after aging at 25 °C for 4 weeks (b).

peaks 2θ ≈ 10° and 19° in the diffraction pattern of the amorphous specimens obtained by compression-molding and cooling the melt to 25 °C (Figure 1a) transform into narrow reflections at 2θ ≈ 10° and 16.2° by aging at 25 °C for several weeks (Figure 1b), indicating crystallization during aging. Both the DSC cooling and successive heating curves of iP3MPD of Figure 2A, recorded at scanning rate of 10 °C/min, present only a glass transition at nearly 35 °C. An endothermic peak near the glass transition is apparent in the heating scan (curve b of Figure 2A). It is well-known that such hysteresis behavior of glasses is correlated with the formation of the glass and of the equilibrium liquid at different cooling and heating rates and/or when internal stresses or strain of any kind are frozen into the glass which are released in the vicinity of the glass transition.28−30 In the former case the endothermic effect in the heating scan increases with decreasing cooling rate,29 whereas in the latter case annealing at temperatures below the glass-transition changes the glass to a more stable state otherwise produced on slow cooling without stress and increase of the endothermic hysteresis with increase of the annealing time has been observed.28,30 Similar behavior was found for iP3MPD in the DSC heating curves of Figure 2B recorded at the same heating rate of 10 °C/min of B


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the nonaged sample is amorphous and the molecular mass is not very high, it shows good ductility (εb > 500%) and strain hardening (σb ≈ 10 MPa). A yielding point at strain εy = 10− 20% and stress σy = 5 MPa is clearly observed even in the nonaged compression-molded initially amorphous sample (curve a of Figure 3). We have checked that under the adopted conditions all samples deform uniformly at room temperature, and any indication of shear banding has not been observed, neither near the yield nor near the breaking point. The yielding phenomenon followed by softening is more evident for the aged samples, and the yield stress increases with increasing the aging time up to 10 and 15 MPa (curves b and c in Figure 3). This phenomenon occurs even for short aging time (the highest aging time is lower than 5 days), not enough to allow crystallization, which occurs for aging time longer than 1 month, as shown in Figure 1. Therefore, the increase of yield stress followed by softening upon physical aging is not correlated with the crystallization. Since cooling from the equilibrium liquid at low cooling rates produces aging at temperature below glass transition, the data of Figures 2B and 3 indicate that upon physical aging the endothermic peak near the glass transition in the DSC heating scans develops (Figure 2B), and simultaneous increase of yield stress is observed (Figure 3). This is in agreement with what observed in the literature that upon aging of amorphous polymers simultaneous increases of DSC enthalpy overshoot close the glass transition and of yield stress are observed.31−40 Both effects have been rationalized in terms of an evolution of the potential energy upon aging,35−39 resulting in local attraction of individual atoms to their neighbors, that is, a local densification or formation of local order upon physical aging.39,40 The local order is destroyed or relaxed by heating the aged glass to induce molecular mobility at the glass transition. This results in the growing of the endothermic peak in the DSC signal close to the glass transition.39 Similar to temperature, an applied stress leads to an increase in segmental mobility41,42 with a resulting decrease in viscosity leading to yielding, the onset of flow, induced by stress (stress-induced glass transition). Therefore, if segmental mobility is increased not by temperature, but by stress, a growing peak in the stress−strain curves is observed and interpreted as an increase in yield stress followed by softening.39 The stress−strain curves of Figure 3 of nonaged and aged samples present a second maximum at strain of nearly 200% after the yield point observed at low deformation of 10−20%. To explain this unusual second maximum, the structural evolution of iP3MPD during deformation has been followed recording wide- and small-angle X-ray diffraction patterns during stretching. The bidimensional wide-angle X-ray diffraction patterns, and the corresponding profiles along the equator, of the compression-molded amorphous nonaged film of iP3MPD, before stretching (sample of Figures 1a and 3a), and deformed at 300% strain, are shown in Figure 4A,B. As in Figure 1a, the diffraction image of the nonstretched sample of Figure 4A presents two broad and diffuse haloes at 2θ ≈ 10° and 19°, the halo at 2θ ≈ 10° being of higher intensity. In the diffraction image of the specimen stretched at 300% strain of Figure 4B the strong reflection at 2θ ≈ 10° becomes less broad and is more concentrated on the equator, whereas the weak reflection at 2θ ≈ 19° is slightly concentrated on a layer off the equator. The increase of the intensity of the reflection at 2θ ≈ 10° and the decrease of that of the reflection at 2θ ≈ 19° on the profiles read on the equator of Figure 4B′ are a clear

Figure 2. (A) DSC cooling curve down to low temperature at 10 °C/ min of iP3MPD (a) and successive heating curve at 10 °C/min (b). (B) Heating curves recorded always at 10 °C/min of samples cooled at the indicated different cooling rates (c.r.). The endothermic enthalpy overshoot near to the glass transition in the second heating scan is evidenced. The scale bar of the heat flow in W/g is indicated.

samples cooled at different cooling rates. It is apparent that the endothermic peak is almost absent in the sample cooled at the highest cooling rate (by quenching at low temperature, curve a of Figure 2B) and increases with decreasing cooling rate (curves b−f of Figure 2B). The stress−strain curve of the amorphous iP3MPD prepared by compression-molding (Figure 1a) is shown in Figure 3. The stress−strain curves of the same amorphous sample aged at 20 °C for different aging time are also shown in Figure 3. The mechanical parameters determined from the stress−strain curve of the nonaged sample are reported in Table 1. Although

Figure 3. Stress−strain curves of compression-molded samples of iP3MPD cooled from the melt to 25 °C (a) and aged at 20 °C for different aging times of 3 h (b) and 3 days (c). The curves are reported as a function of strain ε = 100[(L − L0)/L0] and of logarithm of the draw ratio ln(L/L0). C


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Table 1. Elastic Modulus (E), Stress (σb) and Strain (εb) at Break, Stress (σy) and Strain (εy) at the Yield Point, Tension Set (tb) and Elastic Recovery (rb) at Break, Percentage of Deformation Which Is Recovered after the Break (Rb), Determined from the Stress−Strain Curve of Unoriented and Initially Amorphous Compression-Molded Nonaged Films of iP3MPD of Figure 3a and of the Stress-Relaxed Specimens of Figure 6A. The X-ray crystallinity (xc) is also reported. sample

xc (%)

E (MPa)

σy (MPa)

εy (%)

σb (MPa)

εb (%)

tb (%)

rb (%)

Rb (%)

amorphous compression-molded stress-relaxed specimens

0 0

70 ± 10 85 ± 11

5±1 11 ± 2

20 ± 3 5±1

13 ± 1 31 ± 8

540 ± 30 83 ± 10

104 ± 7 0

205 ± 7 83 ± 10

80 100

indication of the polarization of the first reflection at 2θ ≈ 10° on the equator and of the second reflection at 2θ ≈ 19° on a layer off the equator. Even though a slight decrease of broadening of the reflection is observed, the still broad diffraction peaks suggest that a mesophase of iP3PMD, characterized by structural disorder and, probably, slight order in the chain conformation, crystallizes by deforming the amorphous sample (Figure 4B).24 This could explain the peak at 200−300% deformation in the stress−strain curves of Figure 3, which could be interpreted as a second yielding. Therefore, this noncommon mechanical behavior with a double yielding may be interpreted as a result of events typical of yielding at low deformation εy ≈ 10% that involve alignment of amorphous chains and softening, followed by stress-induced crystallization of the mesophase at higher deformation (200−300%) and successive strain hardening. Finally, the elastic recovery rb of 205%, observed after breaking from 500% strain in the stress−strain curves of Figure 3 (Table 1), indicates that the amorphous or aged samples of iP3MPD show elastic properties. In fact, from a draw ratio at breaking of Lf = 6L0, the final length after breaking is Lr = 2L0 (tension set ≈100%), and the percentage of deformation which is recovered after the break is Rb = 80% (Table 1).

The elastic behavior has been further studied by performing successive cycles of tensile deformation and relaxation of previously stretched and relaxed specimens, prepared as in Figure 5A by stretching the compression-molded and aged films up to 300% deformation (continuous line of Figure 5A) (as in Figure 4A,B) and then releasing the tension (dashed line in Figure 5A). In this procedure, the elastic recovery from Lf = 4L0 (ε = 300%) observed 10 min after releasing the tension was ≈70%, and the residual deformation (tension set) was ≈140% (length of the relaxed specimens Lr = 2.4L0). The diffraction image of Figure 4C of the specimens stretched up to 300% deformation and relaxed by releasing the tension shows that the mesophase formed upon deformation (Figure 4B) melts upon releasing the tension, and the stress-relaxed specimens are amorphous. The stress−strain hysteresis cycles of tensile deformation and relaxation on these stress-relaxed specimens are reported in Figure 5B. In each cycle, the specimens are stretched up to εmax ≈ 300%, that is, a deformation of 60−70% with respect to the relaxed dimension Lr = 2.4L0 of the specimens, corresponding to the elastic recovery observed during the preparation of the stress-relaxed specimens (dashed curve of Figure 5A). It is apparent that a perfect elastic recovery is observed after each cycle with the tension set after the end of each cycle that decreases in

Figure 4. Wide-angle X-ray diffraction images (A−C) and equatorial profiles (A′−C′) read along the dotted horizontal line shown in (A) of the amorphous compression-molded nonaged sample of iP3MPD before stretching (A, A′) (sample of Figure 1a) and of the sample deformed at 300% strain (B, B′) and after release of the tension (C, C′). The patterns B and C are the same as those of stress-relaxed specimens recorded during consecutive steps of deformation and relaxation of Figure 5B. Stress-relaxed specimens, prepared by stretching the compression-molded films up to 300% deformation and then releasing the tension as in B and C, are stretched again up to the original maximum deformation εmax = 300% (corresponding to nearly 70% of deformation with respect to the length Lr of the stress-relaxed specimens achieved in the cycles of Figure 5B), and the diffraction pattern B, B′ is recorded keeping the specimen under tension. Then, the tension is removed, and the diffraction image C, C′ is recorded. The directions of stretching and of the equator and meridian in the WAXS patterns are indicated. D


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step. This is also confirmed by the clear yielding peak even in the stress−strain curve of Figure 6A recorded up to the break

Figure 5. Stress−strain curve of the compression-molded and aged samples of iP3MPD stretched up to 300% deformation (A, continuous line) and successive relaxation after releasing the tension (A, dashed line) for the preparation of stress-relaxed specimens and stress−strain hysteresis cycles recorded on the stress-relaxed specimens prepared in A (B). Three successive cycles of stretching up to εmax ≈ 300%, that is, a deformation of 60−70% with respect to the relaxed dimension Lr = 2.4L0 of the specimens, and relaxation at controlled rates are shown (B). The curves are reported as a function of strain ε = 100[(L − L0)/L0] and of logarithm of the draw ratio ln(L/L0).

Figure 6. Stress−strain curve (A) and DSC heating curve at 10 °C/ min (B) of the stress-relaxed specimens of iP3MPD after the hysteresis cycles of Figure 5B. The stress−strain curve is reported as a function of strain ε = 100[(L − L0)/L0] and of logarithm of the draw ratio ln(L/L0).

on the specimens after being experienced the mechanical hysteresis cycles of Figure 5B (Table 1). The obtainment of an amorphous phase with local order is also indicated by the endothermic enthalpy overshoot near to the glass transition in the DSC heating curve of stress-relaxed specimens of Figure 6B, which is similar to the overshoot observed during the heating scans in the DSC curves of Figure 2B. It is also worth remarking that since the glass-transition temperature of iP3MPD is about 35 °C, slightly above room temperature, the elasticity of iP3MPD is an unusual case of a polymer that shows rubbery elastomeric properties below the glass-transition temperature. The stretching behavior of aged samples of iP3MPD and the yielding during successive stretching and relaxation cycles have been also studied by collecting X-ray scattering data at wide (WAXS) and small (SAXS) angle in real time during two consecutive cycles of deformation at rate of 10 mm/min and recording simultaneously the corresponding stress−strain curves. WAXS data have been recorded only along the equator. The bidimensional SAXS and WAXS images and the stress−strain curves recorded in situ during tensile deformation of a sample of iP3MPD compression-molded and kept at 25 °C for 3 days are shown in Figure 7. The sample is first stretched up to ≈150% deformation (patterns A−F, A′−F′ of path 1 in Figure 7). The corresponding stress−strain curve (path 1 of Figure 7K) shows a well-pronounced yielding because of the aging. The tension is then released up to reach a null force. The sample recovers only minimally the initial dimensions, achieving a deformation of 123% (patterns G, G′ of Figure 7). Successively, the sample is stretched immediately

successive cycles and is almost zero after the third cycle (Figure 5B). However, an unusual yielding point is observed in each cycle, not common when oriented elastic specimens that have already experienced the irreversible plastic deformation during the first stretching of the unoriented compressionmolded sample (as in Figures 3 and 5A) are deformed again. This peaked yielding phenomenon could be explained by the diffraction images recorded during successive stretching and relaxation by releasing the tension (Figure 4B,C) that indicate, as discussed above, formation of the mesophase upon deformation at 300% strain of amorphous samples (Figure 4B) and its melting upon releasing the tension and elastic recovery (Figure 4C). Similar reversible phase transition is observed during the cycles of tensile deformation and relaxation of Figure 5B. The elastic return is associated with melting of the mesophase and loss of chain orientation with obtainment of unoriented amorphous sample. Crystallization of the mesophase upon stretching (Figure 4B) and its melting after relaxation (Figure 4C) are reversible transformations that occur in the successive cycles of Figure 5B. The yielding observed during stretching in each cycle of Figure 5B is probably correlated to the crystallization of the mesophase occurring in every mechanical cycle upon stretching and to a sort of recovery of the state of the amorphous that occurs in every mechanical cycle by relaxing the stress and that is somehow reminiscent, at the local scale, of the partial order of the mesophase formed in the stretching E


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Figure 7. Bidimensional SAXS (A−J) and WAXS (A′−J′) patterns, collected during two consecutive step cycles of deformation (deformation rate of 10 mm/min), and corresponding stress−strain curves (K) recorded in situ during deformation of a compression-molded sample of iP3MPD aged at 25 °C for 3 days. The sample is first stretched up to 151% deformation (A−F, A′−F′, path 1), then the tension is released up to measure a null force (G, G′), it is stretched again from 123% deformation up to 211% deformation (G−I, G′−I′, path 2), and, finally, the tension is removed (J, J′). In both SAXS and WAXS patterns the equator is horizontal and the stretching direction is vertical.

again from 123% strain up to 211% deformation (patterns G− I, G′−I′ of path 2 in Figure 7, and path 2 of the stress−strain curve of Figure 7K), and then the tension is released (patterns J, J′ of Figure 7) up to measure a null force. Therefore, at variance with the stress−strain hysteresis cycles of Figure 5B, the sample is stretched and relaxed continuously during the two cycles (Figure 7K), without introducing any delay time after each cycle. As a consequence, the nonaged sample shows no yielding during the stretching in the path 2 of Figure 7K. The equatorial WAXS images of Figure 7A′−J′ reveal more clearly and in more detail the structural transformations which occur during the stretching steps, already evidenced by the diffraction images of Figure 4. The monodimensional equatorial profiles of the images of Figure 7A′−J′ are shown in Figure 8. The two halos on the equator, centered at q ≈ 7 and 13 nm−1 (corresponding to the haloes at 2θ ≈ 10° and 19° in the images of Figure 4), are clearly visible in the image of Figure 7A′ of the undeformed sample (curve a of Figure 8). The intensity of the first halo at q ≈ 7 nm−1 increases and becomes concentrated on the equator with increasing deformation, whereas the intensity of the second halo at q ≈ 13 nm−1 decreases (Figure 7B′−F′ and profiles b−f of Figure 8) and disappears almost completely at the maximum deformation achieved in both paths 1 and 2 (Figures 7F′ and 7I′ and profiles f and i of Figure 8). As observed in Figure 4, the increase of the intensity of the reflection at q ≈ 7 nm−1 and the disappearance of the reflection at q ≈ 13 nm−1 on the equator of the WAXS images of Figure 7 are a clear indication of the polarization of the first reflection at q ≈ 7 nm−1 on the equator and of the second reflection at q ≈ 13 nm−1 on a layer off the equator. This, in turn, suggests that the mesophase crystallizing by stretching is characterized probably by slight order in the chain conformation. Upon releasing the tension, the intensity of the halo at q ≈ 7 nm−1 becomes more isotropic, whereas the intensity of the halo at q ≈ 13 nm−1 increases (Figure 7G′, J′ and profiles g, j of Figure 8). This demonstrates that the mesophase crystal-

Figure 8. WAXS profiles integrated over the equator extracted from the WAXS images of Figure 7A′−J′ reported as a function of the inplane component qxy of the scattering vector q.

lized by stretching melts after releasing the tension, and therefore the crystallization of mesomorphic domains induced by stretching is reversible in consecutive step cycles of deformation and release of the tension. Reversible changes also occur at SAXS length scale. In fact, in the SAXS images of Figure 7A−J the SAXS intensity tends to polarize on the equator with increasing deformation (Figures 7A−F and 7H,I) and becomes more isotropic upon releasing the tension (Figure 7G). The polarization of the SAXS intensity on the equator observed at the maximum deformation of 151% achieved in the path 1 (Figure 7F), and 211% achieved in the path 2 (Figure 7I), may be also evidenced by comparing the SAXS profiles on the equator and the meridian at 0, 151%, and 211% deformation in Figure 9. It F


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Figure 9. SAXS intensity integrated over the equator and meridian extracted from the SAXS images of Figure 7A−J reported as a function of the in-plane component qxy of the scattering vector q for the equatorial profiles and of the vertical component qz of the scattering vector q for the meridional profiles.

is apparent that while the equatorial and meridional profiles are coincident for the undeformed sample, the intensity on the equator becomes almost 1 order of magnitude higher than the intensity on the meridian at 151% and 211% deformations (Figure 9). This is the hallmark that deformation produces formation of elongated domains with electron density higher than the surrounding matrix oriented with the long axis in the stretching direction. As sketched in Figure 10, these domains are arranged with no lateral and no longitudinal order and have average dimensions higher than 2π/qmin ≈ 100 nm, as evidenced by the remarkable increase of intensity on the equator for the stretched samples, due to some kind of “structuring” at length scale larger than 100 nm. As shown by the WAXS data of Figures 4 and 7, these domains are characterized by crystals of the mesophase with chains in an ordered helical conformation.

Figure 10. Model of elongated domains of the mesophase of iP3MPD formed by stretching. For simplicity, the domains have been drawn of elliptical shape, and it is assumed that the chains are in an ordered helical conformation and the domains are arranged with no lateral and longitudinal order.

glass due to a local densification or formation of local order in the amorphous glass upon physical aging at temperatures below the glass transition. The local order is relaxed upon heating of the glass, and a yielding peak results in the stress− strain curve by application of tensile stress. The second stress peak in the stress−strain curve at higher strain of 200−300% is instead correlated with the crystallization of a disordered mesomorphic form during stretching. Therefore, this uncommon behavior upon stretching with a double yielding may be correlated with events typical of yielding at low strain εy ≈ 10−20% with alignment of amorphous chains and softening, followed by stress induced crystallization of the mesophase at higher deformation (200− 300%) and successive strain hardening. Amorphous or aged samples of iP3MPD show elastic behavior. During successive cycles of tensile deformation and relaxation a yielding peak is still observed during every stretching step, and is correlated with the crystallization of the mesophase upon stretching that melts during elastic recovery. SAXS and WAXS patterns recorded simultaneously during deformation and relaxation have indicated that stretching produces formation of elongated domains of the mesophase, of few hundred nanometers size, oriented with the long axis in the elongation direction and arranged with no lateral and no longitudinal order. The elasticity of iP3MPD is an unusual case of a polymer that shows rubbery elastomeric properties below the glasstransition temperature.

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CONCLUSIONS The mechanical behavior of iP3MPD and the structural and morphological transformations that occur upon stretching and relaxation have been studied. iP3MPD is a crystalline polymer with glass-transition temperature of ≈35 °C. It crystallizes from the amorphous by aging at 25 °C. An endothermic peak near to the glass transition, typical of hysteresis behavior of glasses, is observed in the DSC heating curves of amorphous samples. This endothermic effect increases with decreasing the cooling rate at which the glass is achieved from the equilibrium liquid, which corresponds to an aging of the glass at temperatures below the glass transition. The stress−strain curves of amorphous iP3MPD have shown absence of viscous flow but high ductility, strain hardening, and elastic behavior. An unusual double yield point during deformation has been observed. The yielding phenomenon at low deformation followed by softening is more evident for the samples aged at temperatures below glass transition and the yield stress increases with increasing the aging time. The simultaneous increase of yielding stress and of the endothermic hysteresis effect near to the glass transition upon aging can be explained with the creation of internal stresses frozen into the G


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(12) De Rosa, C.; Ruiz de Ballesteros, O.; Auriemma, F. Mechanical properties of helical and mesomorphic forms of syndiotactic polypropylene at different temperatures. Macromolecules 2004, 37, 7724. (13) De Rosa, C.; Auriemma, F. Structure and properties of syndiotactic polypropylene: a highly crystalline thermoplastic elastomer. Prog. Polym. Sci. 2006, 31, 145. (14) De Rosa, C.; Auriemma, F.; Ruiz de Ballesteros, O. The Role of Crystals in the Elasticity of Semicrystalline Thermoplastic Elastomers. Chem. Mater. 2006, 18, 3523. (15) Auriemma, F.; De Rosa, C.; Esposito, S.; Mitchell, G. R. Polymorphic “Super-Elasticity” in Semicrystalline Polymers. Angew. Chem., Int. Ed. 2007, 46, 4325. (16) De Rosa, C.; Auriemma, F. Single-Site Metallorganic Polymerization Catalysis as a Method to Probe the Properties of Polyolefins. Polym. Chem. 2011, 2, 2155. (17) De Rosa, C.; Auriemma, F. From entropic to enthalpic elasticity: novel thermoplastic elastomers from syndiotactic propylene-ethylene copolymers. Adv. Mater. 2005, 17, 1503. (18) De Rosa, C.; Auriemma, F. Mechanical Properties of Syndiotactic Propylene-Ethylene Copolymers. Macromolecules 2006, 39, 249. (19) De Rosa, C.; Auriemma, F.; Corradi, M.; Caliano, L.; Ruiz de Ballesteros, O.; Di Girolamo, R. Mechanical properties and elastic behavior of syndiotactic propene-butene copolymers. Macromolecules 2009, 42, 4728. (20) Ricci, G.; Leone, G.; Boglia, A.; Bertini, F.; Boccia, A. C.; Zetta, L. Synthesis and Characterization of Isotactic 1,2-Poly(E-3-methyl1,3-pentadiene). Some Remarks about the Influence of Monomer Structure on Polymerization Stereoselectivity. Macromolecules 2009, 42, 3048. (21) Ricci, G.; Forni, A.; Boglia, A.; Motta, T.; Zannoni, G.; Canetti, M.; Bertini, F. Synthesis and X-ray Structure of CoCl2(PiPrPh2)2. A New Highly Active and Stereospecific Catalyst for 1,2 Polymerization of Conjugated Dienes When Used in Association with MAO. Macromolecules 2005, 38, 1064. (22) Ricci, G.; Forni, A.; Boglia, A.; Sommazzi, A.; Masi, F. Synthesis, structure and butadiene polymerization behavior of CoCl2(PRxPh3‑x)2 (R = methyl, ethyl, propyl, allyl, isopropyl, cyclohexylx = 1, 2). Influence of the phosphorous ligand on polymerization stereoselectivity. J. Organomet. Chem. 2005, 690, 1845. (23) Ricci, G.; Leone, G.; Boglia, A.; Boccia, A. C.; Zetta, L. cis-1,4alt-3,4 Polyisoprene: Synthesis and Characterization. Macromolecules 2009, 42, 9263. (24) De Rosa, C.; Auriemma, F.; Santillo, C.; Scoti, M.; Malafronte, A.; Zanchin, G.; Pierro, I.; Leone, G.; Ricci, G. Crystal Structure and Properties of Isotactic 1,2-Poly(E-3-Methyl-1,3-Pentadiene). Macromolecules 2017, 50, 5412. (25) Hammersley, A. P.; Svensson, S. O.; Hanfland, M.; Fitch, A. N.; Häusermann, D. Two-Dimensional Detector Software: From Real Detector to Idealised Image or Two-Theta Scan. High Pressure Res. 1996, 14, 235. (26) Hammersley, A. P.; Svensson, S. O.; Thompson, A. Calibration and correction of spatial distortions in 2D detector systems. Nucl. Instrum. Methods Phys. Res., Sect. A 1994, 346, 312−321. (27) Dyadkin, V.; Pattison, P.; Dmitriev, V.; Chernyshov, D. A new multipurpose diffractometer PILATUS@SNBL. J. Synchrotron Radiat. 2016, 23, 825. (28) Wunderlich, B. Macromolecular Physics; Academic Press: 1976; Vol. 2, p 362. (29) Wunderlich, B. A thermodynamic description of the defect solid state of linear high polymer. Polymer 1964, 5, 125. (30) Weitz, A.; Wunderlich, B. Thermal analysis and dilatometry of glasses formed under elevated pressure. J. Polym. Sci., Polym. Phys. Ed. 1974, 12, 2473. (31) Hodge, I. M. Enthalpy relaxation and recovery in amorphous materials. J. Non-Cryst. Solids 1994, 169, 211. (32) Hutchinson, J. M. Physical aging of polymers. Prog. Polym. Sci. 1995, 20, 703.

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.7b01661.

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Experimental details of preparation of stretched oriented specimens, X-ray diffraction measurements, and analysis of mechanical properties (PDF)

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (C.D.R.). ORCID

Claudio De Rosa: 0000-0002-5375-7475 Finizia Auriemma: 0000-0003-4604-2057 Anna Malafronte: 0000-0002-7854-5823 Giuseppe Leone: 0000-0001-6977-2920 Giovanni Ricci: 0000-0001-8586-9829 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS Financial support from “Ministero dell’Istruzione, dell’Università e della Ricerca” (project PON-DIATEME 2007-2013) and from CARIPLO Foundation (Crystalline Elastomers Project) is gratefully acknowledged. We thank Daniel Hermida-Merino for his experimental assistance at the Dutch−Belgian Beamline (DUBBLE) of ESRF (Grenoble, France).

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REFERENCES

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Mechanical Properties of Isotactic 1, 2-Poly (E-3-methyl-1, 3

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