Article Cite This: Macromolecules XXXX, XXX, XXX−XXX
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Interplay between Crystallization and Entanglements in the Amorphous Phase of the Crystal-Fixed Polymer Poly(ϵ-caprolactone) Ricardo Kurz,† Martha Schulz,† Felix Scheliga,‡ Yongfeng Men,§ Anne Seidlitz,† Thomas Thurn-Albrecht,† and Kay Saalwächter*,† †
Institut für Physik, Martin-Luther-Universität Halle-Wittenberg, 06099 Halle (Saale), Germany Institut für Technische und Makromolekulare Chemie, Universität Hamburg, Bundesstr. 45, 20146 Hamburg, Germany § State Key Laboratory of Polymer Physics and Chemistry, Changchun Institute of Applied Chemistry, Chinese Academy of Sciences, Renmin Street 5625, 130022 Changchun, P. R. China Downloaded via KAOHSIUNG MEDICAL UNIV on July 26, 2018 at 08:40:20 (UTC). See https://pubs.acs.org/sharingguidelines for options on how to legitimately share published articles.
‡
ABSTRACT: This work focuses on the influence of amorphous-phase entanglements on the semicrystalline morphology of poly(ϵ-caprolactone). This polyester is classified as crystal-fixed; i.e., it displays no translational chain dynamics in the crystals. We study a wide range of well-entangled samples with molecular weights up to several million using polarization microscopy to assess the lamellar growth rate, small-angle X-ray scattering and proton time-domain NMR to characterize the morphology, and proton multiple-quantum NMR, auxiliary carbon-13 NMR, rheology, and tensile deformation to assess the entangled chain dynamics in the melt and semicrystalline states. We demonstrate a significant increase in the density of entanglements in the amorphous phase relative to the melt. The dependencies of our observables on crystallization temperature and molecular weight suggest that entanglements control the thickness of the amorphous layers. This is rationalized by an only slowly relaxing exclusion zone with enhanced entanglement density acting as an entropically repulsive layer between adjacent lamellae.
I. INTRODUCTION When cooled below a specific temperature, stereoregular polymers will undergo a phase transformation from their random-coil structure in the melt to a semicrystalline (SC) state. During this process, well-ordered lamellae of a certain thickness are formed, separated by regions that remain amorphous and more mobile and partly consist of chains connecting both phases.1 Depending on their nanostructure, SC polymers show a wide range of mechanical properties, so understanding the determining factors of the final structure is an important but as yet unsolved problem in polymer science. Previous work has in fact mostly focused on the factors governing the chain-folding structure and the actual thickness of the crystalline lamellae.2 Naturally one can expect an interplay between crystalline structure formation and the dynamics of the amorphous fraction during crystallization, the latter possibly constraining crystal growth. In one of the few theoretical accounts, Iwata3 explained a decrease of crystallinity with increasing molecular weight (MW) by a successively increasing the amount of topological constraints hindering the formation of lamellae. A theory employing a similar (entropic) argument has very recently been used to explain the finite size of crystallites appearing in straininduced crystallization of natural rubber and could be convincingly confirmed by comparison with experimental data.4 Note that recently Schmidt-Rohr and co-workers5 have also stressed an important role of MW in a sense that chain ends are relevant in resolving potential density anomalies in the crystal−amorphous interphase. Such an MW effect is, however, © XXXX American Chemical Society
readily offset by a chain tilt in the crystal. Generally, it seems amorphous regions have rarely been addressed in the data analysis and are for the most part not explicitly considered in crystallization theories. On the other hand, it is by now well-known that the mechanical properties of SC polymers depend significantly on variations within the amorphous phase. Mandelkern and coworkers highlighted the impact of the interlamellar spacing,6 and stress−strain curves7,8 were shown to reflect a potential increase of entanglement (or, more generally, constraint) density in the amorphous regions as compared to the bulk melt.9,10 Extending such work and using nuclear magnetic resonance spectroscopy (NMR) as a molecular-level probe, we here focus on the role of the amorphous phase in structure formation during crystallization. We shall demonstrate experimentally that the final topological state in the amorphous layers and the overall morphology as well as the macroscopic mechanical properties are critically related. Already early on, the question whether reorganization of an entangled melt is possible during or after crystallization was critically discussed. Initially, Flory and Yoon argued that the highly entangled state must be largely conserved in the semicrystalline state,11 since disentanglement should be prevented by fixation of the chains in the crystalline lamellae. This was contradicted by DiMarzio et al., who compared two Received: April 16, 2018 Revised: July 12, 2018
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DOI: 10.1021/acs.macromol.8b00809 Macromolecules XXXX, XXX, XXX−XXX
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Macromolecules different models12 and claimed having found fundamental inconsistencies in the approach of Flory and Yoon, concluding that disentanglement times can be up to 5 decades faster than initially approximated. Klein and Ball13 investigated deposition rates of molecular sequences to the growing lamellae and applied the reptation model and inferred that disentanglement during crystallization should indeed be possible. This by now classical debate remains essentially unsettled, in particular whether a fast “reeling-in” process exists or not. Later, Rault and Robelin-Souffaché14 took up these ideas and considered the interplay between lamellar growth kinetics and the time scale of relaxation in the melt. Following this approach, they distinguished between two crystallization regimes, depending on whether chains can (slow regime) or cannot relax during crystallization (fast regime). On this basis, they were the first to propose a direct relation between the entanglement spacing and the semicrystalline morphology. However, only quench-cooled samples were studied (as it is required by their experimental approach), leaving the question open whether such ideas apply in isothermal crystallization. Up to the present day an exact understanding of the behavior of the amorphous phase and therefore its impact on the final structure is still under discussion. In this context, Hu and Schmidt Rohr15 suggested a classification into crystal-mobile and crystal-fixed polymers, i.e., polymers that do and do not exhibit a so-called αc-relaxation process related to chain motion in the crystal, respectively, the latter providing a specific mechanical coupling of the two phases. Indeed, the αc relaxation was found to be related to ultradrawability in the SC state. The αc yes/no classification essentially mirrors Boyd’s much earlier classification of high- and intermediate-/low-crystallinity polymers, respectively.16 We have recently found that the morphologies of crystal-mobile and crystal-fixed polymers differ qualitatively in that the former have a well-defined amorphousdomain thickness, while the latter have a better-defined crystal thickness.17 The crystal-mobile case, prominently represented by poly(ethylene), features high crystallinity and consequently thin and highly constrained amorphous layers that were the subject of earlier NMR studies.18,19 Here, we focus on the crystal-fixed case and explore the amorphous phase in more detail in order to rationalize our previous finding. Generally, the amorphous phase of long-chain crystal-fixed polymers can be viewed as a disordered melt-like state with additional constraints through chain fixation in the crystallites and a possibly increased entanglement density. This suggest an elastomer-like scenario, in contrast to “simple” reptative dynamics as described by the tube model.20,21 In the constrained amorphous phase only much slower arm retraction22,23 rather than reptation20 prevails, and it seems reasonable to apply this consolidated knowledge on chain dynamics over multiple length and time scales. NMR methods are in fact well suited to provide a molecularlevel probe. Specifically, anisotropic NMR interactions such as dipole−dipole couplings (DDC) between abundant protons are the basis of detailed insights into the degree of motional restrictions to fluctuating polymer chains.24,25 DDCs are progressively averaged by fast yet anisotropic segmental motion, and the measurement of residual DDCs in terms of a monomeraveraged residual DDC constant, Dres, opens avenues to probing the length of subchains between cross-links or entanglements,26,27 and also their change upon stretching.28 Simple Hahn echo experiments have been used extensively29,30 to measure Dres but are subject to ambiguities in the data analysis.31
Figure 1. NMR-observable segmental orientation autocorrelation function C(t = τDQ) reflects the chain dynamics across all regimes of the tube model,33,34 i.e., (I) Rouse, (II) contrained Rouse, (III) reptation, and (IV) terminal. Changes arising from constraints to chain motion within the amorphous phase of an SC polymer, where regimes II and III are potentially modified (being replaced by arm retraction of singly end-fixed chains or a fraction of nonrelaxing tie chains), are illustrated and are the subject of this work. Only a small interval can be covered in real time as defined by the duration of the used DQ pulse sequence, τDQ.
In recent years, low-resolution multiple-quantum (MQ), or more specifically double-quantum (DQ), NMR has been established as the method of choice.25,32 MQ NMR provides direct access to the segmental orientation autocorrelation function (OACF) C(t),35 the amplitude of which is proportional to the apparent Dres2, which approaches a low and time-constant value in networks (constant anisotropy level posed by chain-end fixation) or remains time and thus temperature dependent in systems where the chains can relax. This is illustrated in Figure 1. In entangled polymer melts, all regimes of the tube model could be probed by this technique.33,34 The concept of time−temperature superposition (TTS) must however be used, as the experimental time window accessible by the pulse sequence duration τDQ is rather limited, i.e., to an interval of about 0.1−1 ms for the pulse sequence used in this work. Thus, the temperature-dependent apparent Dres2(T) extracted for the beginning of this interval can be used to map out C(t), as is also apparent from Figure 1. For an estimation of the entanglement density, subject to potential changes in the amorphous phase of SC polymers, we will need to evaluate the amplitude of C(t) at the entanglement time (τe) separating the tube model regimes I and II:33 ij 3 yz zz C(τe) ∝ Dres,e 2 ∝ Se 2 = jjj j 5Ne zz (1) k { Se is the segmental order parameter associated with tube/ entanglement constraints, and Ne ∝ Me is the number of segments per entangled strand, which is proportional to the entanglement molecular weight Me.33 Here, we present a comprehensive study of a wide range of samples of poly(ϵ-caprolactone), which is a crystal-fixed polymer.36 We focus on well-entangled samples spanning about 2 decades in MW, including the probably highest MW samples studied so far, benefiting from recent synthetic progress. After isothermal crystallization from the melt, we characterize 2
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DOI: 10.1021/acs.macromol.8b00809 Macromolecules XXXX, XXX, XXX−XXX
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Macromolecules
crystallinity as a function of time. At 50 °C the crystallization of the two highest-MW samples was not completed within reasonable time. Small-angle X-ray scattering on a Kratky compact camera equipped with a focusing X-ray optics and a temperature-controlled sample holder was employed to determine the long period L and the sizes of the crystalline (dc) and amorphous (da) domains. The latter were obtained from a newly developed analysis approach of the interface distribution function (see refs 17 and 44 for details). Polarization microscopy using an Olympus BX51 was used to determine the speed of lamellar growth as approximated by the speed of the spherulitic growth front. Oscillatory shear rheology on an Ares G2 from TA Instruments with a parallel-plate geometry (8 mm diameter) was used to estimate the terminal relaxation times τd of the samples up to PCL370 at a reference temperature Tref = 70 °C. τd = 1/ωx was taken from the crossover frequency ωx of the real and imaginary parts of the complex shear modulus master curves G*(ω) as measured from 0.1 to 200 rad/s (after melting at about 70 °C) in a range between 50 and 90 °C, taking into account sufficiently slow crystallization and the limited thermal stability of the samples, respectively. The shift factor aT was roughly the same for all samples and followed the WLF equation
Table 1. Overview of PCL Samples Studied samplea
PDI
synthesis/sourceb
⟨r2⟩01/2/nm
PCL49 PCL50 PCL66 PCL90 PCL92 PCL138 PCL232 PCL370 PCL580d PCL2000d PCL4500 PCL5700
1.1 1.54 1.51 1.1 1.48 1.42 1.2 1.8 2−3 2−3 2−3 2−3
this work SPP this work this work SPP SPP this work MLUc this work this work this work this work
13.0 13.1 15.1 17.6 17.8 21.8 28.2 35.6
a
The sample name encodes the GPC-based Mw,PS in kg/mol vs PS standard. bSPP: Scientific Polymer Products, Inc., Ontario, NY; MLU: Martin-Luther-University. cIn-house synthesis following similar procedures as in ref 37. dContains ∼10% low-MW fraction.
the morphological parameters, especially the amorphousdomain thickness da, by small-angle X-ray scattering (SAXS) and proton time-domain NMR as a function of crystallization temperature Tc and molecular weight Mw. 1H MQ NMR, aided by additional 13C NMR experiments, was used to assess the entangled chain dynamics in the amorphous phase, complemented by tensile testing hyphenated with SAXS to assess the relation between amorphous-phase entanglements and mechanical properties. We conclude with a tentative explanation of our findings.
log a T = C1
T − Tref T − TV
(2)
where C1 = 2.85 and TV = −110 °C ≈ Tg − 50 °C. Using also data for lower MW (for which the crossover is outside of our frequency range), we could confirm an MW scaling of the zero-shear viscosity η = lim (ωG″) with an exponent of 3.6 (the small deviation from ω→ 0
the expected value of 3.4 is attributed to limited accuracy). Tensile deformation of SC samples was studied at 30 °C with a portable tensile tester (Linkam TST350) hyphenated with online 2D SAXS (see ref 45 for details). True stress−strain curves were obtained for dog-bone specimen (L0 = 10 mm, b0 = 5 mm, d0 = 0.1 mm) molten for 10 min at 90 °C in a heating press, followed by isothermal crystallization in a water bath. Samples were stretched at a constant speed of 20 μm/s, and the engineering stress σ0 = F/A0 = F/(b0d0) was calculated from the measured force and the initial cross-sectional area. To measure the width b of the sample at the center of the forming neck during deformation, pictures were taken, from which the true stress σ = σ0(b0/b)2 and the true strain λ = (b0/b)2 were obtained (assuming transverse isotropy). We omit here details of the 2D SAXS results46 and just summarize that they evidenced a gradual destruction of the initially disordered SC structure and a transformation into an oriented state with similar crystallinity via recrystallization beyond the yield point, as discussed in detail in a previous publication.47 As originally shown by Haward,7,9 the mechanical behavior of semicrystalline polymers in a certain range beyond the yield point can be approximately described by the parallel combination of a neoHookean entropic spring representing a network with modulus G and an Eyring dashpot with a viscosity η describing the plastic deformation:
II. EXPERIMENTAL SECTION Samples. We have studied PCL samples from different sources as collected in Table 1; the MWs are all specified in kg/mol as part of the sample name. Most relevant to the present work are samples synthesized by one of us (F.S.) at the University of Hamburg, which have been prepared via controlled ring-opening polymerization of ϵcaprolactone at room temperature employing different metal alkyl initiators based on Mg, Zn, and Al following established strategies.38,39 Such systems were found to be highly active at moderate temperatures and useful for the synthesis of ultrahigh-MW products (PCL5700 and PCL4500) as well as low- to medium-MW products with narrow polydispersity of 1.1−1.2, such as PCL232, PCL90, and PCL49. Details of the syntheses are deferred to a separate publication. The MWs were determined by GPC/SEC in THF (lower Mw) and chloroform (Mw ≥ 580 kg/mol), the latter relying on an in-house preparation of suitable column materials for large MWs. We have chosen to report and use (with one exception, see further below) Mw,PS vs polystyrene (PS) standard. A conversion factor to the true Mw has been published for PCL in THF to be around 0.56 for Mw above ca. 10 kg/mol.40,41 The factor is likely somewhat larger for highMw PCL in chloroform. Our NMR investigations revealed an ∼10% fraction of rather mobile material remaining in the amorphous phase of the two samples PCL580 and PCL2000. Table 1 also lists the weight-averaged end-to-end distances of the respective chains in the melt as estimated from data on the unperturbed dimensions of PCL,41 ⟨r2⟩0 = Mw,PS × 3.43 nm2 mol/kg (this relation includes the mentioned MW calibration factor of 0.56). As a reference for entangled dynamics in the melt, the plateau modulus G0N of PCL at 140 °C is 1 ± 0.1 MPa,42 corresponding to an entanglement MW 4 Me = 5 ρRT /G N0 43 of 2500 ± 250 g/mol. With this, we can estimate the end-to-end distance of an entangled chain segment, i.e., the tube diameter, to about 3.9 nm. All samples were crystallized isothermally in a range between Tc of 30 and 50 °C (most samples just at these two temperatures), and care was taken to achieve complete crystallization by monitoring the overall
σ = G(λ 2 − 1/λ) + σvisc
(3)
The viscous contribution σvis = ηλ̇ is independent of the absolute value of the strain λ. More extended models for the mechanical behavior of semicrystalline polymers have meanwhile been developed, including especially a more detailed description of yielding.48,49 However, there is wide agreement in the literature that the modulus G determined from the neo-Hookean range immediately beyond yielding, also called strainhardening modulus, reflects the entanglement density in the amorphous regions.50−52 Fits were thus performed between the yield point and λ = 4, i.e., before the onset of further strain hardening. Our result of interest is G representing the elasticity of the entangled amorphous phase. NMR Spectroscopy. 1H time-domain NMR experiments were performed a 200 MHz Bruker Avance III instrument with a dedicated static probehead featuring a short dead time of only 2.5 μs and a 90° pulse length of the same order. The temperature is regulated by controlled heating of a stream of dried air; the accuracy is estimated to C
DOI: 10.1021/acs.macromol.8b00809 Macromolecules XXXX, XXX, XXX−XXX
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Macromolecules
encodes the time scale of chain motions. The intensity buildup in IDQ(τDQ) reflects the magnitude of residual DCCs (RDCCs) on a time scale of τDQmin (0.1 ms) as quantified by the monomer-averaged quantity Dres, which is temperature-dependent. The amplitude of C(τDQmin) is proportional to Dres2. The shape of both functions thus encodes the shape of C(t). More details of the data analysis are given in the Results and Discussion section. Additional static 13C NMR experiments with 1H decoupling were performed on sample PCL138 (Tc = 30 °C) using a Bruker doubleresonance static probe (5 mm coil inner diameter) on a 400 MHz Bruker Avance II instrument. T1 relaxation times for all resonances in the amorphous phase were determined at different temperatures by means of direct excitation using an appropriate recycle delay of around 2 s to focus on the amorphous fraction, implemented as part of a phasecycle controlled 90° pulse pair storing the signal along ±z for a variable relaxation delay, followed by another 90° pulse and spectral detection. The experiment is similar to the cross-polarization (CP)-based experiment of Torchia56 and allows for robust fitting of the decay to zero intensity, using a single exponential. With a simple saturationrecovery experiment we have verified that the samples are characterized by two fractions of lower T1 (
