Random Copolymers from Polyamide 11 and Polyamide 12 by

Article pubs.acs.org/Macromolecules

Random Copolymers from Polyamide 11 and Polyamide 12 by Reactive Extrusion: Synthesis, Eutectic Phase Behavior, and Polymorphism Lien Telen,† Peter Van Puyvelde,‡ and Bart Goderis*,†,§ †

Polymer Chemistry and Materials Division, Chemistry Department, ‡Department of Chemical Engineering, Applied Rheology and Polymer Processing, and §Leuven Material Research Centre (Leuven-MRC), KU Leuven, B-3001 Leuven, Belgium ABSTRACT: Random copolymers from polyamide 11 (PA11) and polyamide 12 (PA12) were obtained from the parent homopolymers by transamidation reactions during high-temperature reactive extrusion. Concomitantly, the product molecular weight increased by postcondensation reactions. With both species being crystallizable, eutectic phase behavior was observed and explained in terms of the Flory theory for random copolymer melting. Close to the eutectic composition, both species crystallize in a competitive mode, which induces mesomorphic phase formation. Although the crystal centers most likely remain PA11 or PA12 pure, it is argued that foreign comonomer units might be tolerated at the crystal borders, at least for PA11 crystals. To understand the copolymer crystallization behavior in general, large amounts of rigid amorphous material need to be created alongside the crystallites. Because of their low melting point and since the crystallinity and mechanical performance typical of the parent homopolymers are preserved, random PA11/PA12 copolymers are suited for blending with thermally unstable (biobased) substances. decreasing amount of hydrogen bonds per unit of mass.1 Therefore, (biobased) polyamide 11 (PA11) and polyamide 12 (PA12), which have amide groups separated by 10 or 11 CH2 groups, belong to the lowest melting commercially available PAs, with melting temperatures of 185 and 177 °C, respectively. However, these melting temperatures still exceed the 150 °C target. The melting temperature of PA11 or PA12 can be reduced by acting on the polymer chain regularity. Acierno and Van Puyvelde11 reported on PA11 with short aliphatic side chains. At a branched comonomer content of 5% the PA11 melting point dropped by 4 °C, whereas the crystallinity reduced from 22 wt % for the linear species to 17 wt % for the branched counterpart. Introducing branches thus seems to be more effective in reducing the crystallinity than the melting point. A crystallinity reduction is expected since the branched comonomers are noncrystallizable units. Alternatively, the copolymerization with crystallizable units can be considered, such as the copolymerization of PA11 with PA12 units. PA11/PA12 random copolymers have been synthesized at different PA11/PA12 ratios by Johnson and Mathias.12 For a 65/35 PA11/PA12 ratio, they reported a melting onset of 148 °C and a crystallinity of 17 wt %. The crystallinity of this material is comparable to that of the mentioned branched PA11, but the melting point reduction is more pronounced.

1. INTRODUCTION Polyamides (PA), also known as nylons, are thermoplastics with a high strength, stiffness, and abrasion resistance, which make them industrially very relevant.1 The amide groups participate in hydrogen bonding which restricts the interchain mobility and therefore results in a high melting temperatures (180−350 °C) and high strengths.1 PA is often blended with other polymers to transfer its superior properties to the other component. Industrially, blending is commonly achieved via the melt. Many biopolymers, including biobased polymers such as poly(lactic acid) and polyalkanoates or natural biopolymers such as starch and proteins, suffer from being brittle or display a poor thermal stability and may therefore benefit from the addition of PA. Conversely, blending biopolymers to PA may increase its biocontent or facilitate biodegradation. By adding proteins, potentially other functionalities can be introduced, such as biocompatibility and controlled delivery.2 However, proteins and many other biopolymers do not withstand processing above the PA typical melting points.3 For example, wheat gluten has been considered as renewable resource for polymeric materials4−9 but starts degrading above 150 °C.10 To allow for melt blending, one thus necessarily has to use a low melting PA. To e.g. avoid protein degradation, 150 °C seems to be a useful target melting temperature. On the other hand, the PA should display the PA typical mechanical performance and crystallizability for it to be useful in the blend. The melting temperature of PAs decreases with increasing aliphatic chain length between amide groups, i.e., with a © XXXX American Chemical Society

Received: May 7, 2015 Revised: December 7, 2015

A

DOI: 10.1021/acs.macromol.5b00976 Macromolecules XXXX, XXX, XXX−XXX

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torque read-out display and logger. The extrusion temperature was varied between 230 and 400 °C, the extrusion time at 350 °C between 5 and 45 min, and the PA11/PA12 ratio at 350 °C for 30 min between 100/0 and 0/100. Blend compositions are denoted as PA11 wt %/PA12 wt %. A 40/60 blend thus contains 40 wt % PA11 and 60 wt % PA12. Gel permeation chromatography (GPC) measurements on the parent polymers and extrudates were performed on a 1260 GPC Infinity model (Agilent Technologies, Santa Clare, CA) with 1,1,1,3,3,3-hexafluoro-2-propanol (HFIP) (99.5%, across Organics, Thermo Fisher Scientific, Geel, Belgium) as solvent. The Zorbax PSM bimodal LC columns (Agilent Technologies, Santa Clare, CA) were kept at 30 °C. A 1260 refractive index detector (RID) (Agilent Technologies, Santa Clare, CA) was used, and the calibration was performed using PMMA standards. Potassium trifluoroacetate (SigmaAldrich, St. Louis, MO) was added (0.1 M) to the solvent to avoid PA aggregation.25,26 The hydrodynamic radius of PA11 decreases with increasing salt concentration as indicated in the literature with NaTFAc.21 Similar behavior is expected for PA12 and the copolymers by which a qualitative comparison of the GPC curves is allowed.25 The exact molecular weight will not be discussed because PA11 and PA12 behave differently from the PMMA standards used during calibration.27 PAs were dissolved in HFIP (1 mg/mL) at 40 °C for 2 h and left to stand for 24 h at room temperature prior to being measured. Before the actual GPC measurement, the solution was filtered over a Millex-FH filter (PTFE, 45 μm) (Millipore, Billerica, MA). Solutions were injected with a 1 mL glass syringe (Fortuna optima, Poulten and Graf GmbH, Wertheim, Germany). During data processing a linear baseline was subtracted prior to normalizing the RI signal to the integral of the peak of interest. Differential scanning calorimetry (DSC) heating, cooling, and second heating runs were recorded using a Q2000 heat-flux DSC (TA Instruments, New Castle, DE), calibrated with indium and sapphire standards and using hermetically closed aluminum pans. Heating and cooling were performed at 10 °C/min between 0 and 230 °C with isothermal segments of 5 min in between. Higher heating rates were used to demonstrate melting−recrystallization−remelting effects. The heat flow data were converted into specific heat capacities (J g−1 K−1) by subtracting an empty pan signal and normalization to the sample mass and scanning rate. The heat of fusion, ΔH (J g−1), was determined by integration of the signal enclosed between the experimental curve and a linear extrapolation form the melt signal. Normalization of ΔH to the heat of fusion for 100% crystalline material, ΔHref, yields a value for the material crystallinity. ΔHref for PA11 and PA12 were taken from the ATHAS database:28 i.e., 244 and 245 J g−1 for PA11 and PA12, respectively. An identical modus operandi was upheld for the copolymers. The reference heat of fusion for each copolymer was approximated by calculating the weight fraction weighted sum of the pure PA11 and PA12 ΔHref values. Time-resolved synchrotron wide-angle X-ray diffraction experiments were performed at the Dutch-Belgium beamline (DUBBLE) BM26B at the European Synchrotron Radiation Facility (ESRF, Grenoble, France) at a fixed wavelength of 1.03 Å, using a Pilatus 300 K-W detector (Dectris, Baden, Switzerland). A polyethylene standard was used for angular calibration. WAXD patterns were accumulated in consecutive frames of 4 s with 2 s of dead time while heating and cooling at 10 °C/min from 0 to 230 °C in hermitically closed DSC pans using a HFS 191 heating/freezing stage (Linkam Scientific instruments, Surrey, UK). Data were azimuthally averaged using homemade software29 and corrected for the empty sample holder signal, taking into account the sample transmission and direct X-ray beam intensity, measured by an ionization chamber placed downstream from the sample. WAXD patterns are represented as a function the scattering angle, 2θ, converted as if a wavelength of 1.54 Å was used (Cu Kα radiation), to facilitate comparison with published WAXD patterns. The angular range in these experiments relative to using Cu Kα radiation covered 10.8° < 2θ < 37.6°. Complementary static laboratory wide-angle X-ray diffraction experiments were performed at room temperature on samples that were wrapped in aluminum foil and cooled at 10 °C/min from 230 °C

Such low melting copolymers are of interest for the reasons mentioned above but are unfortunately not available commercially. To study the potential of PA11/PA12 random copolymers without having to resort to their synthesis from the constituting monomers, we explored their preparation from the parent homopolymers by means of reactive extrusion. Reactive extrusion uses high temperatures and extrusion shears to promote chemical reactions in polymer blends. PAs undergo transamidation reactions at elevated temperatures. Three types of transamidation reactions are possible: aminolysis, acidolysis, and amidolysis.13−16 These reactions are interchange reactions between the amine or carboxylic acid end groups and an amide link or between amide links mutually. Since water and acid influence the transamidation reaction,17,18 the mechanism most likely involves a hydrolysis of the amide group followed by cross-amidation.17 With increasing degree of transamidation during reactive extrusion the homopolymers are first transformed into block copolymers, later in segmented block copolymers, and ultimately into random copolymers.13,14,16,20 This progressive conversion into a more random structure leads to a progressive decrease in the melting and crystallization temperature. This paper describes the melting point depression and morphology of random copolymers from PA11 and PA12 obtained through high-temperature reactive extrusion. Direct proof of transamidation can normally be obtained by means of size exclusion chromatography (SEC)14 or nuclear magnetic resonance (NMR).19 PA11 and PA12 are, however, too similar for SEC, and the aliphatic chains are too long to distinguish the PA11 and PA12 1H NMR signals.21 The extent of transamidation, ultimately up to full randomization, was therefore traced via the differential scanning calorimetry (DSC) based material melting points. These melting points were compared to those of the random copolymers synthesized by Johnson and Mathias.12 The materials were characterized thermally, morphologically, and mechanically. The Flory theory for the melting point depression in random copolymers served at explaining the evolution of thermal transitions and semicrystalline morphology as a function of the copolymer composition.22−24 When using biobased PA11 as base polymer, the biobased content of potential blends with other biobased polymers remains high. Furthermore, PA11 as well as PA12 represents a valuable combination of typical nylon and polyolefin properties, e.g., low moisture absorption and density accompanied by chemical resistance similar to that of polyamide 6 with lower sensitivity to stress cracking.20

2. EXPERIMENTAL SECTION PA11 (Atochem, Serquigny, France; currently Arkema, Colombes France) and PA12 (Sigma-Aldrich Chemie, St. Louis, MO) were dried under vacuum for at least 4 h and stored inside an exicator up to use. From intrinsic viscosity measurements in m-cresol the PA11 weightaverage molecular weight, Mw, is 51 500 g/mol.11 No such information is available for PA12. PA12 was received as reactor powder and PA11 as extruded granulates. Thermogravimetrical analysis (TGA) on the parent polymers was performed using a TGA 2950 (TA Instruments, New Castle, DE) under a nitrogen atmosphere (60 mL/min) while heating at 5 °C/min from 25 to 550 °C. Different ratios of PA11 and PA12 were weighed, mixed in the solid state, and extruded at 100 rpm under a N2 atmosphere using a corotating fully intermeshing recirculating mini twin screw extruder of 5 cm3 (DSM Xplore, Geleen, The Netherlands), equipped with a B


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Macromolecules using the mentioned Linkam HFS 191 heating/freezing stage. These measurements were performed with an XeuSS X-ray camera (Xenocs, Sassenage, France), comprising a GeniX 3D molybdenum ultralow divergence X-ray beam delivery system (wavelength, λ = 0.71 Å) at a power of 50 kV−1 mA, a collimating assembly based on scatterless slits, a sample stage, a He flushed flight tube, and a Mar345 image plate detector (MARresearch, Norderstedt, Germany). The scattering angles were calibrated using silver behenate and polyethylene, and the patterns were processed following a method similar to that of the synchrotron WAXD experiments. Transmissions were obtained by measuring the direct beam intensity with a photodiode placed downstream from the sample. The angular range in these experiments relative to when Cu Kα radiation would have been used covered 2.4° < 2θ < 34.6°. Samples were also investigated by polarized optical microscopy (POM) using a BHS microscope (Olympus, Tokio, Japan) in conjunction with a TMS 600 hotstage (Linkam Scientific Instruments, Surrey, UK). Samples were cooled from the melt to room temperature at 1 °C/min. Pictures were taken every 0.5 °C using a TK-C1381 color camera (JVC, Kanagawa, Japan) interfaced with Leica Qwin software (Solms, Germany). Finally, the extrudates were injection molded into standardized dumbbell-shaped tensile bars (90 mm, 30 mm gauche length × 5 mm × 1.5 mm) using a microinjection molding machine (DSM Xplore, Geleen, The Netherlands) with a barrel capacity of 2.45 cm3 with the reservoir put 15 °C above the polymer end melting point, the mold at 120 °C, and using a pressure of 9 bar. The tensile bar mechanical properties were tested using an Instron 5985 tensile tester (Norwood, MA) (1 kN force cell) at 10 mm/min up to sample fracture. Initial sample width and thickness were measured using a caliper. The initial grip separation was 37 mm for all experiments. A speckle pattern was spray painted on the sample. During deformation a camera (Limes, Messtechnik and Software GmbH, Krefeld, Germany) took pictures every 0.5 s. The axial strain developed during deformation was obtained from the time evolution of the speckle pattern using a digital image correlation algorithm (VIC-2D software, Correlated Solutions, Inc., Columbia, SC) using a subset of 23 and a step size of 3, working incrementally. The mechanical properties were calculated based on the surface strain averaged over the entire sample area, omitting the areas that were less than 1.5 times the width away from the clamps. Engineering strain and stress at break were defined as the maximum values just before fracture and the modulus was extracted from the slope in the engineering stress−strain data between 0% and 2% strain. All measurements were performed at least in triplicate, and average values with their standard deviations are reported.

a N2 atmosphere. To avoid severe degradation, extrusion should thus not happen at temperatures above 350 °C. The extruder torque steeply increased to a certain value during filling. This value decreased with increasing temperature as illustrated for the 50/50 PA11/PA12 blend in Figure 2 for a

Figure 2. Torque evolution during extrusion of the PA11/PA12 50/50 blends at the indicated temperatures.

selection of extrusion temperatures. This decrease reflects the melt viscosity reduction with temperature. However, after the steep filling increase, a steady increase in torque was observed with time. As the extruder screw rotation speed remained constant at 100 rpm, the torque increase must be coupled to a melt viscosity increase. Such an effect was seen for the blends as well as for pure PA11 and PA12 and can be related to an increase in molecular weight by postcondensation reactions.14 The torque increase is faster with increasing temperature, suggesting a faster postcondensation at higher temperatures. However, at 400 °C, the initial torque was very low and tended to decrease rather than to increase. As typical odors of degradation escaped the extruder, extrusion was stopped after 10 min. In fact, given the TGA results, degradation, in particular of PA12, was to be expected. The normalized GPC traces for PA11 and PA12 prior to and after extrusion at 350 °C for 30 min are displayed in Figure 3. Extrusion results in a tailing toward smaller retention volumes (higher molecular weights) for PA11 whereas a broadening is observed for PA12. In both cases, the polydispersity increases. It can be concluded that postcondensation is occurring for both PA11 and PA12 while a limited amount of lower molecular weight chains is formed for PA12, likely as a result of degradation. Note that 350 °C is the borderline temperature for PA12 degradation according to the TGA experiment. Figure 4 contains the GPC signals of the extruded 50/50 blend at 350 °C for 30 min and at 400 °C for 10 min. Within experimental error, the blend GPC profile of the former is identical to the average of the pure PA11 and PA12 profiles. The profile of the blend extruded at 400 °C occurs at significantly higher retention volumes, i.e., at lower molecular weights due to sever degradation as expected from the TGA result. From the GPC experiments it can be concluded that very little degradation occurs at temperatures up to 350 °C for an extrusion time up to 30 min. Using these extruder conditions on pure PA11 and PA12 results in altered thermal behavior compared to when not extruded as illustrated in Figure 5. The crystallization onset temperatures are unaffected, but the crystallization peak temperatures shift to considerably (PA12) or slightly (PA11) higher temperatures.

3. RESULTS AND DISCUSSION At first the thermal stability of the parent polymers PA11 and PA12 was investigated using TGA, the result of which is shown in Figure 1. The 1% weight reduction level occurs at 353 °C for PA12 and at 398 °C for PA11 while heating at 5 °C/min under

Figure 1. TGA weight evolution of PA11 (solid line) and PA12 (dashed line) under a N2 atmosphere while heating at 5 °C/min. C


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Figure 3. GPC results of PA11 (A) and PA12 (B) prior to (full lines) and after (dashed lines) extrusion at 350 °C for 30 min.

Figure 4. GPC results of PA11/PA12 50/50 extruded at 350 °C for 30 min (dotted line) and the 50/50 blend extruded 10 min at 400 °C (dashed line). The full line represents the average of the GPC profiles of pure PA11 and PA12 extruded at 350 °C.

Figure 6. POM images at selected temperatures during cooling at 1 °C/min for PA11 prior to (A) and after extrusion at 350 °C for 30 min (B). The images (C) and (D) are the counterparts of (A) and (B) for PA12. The size of the square micrographs is 550 × 550 μm.

The DSC endothermic melting signals during subsequent heating for extruded as well as nonextruded PA11 and PA12 are bimodal (Figure 5). Experiments at different heating rates, as illustrated for extruded PA11 in Figure 7A, demonstrate that this is related to melting−recrystallization−remelting (MRR) phenomena as the second endothermic peak decreases in favor of the first one with increasing heating rates. Recrystallization is known to be suppressed when higher heating rates are used.30 The recrystallization of nonextruded PA12 is accompanied by cold crystallization since a net exothermic signal is seen prior to the onset of global melting (Figure 5). Recrystallization of the extruded samples is rather limited. These materials crystallized at higher temperatures into more perfect crystals, which have lower tendency to recrystallize during heating. Second, the recrystallization rate may have reduced because of the molecular weight increase by postcondensation. The impact on the thermal behavior of altering the extrusion temperature for pure PA11 and PA12 is moderate. In contrast, a drastic effect is seen for the polymer blends as illustrated for the 50/50 blend in Figure 8. At low extrusion temperature two

Figure 5. Cooling (bottom) and second heating (top) DSC runs of PA11 and PA12 prior to (full lines) and after extrusion at 350 °C for 30 min (dashed lines). The PA11 data are shifted by 5 J g−1 K−1 for clarity.

The POM images in Figure 6 demonstrate that after extrusion a higher number of smaller spherulites is obtained at a given temperature compared to when cooled at 1 °C/min from the melt without extrusion. The spherulites of nonextruded PA12 at 150 °C are grouped in patches, whereas more and evenly distributed spherulites are seen at that temperature for extruded PA12. The spherulites for PA11 after extrusion are almost too small to be observed. Extrusion thus enhances primary nucleation, likely as a result of dispersing nucleating foreign particles that were already present in the materials. The limited effect for PA11 probably stems from the fact that it was received as extruded granulate, whereas PA12 was a reactor powder. D


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Figure 7. DSC heating runs at 10 °C/min (dashed line), 20 °C/min (gray solid line), or 40 °C/min (black solid line) after cooling at 10 °C/min. PA11 extruded at 350 °C for 30 min (A), the PA11/12 50/50 physical blend extruded at 230 °C for 30 min (B), and the PA11/12 30/70 (C), 40/60 (D), 50/50 (E), and 70/30 (F) extruded at 350 °C for 30 min.

mirrored in the heating runs, where the two melting maxima approach but do not merge even after extruding at 350 °C. Figure 7E compares the DSC heating curves at 10, 20, and 40 °C/min after cooling at 10 °C/min for the sample that was extruded at 350 °C and demonstrates that in this case too MRR is at the origin of the double melting behavior: i.e., the second melting peak disappears when increasing the heating rate. In contrast, the double melting behavior remains for the sample extruded at 230 °C, with the split between the two peaks, however, being much reduced (Figure 7B). It thus seems that the double melting behavior for this sample is a reflection of the double crystallization behavior but reinforced by MRR. The progressive change in thermal behavior is interpreted as the progressive change from a physical blend of the two homopolymers at low extrusion temperature (with the two components crystallizing and melting independently in separate DSC maxima), over a segmented copolymer stage at intermediate extrusion temperatures toward a random copolymer formation at the highest extrusion temperature (with the polymer specific segments crystallizing in a single exothermic event). The change in behavior is most pronounced at temperatures above 250 °C, suggesting that a temperature of at least 270 °C is needed to efficiently exceed the activation energy for transamidation reactions. Furthermore, the time allowed during reactive extrusion is important. Varying the extrusion time for the 50/50 blend at 350 °C gradually from 5 min to ultimately 45 min makes the crystallization peak

Figure 8. Cooling (bottom) and subsequent (second) heating (top) DSC runs of PA11/PA12 50/50 after extrusion for 30 min at varying temperatures. All the curves are shifted incrementally with 2 J g−1 K−1 for clarity except for the curves related to the sample extruded at 230 °C. Right side temperature labels refer to the extrusion temperature for that particular sample.

exothermic crystallization maxima are observed, which upon increasing the extrusion temperature merge into a single peak at considerably lower temperatures. This behavior is only partially E


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Figure 9. Time-resolved WAXD data during cooling (lower panels) and subsequent heating (upper panels) at 10 °C/min for PA11 extruded at 350 °C for 30 min (A), PA12 extruded at 350 °C for 30 min (B), the 50/50 blend extruded at 230 °C for 30 min (C), and the same blend extruded to full randomization at 350 °C for 30 min (D). The WAXD patterns are top viewed, with the intensities represented using a gray scale. The different polymer specific polymorphic regimes are indicated in the WAXD patterns. The corresponding DSC curves are represented in the middle row with the cooling and heating runs plotted respectively down and upward.

temperature asymptotically drop from 144 °C at 5 min to a plateau at 133 °C after 30 min (data not shown). A steady state is thus obtained after 30 min extrusion at 350 °C. The evolution of the thermal behavior of blends with different composition but processed in the same way will prove that this state corresponds to having reached the random copolymer state (see further below). The simultaneous crystallization of the PA11 and PA12 segments for the 50/50 random copolymer does not imply that these segments cocrystallize into compound crystals. Evidence for their separate crystallization is provided via the timeresolved WAXD experiments. In Figure 9 the WAXD patterns of PA11, PA12, the physical 50/50 blend (30 min extruded at 230 °C), and the 50/50 random copolymer (30 min extruded at 350 °C) during cooling and subsequent heating at 10 °C/ min are compared. The DSC curves are shown in the middle panels adjacent to the corresponding WAXD data. At room temperature, crystalline PA11 can be organized as a triclinic α phase or a metastable pseudohexagonal δ′ phase.31,32 The α to δ′ phase ratio is a function of the crystallization temperature, with higher temperatures leading to higher α phase shares.33 Quench cooling leads to the metastable δ′ phase only. Literature agrees on the fact that the α phase is composed hydrogen-bonded molecular sheets, stacked side-by-side with a progressive shift.34 Disagreement exists on whether the chains within the sheets are arranged in a parallel34−37 or antiparallel fashion.31,38−40 For odd-numbered nylons (such as PA11) both arrangements allow for the creation of fully hydrogen-bonded sheets with the aliphatic chain fragments in a fully stretched, alltrans configuration.34 It can therefore be speculated that these arrangements coexist. On the other hand, the antiparallel arrangement might be preferred since this type of arrangement is compatible with the folding of chains within the sheets.34 WAXD powder patterns of melt crystallized samples do not allow discriminating between the two options. The pseudohexagonal δ′ phase is composed of a hydrogen-bonded network, rather than a stacking of hydrogen-bonded sheets.37 As the hydrogen bonds are configured in layers, parallel to the lamellar

surface, this network is occasionally referred to as a 2D network.41 At room temperature, the α crystal structure produces WAXD reflections at 20.5° and 23° 2θ Cu Kα, of which the former corresponds to the intersheet distance and the latter to the distance between chains in a sheet. Baeten et al.33 reported that the angular split between these reflections reduces when higher amounts of the δ′ phase are present. During heating, the α phase peak at 23° shifts to smaller angles as a result of thermal expansion until reaching the Brill transition temperature, where the α phase (reversibly) converts into the hexagonal δ phase, displaying a single reflection at intermediate angles.31 Above the Brill temperature the interchain distance, within a hydrogen-bonded sheet, and the intersheet distance remain equal.36 The δ′ structure produces a single peak at 21.1° 2θ Cu Kα in WAXD and does not (reversibly) convert into the α phase at low temperatures. This property allows discriminating the δ′ phase from the δ phase.37 Figure 9A shows that PA11 crystallizes into the δ phase and converts during further cooling at 10 °C/min into the α phase, slightly below 100 °C. During heating, this sequence is reversed with a Brill transition close to 100 °C. Little or no sign of the δ′ phase is present. An important diagnostic reflection for indicating the presence of PA11 crystallinity, particularly useful for the analysis of blends and random copolymers, is the 001 reflection at approximately 7° 2θ. This reflection is related to the monomeric repeat distance, irrespective of the polymorphic form.36 Complementary static WAXD experiments had to be conducted since the 001 angular range was excluded in the synchrotron WAXD experiments (see Experimental Section). Figure 10 represents the static room temperature WAXD patterns of PA11, PA12, a selection of random copolymers (30 min extruded at 350 °C), and the physical 50/50 blend (30 min extruded at 230 °C). The solid line through the experimental data (open circles) of PA11 in the left side panel is a Gaussian approximation of the 001 reflection, fitted to the experimental data using a least-squares procedure. The Gaussians’ full width F


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Table 1. Bragg Spacings (d001) and Crystal Sizes (D001) Based on the Gaussian Approximations to the PA11 and PA12 Shares in the 001 Reflections Depicted in the Left Side Panel of Figure 10; %PA11 and %PA12 Are the Relative Areas Occupied by Respectively the PA11 and PA12 (001) Peaks d001 [Å] PA11 PA11 PA11/PA12 copol PA11/PA12 copol PA11/PA12 copol PA12 PA11/PA12 blend

at half-height (in radians), β, was used to calculate the crystal size, D001, via the Scherrer equation: 0.9λ β cos θ

PA11

PA12

48.4 50.9

%PA11 %PA12

70/30

12.1 12.7

100 100

0 0

50/50

12.7

14.7

50.9

59.0

90.9

9.1

30/70

12.7

14.7

50.9

59.0

18.8

81.2

50/50

12.4

15.4 15.1

49.5

61.5 60.6

0 66.5

100 33.5

happens without progressive shear.45 The chains are in a planar zigzag, except for the methylene units close to the amide bonds, by which the latter are inclined by 60° with respect to the aliphatic planes.46,47 This twisting results in rather identical intersheet and intrasheet chain separation distances and hence a single reflection in WAXD at an angular position identical to that of the γ′ phase at room temperature.42 However, the γ and γ′ phase can be discriminated at higher temperatures, since upon heating the γ phase reversibly converts into the α′ phase whereas the γ′ phase does not. The α′ phase exhibits two characteristic WAXD reflections, but the angular split is not as pronounced as for the α phase.42,43 The PA12 α phase is only obtained under special conditions (high pressure, drawing)42 and is not relevant to the present work. PA12 in Figure 9B starts crystallizing in the α′ phase and converts into the γ phase during cooling. During heating the order is reversed. No polymorphic changes occur in the melting peak ranges of PA11 and PA12, confirming that their bimodal melting is associated with MRR, rather than with polymorphic transitions (see also Figure 7A). In addition, PA12 exhibits a low angle reflection at 5.7° 2θ Cu Kα as illustrated in Figure 10. Just like the PA11 001 reflection, this peak is related to the molecular repeat length. In fact, this reflection shouldat least for the γ phasebe indexed as 002 because the actual repeat along the chain direction in the PA12 γ phase involves two monomer units because of an alternating amide bond orientation.46 However, as this reflection also exists for the disordered γ′ phase and as for this phase a doubled molecular repeat is less obvious, we refer to this reflection as the PA12 001 reflection for both the γ and γ′ phase. The PA12 001 reflection is approximated less ideally by a Gaussian (Figure 10), but nevertheless the D001/d001 ratiojust like for PA11 reveals a lamellar thickness of four unit cells (Table 1). The PA12 lamellar thickness, D001, corresponds well to earlier reported values based on transmission electron microscopy.48 All the temperature-dependent crystalline features of pure PA11 can be found in the WAXD patterns of the 50/50 blend extruded at 230 °C (Figure 9C). For PA12 the γ phase typical reflection is present, but no α′ phase is seen at higher temperatures. This γ reflection should therefore be associated with the pseudohexagonal γ′ phase rather than with the γ phase. Given that the onset of crystallization in this physical blend of the homopolymers approaches that of pure PA11 (165 °C), and since pure PA12 only starts crystallizing at approximately

Figure 10. Static room temperature WAXD data after cooling at 10 °C/min from 230 °C for (from top to bottom) PA11, three random copolymers of the indicated composition, and PA12 extruded at 350 °C for 30 min. The bottom pattern corresponds to the 50/50 blend extruded at 230 °C for 30 min. The right side panel represents the full patterns whereas magnifications of the low angle 001 reflection range (represented using open circles) after a baseline correction are depicted in the left side panel. The magnified 001 reflections are fitted to the sum (full line) of PA11 (dotted line) and PA12 (dashed line) contributions, modeled as Gaussians. No PA12 share is present in pure PA11 and the PA11/PA12 70/30 copolymer whereas a PA11 share is only absent in PA12.

D001 =

PA12

D001 [Å]

(1)

with θ half the scattering angle at the peak maximum and λ the X-ray wavelength. This crystal size estimate corresponds to the lamellar thickness. The Gaussians’ peak position was used to calculate via Bragg’s law the monomeric repeat distance, d001. The least-squares fitting of the Gaussian approximation was constrained such that the ratio D001/d001 yields an integer value, representing the number of unit cells in the lamellar thickness direction. The results are collected in Table 1. The PA11 d001 value (12.1 Å) is identical to what has been reported in the literature32 but lower than the actual monomeric length (14.9 Å) because of chain tilting in the crystal lattice.36,40 Dividing D001 by d001 reveals that the PA11 lamellar thickness is composed of 4 unit cells. PA12 has at least four polymorphs (α, α′, γ, and γ’).42,43 The γ phase is the disordered, pseudohexagonal PA12 equivalent of the PA11 δ′ phase12 and is usually obtained via quench cooling.42 This phase exhibits a single reflection at 21.1° 2θ Cu Kα which persists during heating up to melting.44 The PA11 δ′ phase behaves in exactly the same way. The most stable phase for PA12 is the pseudohexagonal γ phase45 and is obtained by cooling more slowly from the melt.42 In this γ structure, hydrogen-bonded sheets of parallel chains are stacked alternatingly with sheets containing chains of the opposite direction. This structure implies that chain folds run from one sheet to another and not along a given sheet. Sheet stacking G


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PA11 which crystallizes first in all cases, by which confinement effects are only observed for PA12 via absence of the α′ phase and hence also the γ phase. Indeed, the split of the PA11 α phase reflection is as large as for pure PA11, pointing to very low concentrations of the PA11 δ′ phase in the physical blend. In summary, PA11 and PA12 seem to crystallize separately in the 50/50 physical blend as well as in the 50/50 random copolymer, but their polymorphism is influenced by earlier crystallization of the other component. However, arguments in favor of a (limited) cocrystallization of the different components in the copolymer case will be presented further below. The effect of altering the blend composition on the thermal behavior after extruding for 30 min at 350 °C is illustrated in Figure 11 for a selection of compositions. All materials display a

155 °C, it can safely be stated that PA11 is the first crystallizing species in this blend. It therefore seems that the presence of PA11 crystals in the close proximity of molten PA12 induces crystallization into the less perfect γ′ form, rather than into the PA12 α′ phase. The 001 reflections of PA11 and PA12 are both clearly visible in Figure 10, and both are compatible with a lamellar thicknesses of four unit cells. The crystallization behavior of this blend during cooling at 1 °C/min in POM looks very similar to that of pure PA11 (data not shown). This suggests that PA11 and PA12 are well mixed in the melt and that the segregation induced by the earlier PA11 crystallization does not happen at an interspherulitic level, at most at an interfibrilar level but most likely at the interlamellar (nm) level. The confinement of the PA12 enriched melt in between the PA11 crystals may impede the chain conformational changes, needed for crystallization in the α′ form. It surely reduces the PA12 crystallizability since the PA12 share (33.5%) in the overall crystallinity is only half that of PA11 (66.5%) assuming that the relative 001 peak areas are proportional to the shares in the overall crystallinity (see also Table 1). One can expect even stronger mutual influences during the crystallization of the (connected) PA11 and PA12 segments of the random copolymers. As expected, the crystallization behavior of this copolymer during cooling at 1 °C/min in POM also looks very similar to that of pure PA11 and suggests homogeneity of the system (data not shown). However, contrary to for the physical blend, crystallization of the 50/50 random copolymer involves formation of at least a small fraction of the PA12 α′ phase, as two reflections are seen in WAXD at the onset of crystallization (Figure 9D). The (rather faint) small-angle peak merges with the stronger high angle peak during cooling, indicating conversion into the γ phase. Upon subsequent heating the α′ phase appears again at about 115 °C. PA11 in this copolymer crystallizes in both the δ and δ′ phase since the split of the reflection at the conversion from the δ to the α phase is very limited, as expected for mixed δ and δ′ phases.33 During heating, the inverse behavior is observed. However, since the δ′ and δ WAXD signatures at high temperatures cannot be discriminated, it cannot be excluded that the δ′ form transforms into the δ phase during heating. To emphasize this uncertainty, the (δ) label appears in Figure 9D. As both the δ and δ′ phases are formed for PA11, it is possible that besides α′ also PA12 γ′ crystals are created. The latter can unfortunately not be discriminated from the PA11 δ′ or δ crystals at high temperatures or form the PA12 γ crystals at low temperature. The analysis of the 001 reflection area of this copolymer (Figure 10) suggests that only 9.1% of the total crystallinity is due to PA12. Since PA12 α′ crystals are visible in Figure 9D, the amount of PA12 γ′ crystals might be very small. Note that the 001 peak widths keep on revealing PA11 and PA12 crystal thicknesses of four unit cells. The data in Table 1 furthermore reveal that the PA11 001 reflection for the 50/50 random copolymer appears at slightly lower angles (d001 = 12.7 Å) compared to for the pure PA11 (d001 = 12.1 Å). This might be related to the dominant presence of the δ′ phase over the α phase as Nair et al. reported earlier that the δ′ phase d001 value is larger than that of the α phase.36 It is suggested that for the 50/50 random copolymer PA11 as well as PA12 can crystallize (locally) first, by which constraints are put on the crystallization of the remaining species. When PA11 or PA12 crystallizes first, this happens in the γ or α′ form, respectively. When crystallizing as second component, the confinement induces formation of the γ′ and δ′ phases. For the physical blend, it is

Figure 11. Cooling (bottom) and subsequent heating (top) DSC runs for randomized PA11/PA12 blends after extrusion for 30 min at 350 °C. The amount (wt %) of PA11 in the blends is indicated at the right side of the curves. The curves are shifted incrementally with 5 J g−1 K−1 for clarity except for the curves related to the 20% sample. Dashed and full lines are used alternatingly for better discrimination of the curves.

single exothermic crystallization peak and a bimodal endothermic melting peak due to MRR as illustrated for a series of copolymers in Figures 7 (the second melting peak reduces in favor of the first one with increasing heating rate). The crystallization onset temperature passes a minimum value for a PA11 wt % close to 40% as illustrated in Figure 12 (black squares). The same trend is seen for the melting points (black spheres) and resembles the evolution of the liquidus lines in an eutectic phase diagram of binary melt miscible, crystallizable systems with rather similar melting points of the two pure components.30 In (random) copolymers, the noncrystallizing comonomer units assume the role of the solvent and are at the origin of the progressive melting point depression with increasing comonomer content. In such eutectic phase diagrams, the melting point to be reported for the liquidus is the temperature at the end of the broad melting endotherm because only at that point the composition of the melt approaches the copolymer composition. The melting points in Figure 12 are end melting points, except for the melting points of the pure components, which were taken at the DSC peak maxima. The copolymer melting points are the H


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dashed line) was calculated with the ATHAS recommended ΔH (48 400 J mol−1) and T0m (500 K) values.30 The theoretical liquidi are situated above the experimental melting points because polymers tend to crystallize at a given degree of supercooling into metastable lamellar crystallites with melting points lower than those of the infinitely large equilibrium crystals assumed in the Flory formalism. Heck et al. derived that the melting point depression associated with the presence of comonomer and the depression related to having lamellar crystals with finite thickness are simply superimposed.49 Calculating the D001/d001 ratio from the data in Table 1 for the 30/70 and 70/30 PA11/PA12 random copolymers after cooling at 10 °C/min again leads to four unit cells across the lamellar thickness just like for the pure polymers and the 50/50 copolymer. Given that the lamellar thicknesses are identical and following Heck et al.,49 eq 2 should be capable of describing the melting temperatures of the experimental data, provided the equilibrium melting point is replaced by the melting point of the pure components with four unit cells thick lamellar crystals. Accordingly, eq 2 was fitted to the experimental end-melting points by a least-squares minimization procedure in which T0m and ΔH were variables. To adequately describe the PA12 data points (thick dashed curve in Figure 12), the PA12 ΔH value needed to be decreased from 48 400 to 32 921 J mol−1. For the PA11 rich side, a good fit was obtained while keeping the ATHAS ΔH value (thick full line in Figure 12). Furthermore, T0m was reduced to 175.6 and 184.3 °C for PA12 and PA11, respectively. The PA12 melting point is lower than that of PA11 as expected because of the lower amount of hydrogen bonds per unit of mass for PA12.1 The reason for a reversed order in the ATHAS T0m values (493 K for PA11 and 500 K for PA12) is unclear. Logically, the stronger T0m reduction for PA12 compared to for PA11 makes the eutectic point, i.e., the point at which the liquidi intersect, shift from 53.5 wt % PA11 for the equilibrium lines to 38 wt % PA11 for the experimental ones. Part of the shift is also due to the PA12 ΔH reduction. The thin dotted line over the PA12 rich side of the melting points represents eq 2 with the PA12 ΔH fixed at the ATHAS value and an adjusted T0m such that the melting points close to pure PA12 are well covered. This line intersects the PA11 liquidus at 42 wt % PA11, which is higher than 38 wt %. The shift of the eutecticum to lower PA11 wt % also exists in the crystallization onset values (black squares in Figure 12). The PA12 crystallization liquidus based on eq 2 with the ATHAS ΔH value but with T0m lowered to cover the pure PA12 crystallization onset (dotted line that runs over the black squares) strongly deviates from the experimental PA12 crystallization onsets close to the eutecticum. This mismatch is even larger than for the melting points. The ATHAS ΔH based PA11 crystallization liquidus for PA11 wt % < 50% also (slightly) exceeds the experimental PA11 crystallization onset temperatures. The deviations close to the eutecticum suggest that crystallization is hampered in this eutectic region and therefore occurs at a larger than expected supercooling. Crystallization at larger supercooling often implies the production of less stable crystallites, which may melt and recrystallize during subsequent heating. MRR in the 50/50 (Figure 7E), 40/60 (Figure 7D), and 30/70 (Figure 7C) PA11/PA12 copolymerswhich are compositions close to the eutecticumis indeed quite pronounced. Further away from the eutecticum, MRR is limited as e.g. for the 70/30 copolymer (Figure 7F). In general for the PA11 rich side of the diagram,

Figure 12. Experimental and theoretical transition temperatures of PA11/PA12 random copolymers with varying PA11 wt %. The upper lines represent the equilibrium copolymer liquidi for PA11 (thin solid line) and PA12 (thin dashed line). The lines through the experimental copolymer melting points (full circles) are fits based on eq 2 with the thick dashed curve for PA12 rich copolymers and the thick full line for the PA11 rich side. The melting onsets, as reported by Johnson and Mathias,12 are shown as open circles. The experimental onsets of crystallization are depicted using full squares. The thin dotted lines represent transition temperatures based on eq 2, ATHAS values for ΔH, and adjusted T0m values (see text).

DSC based end melting points, reduced with the temperature difference between the peak and end melting point of the pure components. The PA12 rich wing of the diagram was corrected with the PA12 difference and the right side wing with that of PA11. Following Berghmans,30 this procedure accounts for the intrinsic polymer melting point distribution. The open circles in Figure 11 are the melting peak onsets reported by Johnson and Mathias12 for the random PA11/PA12 copolymers they synthesized. Since the degree of randomization by transamidation cannot be traced by chemical analysis (cf. Introduction), the similarity between the melting points by Johnson and Mathias12 and the copolymers obtained by extrusion for 30 min at 350 °C is taken as evidence for full randomization by extrusion. The remaining difference obviously results from comparing end melting points with onsets. Johnson and Mathias12 did not report end melting temperatures or DSC profiles of their materials. The Flory theory22−24 describes the melting point depression in a copolymer with respect to that of the homopolymer. For a copolymer consisting of A units that crystallize and B units that do not (under the chosen experimental conditions) the copolymer equilibrium melting temperature, T0m,c (K), can be calculated based on the equilibrium melting temperature of the pure polymer, T0m (K), its heat of fusion per repeating unit, ΔH (J mol−1), and the A homosequence propagation probability, PAA, using 1 1 R = 0 − ln PAA 0 ΔH Tm,c Tm (2) with R the gas constant (8.31 mol K−1 J−1). For a random copolymer PAA is equal to the mole fraction of A units. Using the ATHAS recommended ΔH (41644 J mol−1) and T0m (493 K) for PA11, the PA11 equilibrium copolymer liquidus was calculated based on eq 2 and represented in Figure 12 using a thin full line. Similarly, the PA12 equilibrium liquidus (thin I


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Figure 13. Scenarios for the crystallization of the PA11/PA12 50/50 (A), 30/70 (B), and 70/30 (C) random copolymers. The arrows and number labels are explained in the text. The melting liquidi at the PA12 rich (thick dashed line) and PA11 rich (thick solid line) side up to the apparent eutecticum are extrapolated beyond the eutecticum using dotted lines. The thin lines represent the crystallization liquidi. Panel D contains the composition range enclosed by the gray box where crystallization of one component may trigger crystallization of the other isothermally. The black squares and spheres have an identical meaning as in Figure 12. The open squares represent the crystallization peak maxima of the different copolymers during cooling at 10 °C/min.

only reflects the supercooling, not the state to which the system tends to evolve. Coincidentally, point 4 also lays on the PA12 crystallization liquidus, implying that PA12 can crystallize from the PA12 enriched melt (point 5), thereby leaving a PA11 enriched melt with a composition given by point 6 on the PA12 melting liquidus. Point 6 occurs below the PA11 crystallization liquidus, by which the systems falls into a cascade of alternating PA11 and PA12 segment crystallizations. In theory, the system can fully crystallize isothermally, similar to what is possible at a regular eutectic temperature. However, polymers never do so and get stuck in a semicrystalline state. The successive melt enrichments take place at the crystallization front and cannot diffuse out in the remaining bulk melt because of the segmental connectivity. These enrichments therefore crystallize close to the growth front of the other species. As a result, solidification happens heavily disturbed, at larger supercooling andas discussed abovelarger fractions of the γ′ and δ′ polymorphs are formed. Moreover, there are also good arguments for the creation of a large fraction of rigid amorphous material. At point 4 the melt composition equals 24% PA11. For obtaining such an enrichment in PA12 segments, starting from a 50% composition, the wt % of crystalline PA11 segments at point 3 should be 34% of the total mass. This value by far exceeds the DSC based total crystallinity of this material, which is close to 20 wt %, as illustrated in Figure 14.

differences in crystallization induced morphology seem to be largely wiped out by MRR so that ultimately the PA11 ATHAS ΔH value can describe the end melting point evolution. In contrast, for the PA12 rich copolymers, MRRalthough quite pronouncedseems to be insufficient to wipe out morphological differences. The low experimental PA12 ΔH value might thus be an artifact, and the experimentally observed eutecticum should not be considered as an equilibrium point.32 For the same reason, the ATHAS ΔH values were preserved for both PA11 and PA12 to calculate the DSC based crystallinities (after conversion to J g−1) and not replaced by the values found by fitting eq 2 to the melting points. A qualitative explanation for the disturbed crystallization close to the (apparent) eutectic point can be inferred from the eutectic phase diagram. In Figure 13, the (melting) liquidi are extrapolated using dotted lines and the crystallization onsets are connected by thin full lines to represent the crystallization liquidi. Crystallization peak temperatures (open squares) are also added to show the temperatures of main crystallization for the different copolymer compositions. When cooling the 50/50 copolymer from the melt (point 1 in Figure 13A) to point 2 at the PA11 crystallization liquidus, the PA11 sequences crystallize (point 3) and leave a PA12 enriched melt with a composition given by the PA11 liquidus at point 4. Indeed, the composition of the segregated melt is given by the melting point liquidus since the crystallization liquidus J


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Figure 14. DSC (black squares) and literature based12 crystallinities (open spheres) as a function of PA11 wt % in PA11/PA12 random copolymers. The light and dark gray vertical bars represent the PA11 and PA12 shares to the total crystallinity based on the analysis of the WAXD 001 reflection (see Table 1).

Put differently, from all PA11 segments 68% should be in the crystalline state, which also is much larger than the crystallinity of pure PA11 (Figure 14). However, semicrystalline aliphatic polyamides are known to contain a large fraction of rigid amorphous material,50 and it is suggested to add this rigid amorphous fraction to the segregated solid fraction. This rigid amorphous fraction, however, does not contribute to the melting enthalpy and hence also not to the DSC based crystallinity. For polyamide 6 rigid amorphous fractions have been reported that are as large as or even larger than the crystalline fraction.51 Whether this rigid amorphous fraction exists at the crystal interfaces51 or is buried within the crystallites52 remains a matter of debate. In any case, the rigid amorphous fraction is associated with the crystallites, implying that it can be added to the segregated solid fraction. Conversely, the need of enriching the melt with PA12 segments to allow for its crystallization together with the observation that the PA12 segments actually crystallized (see Figure 9D and related discussion) can be taken as (indirect) evidence for the existence of (segregated) PA11 rigid amorphous material at quantities as large as the PA11 crystalline content. Crystallization of the 30/70 PA11/PA12 copolymer (Figure 13B) is comparable to that of the 50/50 copolymer. In this case the PA12 crystallization liquidus is reached (point 2) after cooling from the melt (point 1). Part of the PA12 segments crystallize (point 3) and leave a PA11 enriched melt with a composition given by point 4. This point lays below the PA11 crystallization liquidus by which the PA11 segments also crystallize and leave a PA12 enriched melt at point 5. Given its position below the PA12 crystallization liquidus, the PA12 segments can crystallize again and so on. In Figure 15A, the synchrotron WAXD patterns are collected for the 30/70 PA11/ PA12 copolymer. The PA12 α′ and γ phases are clearly present. PA11 crystallizes in the δ′ phase as no traces are seen of the α phase. Analysis of the 001 reflection area in Figure 10 and the related data in Table 1 for this copolymer reveal that 18.8% of the total crystallinity at room temperature is due to PA11 (δ′) crystals. Most likely the PA12 γ′ phase is present as well, but this cannot easily be revealed. Crystallization of the 70/30 PA11/PA12 copolymer, which is far away from the eutectic zone, is different. When cooling this copolymer from a high temperature melt (point 1 in Figure 13C) to the PA11 rich crystallization liquidus at point 2, a

Figure 15. Time-resolved WAXD data during cooling (lower panels) and subsequent heating (upper panels) at 10 °C/min for the PA11/ PA12 30/70 (A) and 70/30 (B) blend extruded to full randomization at 350 °C for 30 min. The WAXD patterns are top viewed, with the intensities represented using a gray scale. The different polymer specific polymorphic regimes are indicated in the WAXD patterns. The corresponding DSC curves are represented in the middle row with the cooling and heating runs plotted respectively down and upward.

crystalline PA11 phase is generated (point 3) and a PA12 enriched melt with composition given by point 4. In contrast to for the 50/50 and 30/70 copolymers, crystallization cannot proceed isothermally. The system needs to be cooled down. In principle, little or no supercooling would be needed to trigger further PA11 crystallization, given that PA11 crystals are present as nuclei. In that scenario, the melt would progressively enrich in PA12, following the PA11 liquidus down to reaching point 5 at the PA12 crystallization liquidus where PA12 can crystallize. However, the system cannot arrive at point 5. To reach a melt composition of 24 wt % PA11 segments at point 5, the PA11 solid phase at point 6 should represent 60% of the total product mass. Put differently, from all PA11 segments in the system 86% should be in the solid state. This seems to be over the limit, even if rigid amorphous material is supposed to contribute to the segregation. Therefore, in the 70/30 copolymer only PA11 crystallizes, and the PA12 enriched melt is left in the amorphous state. This is confirmed in Figure 10 and Table 1 where it is shown that the 001 reflection does not contain any contribution from PA12. Figure 15B illustrates that this copolymer crystallizes into the PA11 δ phase, which converts into the α phase upon cooling. The reversed process is seen during subsequent heating. There is no trace of the PA12 α′ phase. The narrower split of the α phase reflections suggests a larger fraction of δ′ phase compared to in pure PA11. This may result from crystallization at larger supercooling.33 It cannot be detected whether or not this δ′ phase recrystallizes during heating into the δ phase. In Figure 13D, the compositional range is highlighted where crystallization can proceed isothermally but heavily disturbed via the cascades sketched for the 50/50 and 30/70 copolymers. The upper left corner of the gray box coincides with the crossing point of the PA11 melting liquidus and the PA12 K


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compositions are represented by the light and dark gray bars, respectively. These bars are obtained by imposing the WAXD based shares (see Table 1) to the DSC overall crystallinity. As lamellae of 4 unit cells are generated and since only 34% of the PA11 sequences are long enough to contribute to such crystals in the PA11/PA12 50/50 copolymer (which corresponds to 52 mol % PA11), virtually all sufficiently long PA11 sequences need to crystallize to reach the 90.9% share in the total crystallinity, which is highly improbable. In the next paragraph, the thesis is defended that PA12 sequences at the borders of PA11 crystals are tolerated and can contribute to the PA11 crystallinity whereas the inclusion of PA11 sequences in PA12 crystals is less likely. Figure 16 represent a sketch of a PA11 hydrogen-bonded sheet in the α structure with a thickness of 4 unit cells. The blue

crystallization liquidus. Its upper right corner occurs at the PA11 crystallization liquidus where a horizontal line drawn from the upper left corner hits it. In the compositional range enclosed by the gray box, the ATHAS ΔH based crystallization liquidi deviate most markedly from the experimental ones (see Figure 12). Note that crystallization onset and peak values of the PA11 rich copolymers in this range are very close, indicating sharp crystallization peaks (a circle is drawn around these three points in Figure 13D). The remarkably sharp crystallization exothermic peak can be seen for the 40 wt % sample in Figure 11 and indicates that crystallization indeed can proceed isothermally via the sketched cascade. The 50/50 sample is a borderline case of which only the crystallization peak falls within the gray zone of Figure 13D. The crystallization peaks pertaining to the PA12 rich side in the gray box are not that narrow. A possible reason could be that the PA12 crystallization is slowed down when the γ′ crystallization interferes with the α′ crystallization. The concurrent growth of the PA11 δ and δ′ phases also has been suggested to slow down the overall crystallization rate,33 but in contrast to the α′ and γ′ phase, the δ and δ′ phase may nucleate each other mutually because of their identical hexagonal symmetry. It can be speculated that the inability to do so contributes to slowing down the PA12 crystallization rate more strongly than in PA11 such that crystallization happens at larger supercooling into more defect crystals, which in turn may lead to the deviations seen in the liquidi slopes for PA12 (Figure 12). In Figure 14, the DSC based crystallinities of all copolymers are collected. For comparison, the crystallinity estimates by Johnson and Mathias12 were added to this figure. The fact that the values of synthesized random copolymers are very comparable to the ones obtained via transamidation reactions underlines once more that full randomization was likely achieved. The crystallinities of pure PA11 and PA12 are comparable with values reported by others.44,53 The crystallinity of random PA11/PA12 copolymers remains relatively high. In random copolymers of other crystallizable polymer pairs the crystallinity reduction often is much more drastic, as e.g. in copolymers of ethylene with propylene, where incorporating 33 wt % propylene in polyethylene is able to make it amorphous.54 The preservation of crystallinity is closely related to the fact that the PA12 and PA11 lamellae are only four monomeric units thick. In order to build crystals from e.g. ethylene copolymers, the ethylene sequences need to be much longer than 4. Sufficiently long sequences for crystallization are readily absent when both monomers are present in close to equal amounts. Using another equation by Flory,24 the sequence length distribution of random copolymers can be calculated: wn = n(1 − PAA )2 PAA n − 1

Figure 16. Schematic representation of a PA11 hydrogen-bonded sheet in an α crystal. The red segments are PA12 units that might interfere in PA11/PA12 copolymers. In A the PA12 unit is positioned in a disturbing way, whereas in B and C the PA12 segments might be tolerated. Arrows suggest chain folding. The blue area highlights a sheet fragment in which all chains are arranged in a parallel fashion. The blue spheres in the insets represent nitrogen atoms. Oxygen atoms are the largest spheres in the sketch, followed by carbon and hydrogen. The sketch is derived from the PA12 structures presented in ref 45.

area corresponds to a sheet fragment in which all chains run parallel. For the hydrogen bonds to fit, a progressive shift of the chains is required.34 One can derive that the PA12 carbonyl group would collide with that of the neighboring chain if a PA11 unit would be replaced by a PA12 unit at the top side of this sheet and if the planar extended chain conformation would be preserved. This effect is highlighted in the A square of Figure 16 and most likely cannot be tolerated in the crystal. This problem does not exist at the other side of the sheet, where because of the progressive shearthe PA12 carbonyl group would surpass that of the neighboring chain as illustrated in the B square. Such a minor defect would most likely be tolerated by which 3 unit long PA11 sequences could also contribute to the 4 unit thick crystals, as illustrated for the chain that runs upward from the B square in Figure 16. The PA12 segment on the other (top) side would simply end up in the amorphous phase. Sequences of only two PA11 units are too short as the problem as illustrated in the A square would arise. Another option, compatible with the inclusion of PA12 segments at the PA11 crystal borders, is illustrated in the C square to Figure 16.

(3)

where wn is the normalized weight fraction of the sequence of n units. At a comomomer mole fraction PAA = 0.5, the fraction of sequences larger than or equal to 4 units equals 31%, which seems large enough to bring a reasonably high crystallinity in PA11/PA12 copolymers. To obtain 60 Å thick polyethylene crystals, one needs ethylene sequences that are 24 units long, of which there is a negligible amount in such a random copolymer. However, a peculiarity is noted when combining the WAXD based PA11 and PA12 crystallinity shares with the DSC based overall crystallinity estimates. In Figure 14 the PA11 and PA12 shares to the total crystallinity at selected copolymer L


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to amost likely intolerablecollision of the carbonyl groups. This intolerance for PA11 segments might be inductive to crystallization at larger supercooling and may explain why the copolymer crystallinity on the PA12 rich side declines more rapidly compared to that at the PA11 rich side in Figure 14. Again, conversion into the mesomorphic γ′ phase may increase the acceptance for PA11 segments. Even in the ideal case where the crystals would be perfectly PA11 or PA12 pure and of the same size and perfection, one would observe a melting point spread for the copolymers since the melt composition around the crystals is continuously changing when melting progresses.30 In the present case this natural spread in the melting point most likely is convoluted with crystal quality aspects. Recall that MRR is observed in all cases (Figure 7). Often MRR is associated with the melting of thinner crystals and the recrystallization into thicker ones that melt at higher temperatures. Given that 4 unit cell thick crystals were persistently found, it may well be that the crystal lateral size and degree of comonomer inclusion rather than the crystal thickness cause crystal quality differences and hence differences in propensity to recrystallize. These aspects need to be addressed in future research. However, it seems that the ultimate crystal quality at the end of the melting traces is comparable for the copolymers since eq 2 can be used to describe the copolymer composition dependent end melting point evolution (and even rather ideally at the PA11 rich side of the diagram in Figure 12). Although increasing the heating rate (partially) prevents MRR, it does not alter the end melting point (Figure 7), suggesting that the most perfect crystals were already created during the preceding cooling run. Note that the melting point depressions in random copolymers of PA6 with PA6.6 are stronger compared to what can be expected from eq 2, the component specific melting points and enthalpies. A similar eutectic phase diagram was obtained by Allen55 with an apparent eutecticum close to 60 wt % PA6 and a melting point of 150 °C (compared to 260 and 220 °C for pure PA66 and PA6). It seems that morphological effects more strongly contribute to the melting point depression in these copolymers. The crystallinity of this 60 wt % PA6 copolymer was 40 wt %, compared to 55−60% for the pure components.56 The 40/60 PA11/PA12 copolymer has the lowest melting point with the end of melting at 151 °C. It has a crystallinity of 16% according to the DSC analysis. To judge its suitability as component in blends with e.g. (bio)polymers, its mechanical properties need to be evaluated. The ultimate strength, ultimate strain, and modulus of the copolymer are all a bit lower compared to the values of the pure extruded homopolymers as shown in Figure 18. This decrease is most likely due to the reduced crystallinity. All in all, this set of properties might be sufficient for a number of applications, in particular those in which blends are considered with thermally unstable (biobased) components. Copolymers comprising more than two different momomers may lead to stronger melting point depression, but also to excessive crystallinity reductions.1 Although the melting temperature of 60/40 PA6/PA6.6 random copolymers may also allow for blending at low temperature, PA11/PA12 copolymers have the extra advantage of being (partially) biobased.

In that case, a 2 unit long PA12 segment makes up a chain fold. Folding results in an antiparallel chain in which the carbonyl and NH group are directionally inverted with respect to the neighboring chain. In odd polyamides this does not hamper the continuation of hydrogen bond formation.34 This is also illustrated at the right side of the blue area in Figure 16 where a fold is made in a pure PA11 chain. In summary, the minimum sequence length for PA11 to crystallize seems to be 3 rather than 4. This brings the amount of sufficiently long sequences for PA11 in the 50/50 copolymer from 34 to 53%. Given that PA11 is present for 50% in this sample, one can calculate that roughly 60% of the sufficiently long sequences need to be in the crystalline state to account for the 16.2% (absolute) contribution to the total crystallinity as depicted in Figure 14. This is still quite a lot. However, the WAXD based PA11 share in the crystallinity is inflated as the PA12 sequences at the PA11 crystal borders contribute to the PA11 crystallinity. For the PA11/PA12 30/70 copolymer (which corresponds to 31.5 mol % PA11), 22.5% of the PA11 sequences are at least 3 units long, implying that PA11 can at most contribute 7.3% (absolute) tot the total crystallinity. The PA11 share in the total crystallinity is 3.7% (absolute; see Figure 14), implying that approximately 50% of the crystallizable PA11 sequences are actually crystalline, which with the arguments given aboveseems reasonable. Finally, the disturbing impact of incorporating PA12 units might be further alleviated by allowing departures from the sheetlike arrangement, i.e., by converting into the δ′ phase, as actually is observed. Incorporating PA11 sequences in PA12 is less obvious as illustrated in Figure 17 where a PA12 sheet is depicted, assuming a γ phase. In any scenario, the introduction of a PA11 segment while the chain conformation is preserved would lead

Figure 17. Schematic representation of a PA12 hydrogen-bonded sheet in a γ crystal. The red segments are PA11 units, which are all positioned in a disturbing way. The blue spheres in the inset represent nitrogen atoms. Oxygen atoms are the largest spheres in the sketch, followed by carbon and hydrogen. The sketch is derived from the PA12 structures presented in ref 45.

4. CONCLUSIONS To transfer the superior PA properties to thermally unstable, biobased or natural polymers via melt blending, the PA melting M


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low melting point and since the crystallinity and mechanical performance typical of the parent homopolymers are preserved, random PA11/PA12 copolymers are suited for blending with thermally unstable (biobased) substances.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (B.G.). Funding

The authors acknowledge the Agency for Innovation through Science and Technology Flanders (IWT Vlaanderen), the Strategic Initiative Materials in Flanders (SIM), and the KU Leuven Industrial research fund for financial support. B. Goderis thanks FWO-Vlaanderen for supporting the ESRF/ DUBBLE Big Science project (G.0C12.13). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors thank the staff at the European Synchrotron Radiation Facility and DUBBLE for assistance during the measurements.

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ABBREVIATIONS DSC, differential scanning calorimetry; HFIP, 1,1,1,3,3,3hexafluoro-2-propanol; MRR, melting−recrystallization−remelting; NMR, nuclear magnetic resonance; PA, polyamide; POM, polarized optical microscopy; RID, refractive index detector; SEC, size exclusion chromatography; TGA, thermogravimetrical analysis.

Figure 18. Modulus (A), stress at break (B), and strain at break (C) of PA11, PA12, and the transamidated PA11/PA12 40/60 random copolymer.

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REFERENCES

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temperature has to be decreased to avoid degradation during blending. Random copolymers from polyamide 11 (PA11) and polyamide 12 (PA12) have suitable melting temperatures and were obtained from the parent homopolymers by transamidation reactions during high temperature reactive extrusion. The melting temperatures of these products compare well to those of the random copolymers synthesized by Johnson and Mathias.12 This agreement is taken as evidence for full randomization by reactive extrusion. Extruding at 350 °C for 30 min was found to be a good compromise in which the transamidation reaction speed was reasonably high and thermal degradation minimal. Postcondensation reactions happened concomitantly, leading to increased polymer molecular weights. In physical blends as well as random copolymers from PA11 and PA12, both species essentially crystallize separately. The copolymer melting point depression can be understood in terms of the Flory theory for random copolymer melting. With both species being crystallizable, eutectic phase behavior was observed. Close to the eutectic composition, crystallization happens in a competitive mode and induces mesomorphic phase formation. To understand the copolymer crystallization behavior, large amounts of rigid amorphous material need to segregate from the melt together with the crystalline material. Although the crystal centers in PA11/PA12 random copolymers remain PA11 or PA12 pure, it seems that foreign comonomer units are tolerated at the crystal borders, at least for PA11 based crystals. At a 40/60 PA11/PA12 composition, a (lowest) copolymer melting temperature of 151 °C was obtained. Because of their N


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Random Copolymers from Polyamide 11 and Polyamide 12 by

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