Bin
Zhang
abc,
Jingbo
Chen
c,
Hui
Zhang
a,
Moritz C.
Baier
d,
Stefan
Mecking
d,
Renate
Reiter
af,
Rolf
Mülhaupt
ef and
Günter
Reiter
*af
aInstitute of Physics, University of Freiburg, 79104 Freiburg, Germany. E-mail: guenter.reiter@physik.uni-freiburg.de
bHermann Staudinger Graduate School, University of Freiburg, 79104 Freiburg, Germany
cSchool of Materials Science & Engineering, Zhengzhou University, Zhengzhou 450002, People's Republic of China
dChair of Chemical Materials Science, Department of Chemistry, University of Konstanz, 78464 Konstanz, Germany
eInstitute of Macromolecular Chemistry, University of Freiburg, 79104 Freiburg, Germany
fFreiburg Materials Research Centre, University of Freiburg, 79104 Freiburg, Germany
First published on 16th January 2015
Applying a slow annealing procedure, we have transformed geometrically simple, faceted polymer single crystals into periodically branched crystals. Interestingly, the period of the branches increased logarithmically with annealing time and depended on crystallization temperature in a similar fashion as the thickness of the lamellar crystal. We tentatively relate the periodic pattern to meta-stable states, differing in crystalline order and thus melting temperature. The degree of meta-stability and its variance depend on lamellar thickness but can also change with the degree of molecular order, causing differences in melting behaviour. Our results propose that periodic variations in thermal stability within a polymer single crystal can be made visible by annealing.
Annealing-induced morphological changes in polymer single crystals at temperatures well below the nominal melting point have been observed previously.6,30–33 Local molecular reorganization processes were identified as the central mechanism which allowed to partially remove chain folds with increasing annealing time or annealing temperature.31,34,35 As a consequence, the mean lamellar thickness increased, leading to enhanced thermal stability.36,37 Pioneering microscopy works38–42 on individual single crystals have demonstrated that, as a mere consequence of mass conservation, such annealing induced local thickening43 is accompanied by the formation of a “Swiss cheese” morphology or “saw-tooth” patterns.44–47 Such spatially varying changes in morphology induced by annealing reflect heterogeneities, i.e., differences in molecular order within these crystals. However, so far annealing did not reveal any evident spatial correlations of these heterogeneities. No relation to the initial growth conditions and to the conditions of annealing could be established.
One reason for not observing correlated patterns induced by annealing may be attributed to the typically small differences in degree of order within a single crystal. Thus, to uncover heterogeneities,48,49 differences in thermal stability have to be amplified first. Here, we have chosen a lengthy series of annealing steps which started at low temperatures and slowly moved to higher temperatures. Annealing has to remove sequentially, and with high accuracy, crystalline parts which have an only slightly lower thermal stability than more perfect ones. Upon proper annealing, less stable regions will melt. Simultaneously, the remaining regions will become even more stable, e.g., thicker. Consequently, the finally surviving regions reflect, to a large extent, the more stable crystalline parts formed already during crystal growth. By applying suitable annealing sequences, we attempted to unveil the less stable parts of a polymer crystal and aim to answer the following question: are heterogeneities within facetted lamellar single crystals distributed homogeneously or are they arranged according to some pattern?
As shown in Table 1, we studied four high molecular weight polyethylene samples, including two homopolymers of protonated polyethylene (PE), a deuterated PE (dPE) and a diblock copolymer of hydrogenated and deuterated polyethylene (hPE-b-dPE) of different molecular weights.52 For all these polymers, single crystals were prepared by a self-seeding method,53 from dilute solutions at appropriate crystallization temperatures (Tc). The polymers were dissolved in p-xylene at rather low concentrations of 0.001 wt% or 0.0001 wt% by heating the solution (ca. 1 ml in a closed glass vial) first to 130 °C for 30 min, i.e., well above the nominal dissolution temperature of ca. 100 °C. This so homogenized polymer solution was then rapidly crystallized and aged for 2 hours at 70 °C in a thermostated bath filled with silicone oil. The solution was subsequently heated to the respective self-seeding temperature Ts, where it was kept for 10 minutes. For growing single crystals, the samples were subsequently quickly transferred to another thermostated bath at a pre-set crystallization temperature Tc. Drops of such solutions containing suspended PE single crystals were then deposited onto a silicon wafer, an often used substrate.47,54–57 After deposition, the solvent was allowed to evaporate completely in vacuum at 25 °C for 12 hours.
Code | Material | M w [g mol−1] | PDI | T s (self-seeding temperature), [°C] | Crystallization concentration, [wt%] | T m b, [°C] |
---|---|---|---|---|---|---|
a 29 mol% ethylene-d4 (determined by IR spectroscopy). b The temperature control program for DSC was as follows: (a) heat the as-received polymer bulk sample from room temperature to 200 °C at a rate of 10 °C min−1. (b) Keep the polymer melt by holding the temperature at 200 °C for 5 min. (c) Cool at a rate of 10 °C min−1 down to −50 °C. (d) Heat the polymer sample from room temperature to 200 °C at a rate of 10 °C min−1. | ||||||
PE1 | PE-b-deuterated PEa | 3.8 × 106 | 1.23 | 105 | 0.0001 | 136.1 |
PE2 | Deuterated PE | 5.76 × 105 | 1.08 | 100 | 0.001 | 132.5 |
PE3 | PE | 3.5 × 106 | 7.5 | 106 | 0.0001 | 135.8 |
PE4 | PE | 7.5 × 105 | 7.5 | 100 | 0.001 | 135.4 |
Thus, some high-resolution AFM images of the morphology evolution after annealing at increasing temperature (Ta) and for prolonged times were performed ex situ at room temperature. For ex situ experiments, the single crystals were annealed first on a Linkam THMS 600 hot stage (Linkam Scientific Instruments, Tadworth, UK) under nitrogen atmosphere at a desired temperature for a chosen time (mostly 20 minutes). The sample was subsequently cooled to room temperature and investigated by AFM. The sample was further annealed on the Linkam hot stage at a slightly higher temperature and re-examined by AFM. This procedure was repeated several times. However, at room temperature the molten polymers re-crystallized and the AFM phase-signal was not able to distinguish them from the crystalline regions existing at elevated Ta. Thus, after a temperature protocol as the one shown in Fig. 1b including a quench to room temperature, AFM measurements could only visualize the morphology obtained after the annealing process but could not indicate the amount of molten polymers and their distribution within the single crystal.
Upon annealing at Ta > 122 °C, thickened parts became more prominent and grew from the periphery of the crystal in the direction towards the interior of the crystal. The resulting branches progressively carved regularly spaced valleys into the single crystal. Eventually, the branches reached the centre of the crystal, i.e. the initial nucleation site. When heating slightly above 130 °C, valleys were refilled by molten polymers. However, the memory of the branched patterns still existed as the branches re-appeared upon cooling (see also Fig. 3). Only after heating above ca. 135 °C, all memory of a crystalline pattern was lost. It is worth noting that annealing the crystal by jumping right away to temperatures above ca. 130 °C caused the formation of “Swiss-cheese”-like patterns rather than branches. Thus, slowly removing less perfect crystalline regions by providing enough time for re-organisation within the crystal seems to be required to generate periodic patterns formed within a single crystal.
In addition to identifying local morphological changes, we also determined the mean height of the crystal, recorded during annealing. These mean values allowed a direct comparison with results frequently obtained via averaging techniques like X-ray scattering.34,47 As shown in the histogram of Fig. 2f, taken from the recorded AFM images, a sequence of stages can be identified, consistent with previous observations.21 For the early stages (low Ta and/or short annealing times (ta)), the mean height of the crystal (its lamellar thickness (ld)) did not change much. Nonetheless, even at such early stages, the morphology already started to change at the crystal periphery. At progressively higher Ta, or after prolonged annealing, a stage of local melting and significant reorganization followed (Fig. 2c and d), causing a shoulder to appear at higher thickness values and evolving into a peak in the height histogram. With increasing Ta, this new peak, which is typically attributed to lamellar thickening,29,31,58,59 slowly shifted to higher values of ld. In combination with the direct space observation by AFM, we can conclude that the lamellar thickening process was accompanied by the formation of a regular pattern within the single crystal. For Ta > 130 °C, the height distribution became rather broad, related to the simultaneous melting of large parts of the crystal and thickening of the remaining parts, the branches.
Using the phase-signal of tapping mode AFM, it is possible to visualize differences in viscoelastic properties within an annealed single crystal. Under the imaging conditions used here, the contrast visible in the phase images reflects crystalline (light colours) and molten (dark colours) areas. As can be seen in Fig. 3a, the AFM phase images indicated that in the course of the appearance of the branches, they carved regularly spaced valleys into the single crystal which were partially filled by molten polymers as indicated by the change in the phase signal (the dark colour of the phase signal is indicating molten polymers).
The short waiting time of 10 min at Ta = 124 °C did not allow for the formation of branched structures within the whole crystal. After raising the temperature to 128 °C, only the thickened parts, i.e. the parts forming the branched morphology, exhibited a solid-like response in the AFM phase-signal. At even higher temperatures, such as 132 °C (Fig. 3c), the branched structure seemed to disappear. Under the chosen imaging conditions, it was not visible anymore in the AFM phase-signal. However, when increasing the tapping force,60,61 it became clear that branches still existed but were “hidden” under a layer of molten polymers. In addition, upon quenching to room temperature, the branched structure became clearly visible (compare Fig. 3f and i).
Upon cooling, all molten polymers re-attached to the remaining branches, making them visible again in the AFM topography images. A series of high resolution AFM high images of PE1 single crystals were taken at room temperature after annealing at various Ta (see Fig. 3g–i). When annealing at temperatures lower than 128 °C, the morphological evolution of the annealing-induced branched pattern was obviously very similar to what was demonstrated in the above in situ AFM images (compare Fig. 2b and c and 3b and e). Crack-like structures could be seen both in the central region and at the crystal edges after annealing at 128 °C. As shown in the AFM height image (Fig. 3h), the cracks formed a branched pattern. In addition, the lamellar thickness within such single crystals increased upon annealing.
Upon quenching the sample to room temperature, re-crystallisation of polymers from the partially molten state in between the branches contributed to the morphological changes. This became particularly visible for annealing temperatures above 128 °C. It can be deduced by comparison of Fig. 3d–f with Fig. 3g–i that re-crystallisation of the molten polymers was guided by the crystalline branched pattern.
Besides varying Ta, we also varied ta while keeping Ta constant. For example, at Ta = 125 °C, the morphological evolution of one solution-grown faceted crystal is shown in Fig. 4a. Already after ta = 2 min, numerous notches formed along the crystal edges. With increasing ta, these notches penetrated further towards the centre of the crystal. Elongated branches formed in each sector of the initial crystal, all of them pointing preferentially towards the centre (i.e. the nucleation site) of the crystal. As shown in the height histogram of Fig. 4b, a second broad peak emerged at the expense of the initial peak. The comparison of Fig. 2 and 4 reveals a kind of time–temperature superposition. Interestingly, the value of the periodic spacing (λ) of the branches increased with ta, reflected by a decrease in number of branches per unit length. This evolution of λ(ta) followed the same trend as the mean crystal thickness (ld) (Fig. 4c), consistent with previous scattering experiments34,36,47 which averaged over a large number of crystalline domains. Both, λ and ld increased approximately with the logarithm of ta.
In order to demonstrate the generality of our approach, we have repeated analogous experiments for a variety of PE molecules of high molecular weight, including commercial products (see Experimental section). For all samples, we grew single crystals having the shape of a truncated lozenge. Applying similar annealing procedures, we could generate analogously branched patterns for all studied samples. As shown in Fig. 5a for a constant annealing time of 20 minutes, λ and the length of the branches increased with Ta. Keeping Ta constant (Fig. 5b) and varying ta, a similar increase of λ was observed.
The initial lamellar thickness (ld0) depends on the growth rate, which, in turn, can be controlled by crystallization temperature Tc.35,62 In addition, the thickness (ld) of a crystalline lamella can vary during growth in the regions behind the growth front or during post-growth annealing stages. In Fig. 6, we demonstrate that λ also depends on Tc. This suggests that a relation between λ and ld0 (or ld) exists.
As has been shown previously,34,58 lamellar thickening induced by annealing is a complex process involving (1) a molecular diffusion process of the detached/molten molecules to more stable, thicker crystalline regions, (2) attachment and incorporation of molecules into these regions, followed by (3) further improvement of stability by removing even more folds and disappearance of less stable regions. Consequently, as shown in Fig. 2–6, annealing temperature and time have a crucial influence on the evolution of morphology and lamellar thickness (ld) during/after growth. The rate of such annealing-induced changes depended on the initial degree of order and the morphology of the single crystal, mainly determined by the crystallization temperature.
To shed some light on the origin of annealing-induced branched patterns, we searched for relations between the conditions under which the faceted polymer single crystals were grown and the periodic patterns observed after annealing. For various annealing procedures and crystallization conditions, we analysed a series of single crystals and compared the evolution of λ with the accompanying changes in the parameters ld0 (see Fig. 7a) or ld (see Fig. 7a). The results are summarized in Fig. 7, from which we can draw the following conclusions. λ was found to be much smaller than the size of the faceted crystals and only about ten times larger than ld. As shown in Fig. 4, for constant Ta, λ increased approximately with log(ta). Because ld also increased with log(ta), a clear relation between λ and ld was found. Single crystals prepared at increasing Tc showed an increase in ld0. As shown in Fig. 7b for annealing times increasing with Tc, λ also increased with Tc, giving rise to an approximately linear relation between λ and ld0. A more complex behaviour was found when keeping ta constant (20 minutes) and varying Ta (from 115 to 130 °C). At low Ta, λ increased rapidly with ld. But at high Ta, λ increased only slowly with ld. At low annealing temperatures (from ca. 115 to 120 °C), changes in morphology could only be observed for crystals grown at low Tc. At such conditions, λ was found to be small, i.e., the distances between branches were short, allowing for relatively efficient transport of polymer chains between branches. Thus, a coarsening of λ was observed even for short annealing times. For thicker lamellar crystals, higher temperatures and longer times were required to allow for the annealing-induced changes in morphology to become visible.
Footnote |
† Electronic supplementary information (ESI) available. See DOI: 10.1039/c4ra10563j |
This journal is © The Royal Society of Chemistry 2015 |