Influence of the Main-Chain Configuration on the Mechanical

Article pubs.acs.org/Macromolecules

Influence of the Main-Chain Configuration on the Mechanical Properties of Linear Aliphatic Polyesters Mark P. F. Pepels,*,† Leon E. Govaert,‡ and Rob Duchateau†,§ †

Laboratory of Polymer Materials, Department of Chemical Engineering and Chemistry, Eindhoven University of Technology, P.O. Box 513, 5600 MB, Eindhoven, The Netherlands ‡ Polymer Technology, Department of Mechanical Engineering Eindhoven University of Technology, P.O. Box 513, 5600 MB, Eindhoven, The Netherlands § SABIC T&I, STC-Geleen, SABIC Europe B.V., Urmonderbaan 22, 6160 AH Geleen, The Netherlands ABSTRACT: Linear aliphatic polyesters have been investigated extensively over the last decades. However, the relation between the configuration of the ester groups in the main chain and the mechanical properties is only poorly understood. Therefore, in this work the influence of the composition of these polymers on the morphology, mechanical properties and relaxation processes for a set of random copolymers of εcaprolactone (CL) and ω-pentadecalactone (PDL) is explored. For these isomorphic copolymers, the crystallinity and lamellar thickness was shown to be independent of the composition over the largest part of the composition range. However, the yield stress does decrease significantly when comonomers were introduced in either of the homopolymer chains. Dynamic mechanical analysis revealed an additional high-temperature relaxation process for the copolymers, which was attributed to mobility of the crystalline phase (αc-mobility) and is responsible for the decrease in yield stress relative to the homopolymers. The origin of this mobility was related to the stacking of the ester groups in the crystal lamellae, which occurs less regular in copolymers and therefore gives rise to lower energy barriers for defect propagation (responsible for αc-mobility). Furthermore, the yield-kinetics of polypentadecalactone and two copolymers were accurately captured using the Ree−Eyring theory, showing the relation between αc-mobility and the contribution of both interlamellar and intralamellar shear to the yield stress.

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INTRODUCTION Aliphatic long-chain polyesters (ALCPEs) are interesting alternatives for polyethylenes in high-end applications, since their structure can be tuned to include additional functionalities, such as degradability, which is difficult to achieve for polyethylene.1−3 The “PE-like” character of this class of polymers originates from their long linear methylene backbone dominating the crystallization behavior and crystal structure, as well as the resulting mechanical properties.4−6 An example is polypentadecalactone, which crystallizes in an orthorhombic unit cell with dimensions very similar as for PE.7 The ester groups fit well in the orthorhombic unit cell, leading to the inclusion of these groups with only a minor associated expansion of the unit cell dimensions compared to PE.7 The most profound effect of the inclusion of ester groups is an enthalpic penalty reducing the heat of fusion of the crystal, which simultaneously decreases the melting temperature.8,9 Nevertheless, the crystallinity remains high (>50%) independent of the amount of ester groups in the backbone of the polymer. Even though there are many reports on the synthesis and thermal properties of ALCPEs,10−15 a systematic investigation into the influence of the amount and configuration of the ester © XXXX American Chemical Society

groups on the mechanical properties has to our knowledge not been performed. Copolymers of ω-pentadecalactone (PDL) and ε-caprolactone (CL) are ideal candidates for such an investigation, since the random copolymers are easily synthesized in large quantities, while the amount of ester groups is readily determined by the ratio of the two monomers copolymerized.16 Furthermore, poly(CL-ran-PDL) copolymers were reported to exhibit isomorphism over the whole composition range, resulting in a steady decrease of the melting temperature upon increased CL content.17,18 In this work, we investigated the influence of the main chain configuration in poly(CL-ran-PDL) copolymers on the crystalline morphology and its effect on the mechanical properties, in particular the yield stress. Furthermore, the origin of this behavior has been studied and related to the molecular structure. Finally, the kinetics of the yield phenomena of the copolymers were investigated. Received: May 20, 2015 Revised: August 4, 2015

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DOI: 10.1021/acs.macromol.5b01089 Macromolecules XXXX, XXX, XXX−XXX

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Figure 1. Structure of the random copolymer and calculation of the mole fraction of ester groups (Xester), in which m and n indicate the mol % of CL and PDL respectively.

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when high conversions of PDL are obtained.19 The random distribution of both monomer units in the obtained copolymers was furthermore confirmed by comparison of the theoretical and measured diads in 13C NMR. The relatively large variety in the obtained Mw-values is likely the result of trace amounts of impurities present in the monomers and in the applied solvent, capable of initiating the ROP. In addition, different chemical compositions of the copolymers also result in different hydrodynamic volumes during SEC measurement, which makes a reliable comparison of Mw values difficult if not impossible. Nevertheless, the obtained molecular weights are all high enough to yield ductile deformation behavior (see below). Before the mechanical behavior was investigated, the thermal and crystalline properties of the copolymers were examined. The DSC thermograms of the first cooling run from the melt and the second heating run are shown in Figure 2a. Similar to the work by Scandola and co-workers,17 also for this composition range, the peak melting- and crystallization temperatures decrease upon increasing the CL content in the copolymers (Figure 2b). This decreasing trend can be regarded as an extension of the data obtained on the ALCPEs investigated in a previous study, ranging from polyethylene to polypentadecalactone.9 Additionally, a broadening of the melting and crystallization peaks is observed for the copolymers relative to the homopolymers. Likely, this is the result of the existence of chain-segments rich in one of the two counits, as can be expected from a statistical copolymer. The melting enthalpy (ΔHm) also shows a similar decreasing trend with increasing CL content, either indicating a decrease in crystallinity, and/or a decrease in the heat of fusion of the

RESULTS AND DISCUSSION Thermal and Crystalline Properties of Poly(CL-ranPDL) Copolymers. Although the isomorphic nature of random copolymers of CL and PDL has been reported previously,17 the chemical composition range of the investigated copolymers was rather limited. Therefore, in this work a series of poly(CL-ran-PDL) copolymers with a broad composition range were synthesized via the ring-opening polymerization (ROP) of CL and PDL using aluminum salen catalysts (Figure 1, Table 1). It has been shown that for this Table 1. Composition and Molecular Weight of the Poly(CL-ran-PDL) Copolymers P(CL-PDL)

XCLa

XPDLa

Xesterb

Mw (g·mol−1)c

ĐMc

P(100−0) P(95−5) P(90−10) P(80−20) P(60−40) P(40−60) P(20−80) P(0−100)

1.00 0.95 0.90 0.80 0.60 0.40 0.20 0.00

0.00 0.05 0.10 0.20 0.40 0.60 0.80 1.00

0.167 0.155 0.145 0.128 0.104 0.088 0.076 0.067

62 189 209 90 147 167 180 101

2.2 3.2 2.9 3.3 2.5 2.4 2.2 2.8

a

Mole fraction of CL/PDL. bMole fraction of ester groups calculated using 1/(XCL·6+XPDL·15). cDetermined using SEC at 160 °C in TCB relative to polyethylene standards.

combination of monomers and catalyst the rate constant of transesterification is similar as the rate constant of polymerization of PDL, which results in completely random copolymers

Figure 2. DSC thermograms of the poly(CL-ran-PDL) copolymers showing the first cooling- and second heating run (a). Melting enthalpy, peak melting- and crystallization temperature as a function of the ester group content (b). Closed symbols are from the work described in this chapter, open symbols are taken from ref 9. B


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Figure 3. Long period, lamellar thickness, from SAXS and TEM, and the crystallinity obtained from WAXD as a function of the mole fraction of CL units in the copolymer (a). TEM images of the samples of interest for the crystalline morphology transformation (b).

To visualize this effect, transmission electron microscopy (TEM) imaging was carried out revealing the lamellar morphologies shown in Figure 3b. The similarity in morphology can clearly be observed for P(0−100) and P(80−20), as well as for the copolymers in between (results not shown here). Also the average lamellar thicknesses obtained from the measurement of 20 lamellae visible on the corresponding TEM picture, remains constant over this composition range. P(100−0) clearly shows thinner lamellae and a different morphology compared to P(0−100). Comparison of the transition region from 80−100 mol % CL reveals the change of a PPDL-type of morphology for P(80−20), to an intermediate morphology having thinner lamellae for P(90− 10), and further to a PCL-type of morphology for P(95−5). The lower melting enthalpies, melting- and crystallization temperatures, crystallinities and lamellar thicknesses, in combination with the changing lamellar morphology observed by TEM suggests that in the region between P(80−20) and (P100−0), the crystal morphology changes from a PDLdominated structure, to a CL-dominated structure. It appears that up to 80 mol % (90 wt %) of the CL units can be incorporated into the PPDL structure without changing its crystalline morphology. Influence of the Copolymer Composition on the (Dynamic) Mechanical Properties. The deformation behavior of the poly(CL-ran-PDL) copolymers was investigated by means of uniaxial tensile testing on die-cut dumbbell samples obtained from compression molding. All the (co)polymers exhibited ductile deformation and high elongations at break, indicating that all the molecular weights are sufficiently high for comparison of the yield stresses between the (co)polymers (Figure 4a). The tensile curves clearly show a decrease in yield stress upon inclusion of CL into the PPDL backbone from 17.5 MPa for neat PPDL, down to 9.2 MPa for P(60−40) (Figure 4b). Upon increasing the concentration of CL units to 80 mol % and higher, the yield stress increases again up to 14.6 MPa for PCL. The yield stress is generally considered as the onset of plastic flow of the crystals, for which the lamellar thickness is the main determining factor in the case of homopolymers.24

crystals due to the inclusion of ester groups. However, both P(90−10) (Xester = 0.145) and P(95−5) (Xester = 0.155) seem to deviate from this trend, having a lower ΔHm compared to P(80−20) and P(100−0). Even though these CL-rich copolymers have less ester groups than PCL, both their peak melting and crystallization enthalpies are lower, showing the deviation of these samples from the trend observed for the largest part of the composition range. The higher melting/ crystallization temperature of PCL cannot be ascribed to the general trend that at very high ester concentration the melting point increases, since this is typically observed at a methylene to ester group ratio of three, i.e., poly(ethylene adipate)20 or poly(γ-hydroxy butyrate).21 Most likely, this increase is caused by the favorable dipole−dipole interaction resulting from the increase in regularity for the PCL homopolymer. Next to the main crystallization peak observed in the DSC cooling curves, also an additional smaller peak is observed at lower temperatures for all copolymer samples up to 80% CL content. Similar as has been shown for LLDPE, this peak likely originates from a secondary crystallization event related to bundle-like ordering processes, which follow the primary crystallization and lamellar insertion (secondary crystallization) in the spherulites.22,23 In order to elucidate the trends observed for the thermal properties, both wide- and small-angle X-ray diffraction/ scattering (WAXD/SAXS) were performed to gain additional insight into the influence of the ester groups on the crystal unit cell, crystallinity and lamellar thickness. All copolymers containing up to 80 mol % CL revealed crystallinities between 51% − 55%, similar to P(0−100), while P(90−10) and P(95− 5) exhibit lower crystallinities of 43% and 46%, respectively (Figure 3a). Interestingly, also the lamellar thicknesses (lc), calculated from the long period obtained from the Lorentzcorrected SAXS patterns, shows a similar trend. Up to 80 mol % CL, the crystal thickness of the copolymers is similar to that of PPDL (lc ≈ 10 nm), indicating that the PDL sequences in the copolymers dominate this parameter. A further increase of the CL content above 80 mol % leads to a rapid decrease of the lamellar thickness toward the value of PCL (7 nm). C


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Figure 4. Stress vs strain curves of the poly(CL-ran-PDL) copolymers (a). Yield stress dependence on the composition of the copolymers (b). Experiments were performed at 20 °C using a strain rate of 1.33 × 10−2 s−1.

Figure 5. Storage modulus (solid lines) and loss modulus (dashed lines) vs the temperature (a). Curves are vertically shifted for better visualization. tan δ vs the temperature (b). Samples were measured using a heating rate of 3 °C·min−1 at a frequency of 1 Hz.

The relaxation corresponding to the Tg can be seen at −59 °C for P(100−0), gradually shifting to higher temperatures for the copolymers up to −28 °C for P(0−100), which is a typical feature of random copolymers. To avoid confusing in a later stage, this relaxation will be referred to as the alpha relaxation of the amorphous phase, αa, while the alpha relaxation related to the crystalline phase, as for instance is observed in polyethylene, will be referred to as αc. Next to these γ-, β-, and αa-transitions, for some copolymers an extra transition is observed in the high temperature regime, manifesting itself as a change of slope for the storage modulus and the observation of a corresponding peak for tan δ. This transition is observed (peak tan δ) for P(20−80) at 62 °C, for P(40−60) at 47 °C, and for P(60−40) at 42 °C, which is in a similar temperature range as the αc-relaxation of polyethylene.27 The αc-relaxation is generally attributed to an increase in mobility of the chains within the lamellae, leading to a steeper decrease of the modulus as a function of temperature.28

However, as shown in Figure 3, the lamellar thicknesses remain constant within the regime exhibiting the decrease in yield stress, while the increase in yield stress for copolymers having 90−100 mol % CL goes hand in hand with a decrease in lamellar thickness. Therefore, the origin of this effect is likely related to some other parameter, overruling the effect of the lamellar thickness. In order to get more insight into the mechanical relaxation processes involved, dynamic mechanical thermal analysis (DMTA) was performed on the poly(CL-ran-PDL) copolymer samples obtained from compression molding (Figure 5). All the polymers exhibit the typical γ-transition around −130 °C related to a local-mode secondary relaxation of a small amount of methylene groups,25 also observed for polyethylene.26 Furthermore, a β-transition around −90 °C can be observed which is related to relaxation processes involving water, commonly observed for polymers containing polar groups and therefore is more pronounced in the CL-rich copolymers.25 D


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Figure 6. tan δ vs temperature for DMTA measurements of P(20−80) at frequencies ranging from 0.01 Hz−10 Hz (a). Log frequency, ln(υ), corresponding to the peak of tan δ as a function of 1/T for P(20−80). Activation energy (Ea,αc) calculated using the slope (ln(υ)/(1/T))·R), in which R is the gas constant.

Figure 7. Complex modulus (G*) and tan δ vs temperature for the quenched and the slowly cooled samples of P(20−80) (a). Corresponding tensile tests (strain rate =1.33 × 10−2 s−1) showing the yield regime (b).

of the slower cooling rate, it is expected that both the lamellar thickness and the crystallinity increase with respect to the quenched sample. The difference in mechanical behavior of quenched and slowly cooled samples can be observed in the corresponding DMTA measurements (Figure 7a). First of all, the overall modulus for the slowly cooled sample is higher than for the quenched sample as a result of the increased crystallinity. Moreover, the αc-relaxation shifts to higher temperatures and becomes obscured by the occurrence of melting of the sample. The increase to higher temperatures of the αc-relaxation when samples are slowly cooled is a general phenomenon for polymers having this relaxation, which results from the increase of the lamellar thickness.28 Both the increase in lamellar thickness and the associated shift of the αc-relaxation lead to a significant increase in the yield stress of P(20−80) (Figure 7b). Discussion on the (Dynamic) Mechanical Properties. From the results described above, it can clearly be seen that the decrease in yield stress presented for the copolymers in Figure

To confirm that the observed extra relaxation for the copolymers containing 20−60 mol % CL is not related to a thermodynamic transition, DMTA-temperature sweeps were performed on the P(20−80) sample at frequencies ranging from 0.01−10 Hz (Figure 6a). Indeed a clear shift in temperature of the peak of tan δ can be observed from 46 °C at 0.01 Hz to 64 °C at 10 Hz, confirming the kinetic nature of this transition, which is therefore likely the αc-relaxation. Furthermore, a typical Arrhenius relation is observed between the frequency and the peak temperature, allowing the derivation of the activation energy of this αc-relaxation (Ea,αc) resulting in a value of 340 kJ·mol−1 (Figure 6b). The measurements described above were performed on samples which were cooled from the melt (quenched) between cold plates, resulting in a cooling rate of ±1 °C·s−1 in the crystallization regime. In order to investigate the effect of cooling rate, P(20−80) was compression molded at 150 °C after which the heating was turned off, allowing slow cooling (SC) of the sample at a cooling rate of ±0.01 °C·s−1. Because E


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Figure 8. Schematic representation of the translational movement of a chain ( × = methylene unit, red ● = ester group) through the crystal and the associated energy landscape for regular (a) and irregular ester stacking (b).

to a more regular system, since CL becomes the most abundant counit. This will likely lead to an increase of the αc-relaxation temperature. However, because simultaneously the melting temperature decreases further, the polymer melts before the αcrelaxation can be observed. This hypothesis also rationalizes the observed increase of the yield stress upon increasing the CL content from 60−100 mol %. The fact that the αc-relaxation is observed at a higher temperature when P(80−20) is slowly cooled can be explained by an increase in lamellar thickness.28 However, it is also likely that part of the shift is due to a more regular stacking of chainsegments with similar ester group distributions within the crystal. Indeed, the periodic stacking of ester groups along the c-axis is relatively sensitive for the cooling conditions.35 Therefore, upon slower cooling, the energy landscape will have an increased amount of high barriers, leading to a shift of the αc-relaxation to higher temperatures. The observed activation energy (340 kJ·mol−1) for the αcrelaxation of P(80−20) is much higher than that of polyethylene, which is typically 120−150 kJ·mol−1.28,36 Possibly, the inclusion of ester groups in the crystal lattice has a similar effect as the inclusion of methyl branches. Indeed, for PE with 12.5 methyl branches/1000 methylene units an increase of activation energy of the αc-relaxation was observed from 147 to 227 kJ·mol−1, which was ascribed to the hindrance of chain-to-chain slippage.36 To summarize, even though ester groups do not appear to have a significant effect on the crystalline structure, irregular stacking of the ester groups seems to increase the crystal mobility, thereby lowering the yield stress of the associated polymer. Yield-Kinetics of Poly(CL-ran-PDL) Copolymers. The mechanical behavior of the poly(CL-ran-PDL) copolymers is further elucidated by investigation of the yield-kinetics using tensile tests at various strain rates and temperatures. First, the PPDL homopolymer (P(0−100)) was studied in order to explain the deformation phenomena. As can be seen in Figure 9a, the yield stress increases upon increased strain rate, both at 20 and 45 °C. Furthermore, two distinct slopes can be observed in these regimes, which indicates that there are two different molecular deformation processes contributing to the yield stress.37−40 The slope at low strain rates corresponds to intralamellar deformation (process I), which is crystal slip within the crystals and lamellar fragmentation (Figure 9b).41,42 The second slope is the result of the sum of process I and process II, of which the latter was proposed to be linked to interlamellar shear and corresponds to the movement of different lamellar crystals with respect to each other, effectively deforming the amorphous phase while keeping the lamellar crystals intact.42 In process II the contribution to the yield

4 is the direct result of an increased mobility relative to the homopolymers, which is likely related to the crystalline αcrelaxation. For polyethylene, the origin of the αc-relaxation is generally attributed to twist defects propagating in the stems (chains) within the crystal lamellae.26 This allows chains to propagate in and out of the lamellar planes, leading to an increased mobility of the normally restricted intercrystalline (amorphous) region. The possibility of the twist defects to propagate is highly dependent on the interactions within the crystal. For example, polyamides, which have planar zigzag conformations similar to PE, do not exhibit an α c relaxation.29,30 This is generally thought to be attributed to the strong interactions (due to hydrogen bonding) between adjacent stems, giving rise to a high energy barrier for the formation and propagation of mobile localized structures (defects).26,31 On the other hand, for linear aliphatic polyesters the interaction between ester groups is weaker than that of amide groups, leading to a less straightforward situation. X,Y-type polyesters (PEX,Y), where X and Y are the number of carbons in the diol and diacid, respectively, seem to exhibit an αcrelaxation depending on the size of X and Y.32 For example, it was shown that PE6,8 and PE10,16 polyesters have a clear loss peak around 40 °C corresponding to the crystalline phase, while no such transition was observed for PE2,4, PE2,6, PE2,8, PE6,4, and PE6,6.32,33 Furthermore, it was suggested that the relaxation temperature was lowered by a decrease of the methylene sequence length.26 As stated before, there is no report among the limited data available for PPDL on the occurrence of an αc-relaxation,25 while PCL, which has been investigated more extensively, was also not reported to display such a relaxation.34 However, our results show that poly(CL-ran-PDL) copolymers having 20−60 mol % CL do exhibit an αc-relaxation. The origin of this effect is likely to be related to the regularity of the ester groups in the polymer chain. PPDL crystallizes in an orthorhombic unit cell in which the ester groups are periodically located together along the c-axis.7 Apparently, the favorable dipole interaction between the ester groups in PPDL result in an energy barrier which is too high for mobile segments (defect) to propagate over, resulting in immobile crystals (Figure 8a). However, upon the introduction of CL, the spacing between the ester groups loses its regularity, resulting in less good ester stacking between stems in the crystal. Consequently, even though there is an increased amount of ester groups present, their less regular stacking results in more but smaller energy barriers (Figure 8b). Therefore, the mobility of the defects in the crystal remains high, resulting in an increased mobility relative to PPDL for copolymers containing 20−60 mol % CL and the observation of an αc-relaxation. A further increase of the CL content leads F


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propagation of screw dislocations within the lamellar planes. This proceeds via nucleation and propagation of dislocations and/or defects. The change in mobility of these defects typically follows an Arrhenius relation. However, also the number of defects present in the crystals becomes larger with increasing temperature,45 leading to an increased mobility of the polymer chains. Therefore, as suggested by van Erp and coworkers, the high value for ΔUI could be related to the collective effect of the thermal activation of defect mobility and an increase of the number of defects.38 Next to the PPDL homopolymer, also P(20−80) and P(40− 60) were investigated with respect to the yield-kinetics (Figure 10). It was shown for isotactic polypropylene (i-PP) that the incorporation of comonomer (e.g., ethylene) does not change the values for Vi and ΔUi of both processes, and only requires the alteration of ε̇0,i. Even though for P(20−80) and P(40−60) the CL units are included in the crystalin contradiction to ethylene units, which are excluded from the i-PP crystalsthis still offers a good starting point. Therefore, for these copolymers eq 1 was applied using Vi and ΔUi of PPDL for both processes and ε̇0,i was varied to obtain satisfactory fits. For P(20−80), this resulted in ε̇0,I = 2 × 10135 s−1 and ε̇0,II = 3 × 1025 s−1, while for P(40−60) this leads to ε̇0,I = 8 × 10137 s−1 and ε̇0,II = 2 × 1026 s−1. It should be noted that for P(20−80) the measurements were performed at 24 °C instead of 20 °C (P(40−60) and PPDL), due to the difference in ambient temperature at the moment of measurement. For the copolymers the onset of process II is observed to shift to higher strain rates compared to PPDL. This is the direct result of the αc-relaxation observed for these copolymers, which increases the mobility of the interlamellar regions, eliminating its contribution to the yield stress. Also between the two copolymers the difference in αc-relaxation is observed, since ε̇0,II for P(40−60) (Tα = 62 °C) is an order of magnitude higher than that of P(20−80) (Tα = 47 °C). Furthermore, the contribution to the yield stress of process I (intralamellar deformation) in the low strain rate regime is significantly lower for the copolymers compared to PPDL, as can be seen from the large difference in ε̇0,I. As stated above, yielding in this regime is related to slippage of the crystal planes and is dependent on the mobility of thermally activated defects

stress originates from stems within different lamellar crystals connected through the amorphous phase by tie-chains that are pulled through the crystal lattice. This contribution becomes negligible at low strain rates/high temperatures, since the αcrelaxation of the crystals allows free movement of the stems in and out of the crystal within the time frame of the experiment. The two processes can be described by the Ree−Eyring modification of the Eyring theory,38,43,44 which assumes that both processes act independently and that their stress contributions can be added:

∑

σ(ε ,̇ T ) =

i = I , II

⎡ ε̇ ⎛ ΔU ⎞⎤ kT sinh−1⎢ exp⎜ i ⎟⎥ ⎝ RT ⎠⎥⎦ ⎢⎣ ε0,̇ i Vi

(1)

in which σ is the yield stress, ε̇ is the applied strain rate, k is Boltzmann’s constant, T is the temperature, R is the gas constant, Vi is the activation volume, ΔUi is the activation energy, and ε̇0,i is the pre-exponential (fit) factor of each of the two processes i. Application of this theory to the experimental yield stresses for PPDL in Figure 9 leads to a satisfactory fit of the yieldkinetics, of which the parameters for both processes are given in Table 2. The activation volume of process I determines the Table 2. Ree−Eyring Parameters for PPDL material

i

Vi (nm3)

ΔUi (kJ·mol−1)

ε̇0,i (s−1)

PPDL

I II

19.9 7.98

906 158

8.25 × 10126 2.13 × 1023

slope in the low strain rate regime, while the activation volume of both process combined governs the slope at high rates. From this, the activation volume of process II is easily derived. It can be seen that at room temperature both process I and II are governing the yield behavior over the largest part of the investigated strain rates, and very slow deformation is required to eliminate the contribution of process II. The fact that the activation energy of process I is much larger than the bond energy of a C−C bond (284−368 kJ·mol−1) seems peculiar, but can be explained by considering the crystal slip associated with this deformation. Yielding is generally considered as the

Figure 9. Strain rate and temperature dependence of the yield stress for PPDL (a). Solid line represents the best fit of eq 1. Dotted lines represent the contribution of process I and II, respectively. Schematic representation of intralamellar shear (process I) and interlamellar shear (process II) (b). G


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Figure 10. Strain rate dependence of the yield stress for P(20−80) at 24 °C (a) and P(40−60) at 20 °C (b). Solid line represents the theoretical yield stress from eq 1 using the Vi and ΔHi of PPDL, while fitting ε̇0,i. Dotted lines represent the contribution of processes I and II, respectively.

exponential factors, ε̇0,I/II. The differences in value for ε̇0,II between the polymers relates to the increased amount of αcmobility of the crystals for PPDL < P(20−80) < P(40−60), and is reflected in the higher required strain rates before the contribution of interlamellar shear is observed. Furthermore, the αc-mobility of the crystals leads to a decrease in yield stress for process I, since there is a parallel between the conformational defects capable of activating screw dislocations and the ones that have been suggested to account for the αc-relaxation. Therefore, the configuration of the ester groups within aliphatic long chain polyesters is a key factor influencing the crystal relaxation processes, and the associated properties such as yield stress and modulus.

and the number of defects. It was shown by Séguéla and coworkers that there is a parallel between the conformational defects capable of activating screw dislocations and the ones that have been suggested to account for the αc-relaxation.46 Indeed, this would explain the large difference for the yield stresses of PPDL (no observable αc-relaxation) and the P(20− 80)/P(40−60) (observable αc-relaxation) in the low strain rate regime.

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CONCLUSIONS Poly(CL-ran-PDL) copolymers with varying amounts of the two monomer units remain highly crystalline over the whole composition range and exhibit a steady decrease in melting temperature with increasing ester group content from PPDL toward that of PCL. This decrease is caused by the inclusion of ester groups in the crystal lattice leading to a decrease of the heat of fusion of the crystals. The crystalline morphology is dominated by the PDL units and remains constant up to 80 mol % of CL in the polymer backbone, after which a further increase of CL leads to a transformation to a CL-dominated morphology. Despite the fact that up to 80 mol % CL the crystallinity and lamellar thickness remained constant, a clear decrease of the yield stress of the copolymers was observed in this regime. Conversely, from 80−100 mol % CL the yield stress increases again to higher values, while simultaneously the lamellar thickness decreases. The origin of this effect was identified by DMTA as an increase in the mobility within the lamellar crystals of the copolymers, as a result of less regular stacking of the ester groups. This leads to lower energy barriers for the propagation of defects in stems within the lamellae, giving αcmobility to the crystals, effectively reducing the resistance to deformation. The yield-kinetics of PPDL could readily be described by the Ree−Eyring equation, taking into account two distinct processes, i.e., intralamellar shear at low strain rates/high temperatures (process I) and an additional process II, corresponding to interlamellar shear at high strain rates/low temperatures. The yield-kinetics of P(20−80) and P(40−60) could be described by application of the determined activation volume and activation energy of PPDL, and by fitting the pre-

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EXPERIMENTAL SECTION

Reagents and methods. All solvents and reagents were purchased from commercial sources (Sigma-Aldrich, BioSolve) unless stated otherwise. p-Xylene (99.9%) was dried over sodium and fractionally distilled under nitrogen and degassed prior to use. PDL and CL were distilled from CaH2 under nitrogen prior to use. The aluminum Schiff base complex (1), [N,N′-bis(salicylidene-1,2-ethylenediimine) aluminum ethyl (1), was synthesized using literature procedure.47 Irganox 1010 and Irgafos 168 were kindly received from BASF. All reactions and preparations were either carried out in an MBraun MB-150 GI glovebox or using proper Schlenk techniques. 1 H NMR and 13C NMR spectra were recorded in 5 mm tubes on a Varian Mercury 400 MHz spectrometer equipped with an autosampler at ambient probe temperature in CDCl3. Chemical shifts are reported in ppm vs tetramethylsilane. A Polymer Laboratories PL XT-220 robotic sample handling system was used as autosampler. High Temperature size exclusion chromatography (HT-SEC) of the copolymers was performed at 160 °C using a Polymer Laboratories PLXT-20 Rapid GPC Polymer Analysis System (refractive index detector and viscosity detector) with 3 PLgel Olexis (7.5 × 300 mm, Polymer Laboratories) columns in series. 1,2,4-Trichlorobenzene was used as eluent at a flow rate of 1 mL·min−1. The reported molecular weights were calculated with respect to polyethylene. The polyethylene calibration curve was calculated from polystyrene standards (Polymer Laboratories, Mp = 950 Da up to Mp = 2.25 × 106 Da) using log(MPE) = 1.0066 × log(MPS) − 0.3588. This conversion has been determined on our system and is used due to the limited availability of polyethylene standards. H


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Macromolecules Differential scanning calorimetry (DSC) analyses of the polymers was carried out on a DSC Q1000 from TA Instruments at a heating and cooling rate of 10 °C·min−1 in the desired temperature range. The reported thermograms, crystallization and melting temperatures correspond to (the peak temperatures of) the first cooling run and second heating run, respectively. Samples were compression molded into plates with dimensions of 40 × 40 × 1 mm by subsequently melting the sample at 150−170 °C and pressing them at 100 bar for 10 min, after which the samples were cooled using water-cooled plates. These samples were used for X-ray analysis and TEM analysis, as well as for tensile testing and DMTA. Wide- and small-angle X-ray diffraction/scattering (WAXD/SAXS) measurements were performed on a Ganesha laboratory instrument equipped with a GeniX-Cu ultra low divergence source producing Xray photons with a wavelength of 1.54 Å and a flux of 1 × 108 ph·s−1. Scattering patterns were collected using a Pilatus 300 K silicon pixel detector with 487 × 619 pixels of 172 μm2 in size placed at a sampleto-detector distance of 480 and 1080 mm, respectively. The beam center and the q range were calibrated using the diffraction peaks of silver behenate. The mass fraction of the crystal phase, calculated using the peak integrals of the crystalline phase divided by the total area in the WAXD spectra, is used as the crystallinity. Since the exact lamellar thickness requires knowledge of the volume percentage of crystals and since the density of the amorphous and crystalline phase of the copolymers are unknown, the mass fraction of crystal (derived from WAXD) × the long period (derived from SAXS) was used to calculate the average lamellar thickness. Therefore, the calculated values are affected by an absolute error, and they must only be considered for comparison between the different samples. TEM analysis was done using the following procedure. Before analysis, the samples (obtained using compression molding, vide inf ra) were trimmed at low temperature (−130 °C) and subsequently stained for 20 h. with a RuO4-solution prepared according to Montezinos et al.48 Ultrathin sections (70 nm) were obtained at −100 °C using a Leica Ultracut S/FCS microtome. The sections were placed on a 200 mesh copper grid with a carbon support layer. The sections were examined in a Tecnai 20 transmission electron microscope, operated at 200 kV. Tensile tests were performed with a Zwick Z100 and a Zwick Z010 tensile tester equipped with a 100 N and 1 kN load cell, respectively. A temperature chamber was used for the experiments performed at elevated temperatures. Dog-bone shaped tensile bars were die-cut from the compression molded plates. A grip-to-grip separation of 20 mm was used with a gauge length (LE) of 12.5 mm. The samples were prestressed to 0.5 N, then loaded with a constant cross-head speed ranging from 10−5 to 1 s−1. In order to complement the yield data in the low strain rate regime, creep experiments were performed by rapidly applying various loads and following the strain over time. Since the relation creep rate−applied stress is directly interchangeable with the relation strain rate−yield stress, the combined experiments were used for the investigation of the yield-kinetics.49 Dynamic mechanical thermal analysis (DMTA) was performed on a TA Instruments Q800 in film tension mode at frequencies ranging from 0.01−10 Hz and a strain of 0.1%. A maximum temperature range of −150 °C up to +100 °C was used at heating rates of 0.1 (0.01 Hz measurement of P(80−20)), 1 (P(80−20) measurements) or 3 °C· min−1 (standard measurement of all polymers). Polymerization Procedure. In all reactions similar amounts of total monomer and catalyst were used. As an example the synthesis of P(60−40) is described. Al(Salen)Et (100 mg, 0.31 mmol), benzylalcohol (16.8 mg, 0.165 mmol), PDL (26.58 g, 110.6 mmol), CL (8.42 g, 73.73 mmol), and p-xylene (50 g) were added to a 100 mL glass crimp cap vial in a nitrogen-filled glovebox. The vial was capped, taken out of the glovebox and placed in an oil bath at 100 °C, while stirring with a magnetic stirring bar for 12 h to allow full conversion of the monomers. After the reaction, the reaction mixture was removed from the vial and the crude sample was analyzed using 1H NMR spectroscopy and SEC. The reaction mixture was grinded, washed 3 times with methanol, and dried under vacuum. After this, the powder was transferred into a solution of Irgafos 168 and Irganox 1010 (0.5 wt

% each relative to the polymer) in acetone and stirred. The solution was allowed to evaporate under ambient conditions after which the powder was dried in a vacuum oven at 40 °C.

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AUTHOR INFORMATION

Corresponding Author

*(M.P.F.P.) E-mail: [email protected]. Notes

The authors declare no competing financial interests

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ACKNOWLEDGMENTS Financial support by SABIC for this work is gratefully acknowledged. The authors thank Ilja Voets (Eindhoven University of Technology) for her help with X-ray analysis and Anne Spoelstra (Eindhoven University of Technology) for her help with TEM analysis. Marc Kanters (Eindhoven University of Technology) is kindly acknowledged the discussions and help regarding the Ree−Eyring fitting protocol.

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REFERENCES

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