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PDL–PEG–PDL blends: Effects of blend morphology on film toughness
European Polymer Journal 69 (2015) 308–318
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Stereocomplexation of PLL/PDL–PEG–PDL blends: Effects of blend morphology on film toughness Sunan Tacha a, Tinnakorn Saelee a, Wootichai Khotasen a, Winita Punyodom a,b, Robert Molloy a,b, Patnarin Worajittiphon a,b, Puttinan Meepowpan a, Kiattikhun Manokruang a,b,⇑ a b
Department of Chemistry, Faculty of Science, Chiang Mai University, Chiang Mai 50200, Thailand Materials Science Research Center, Faculty of Science, Chiang Mai University, Chiang Mai 50200, Thailand
a r t i c l e
i n f o
Article history: Received 25 April 2015 Received in revised form 8 June 2015 Accepted 9 June 2015 Available online 10 June 2015 Keywords: Stereocomplex Poly(lactide) Toughness Atomic force microscopy (AFM)
a b s t r a c t Polylactide (PL) finds its use in a wide variety of applications ranging from medical devices to engineering plastics. However, the use of PL is limited to certain applications due to its brittle characteristic and low heat distortion temperature, which indicate its poor mechanical and thermal properties, respectively. This present work demonstrated the toughening of PL by stereocomplexation adopting poly(D-lactide)–poly(ethylene glycol)– poly(D-lactide) (PDL–PEG–PDL) copolymers with various PDL segment length. The DSC results showed that the complete stereocomplexation was reached when 40 wt% copolymer was blended with poly(L-lactide) (PLL). In addition, the crystallization of PEG was interrupted and, thus, prohibited by either adding PDL or PLL in the system. AFM images showed that, for the first time, the stereocomplex crystallites (Sc-crystallites) formed by the enantiomer pairs were dispersed in a continuous amorphous phase of PL and the non-crystallizing PEG. The increase of PDL segment length in the copolymer leaded to the increase of the Sc-crystallite size, which, consequently, resulted in the increase of the tensile strength of the blended-films. Elongation at break of the films was found to rely on the determined amorphous area. To maximize the toughness of the films, the specific ratio of PDL and PEG segment length in the copolymer must be achieved to provide the optimum balance between Sc-crystallite size and % amorphous area. Ó 2015 Elsevier Ltd. All rights reserved.
1. Introduction Polylactide (PL), one of the well-known aliphatic polyesters, has received much interest due to its enzymatic-/hydroly tic-degradability as well as biocompatibility [1–3]. PL finds its use in a wide variety of applications ranging from medical devices to engineering plastics [4–7]. Since its monomer can be derived from renewable agricultural sources, this polymer is considered as a ‘‘green’’ alternative thermoplastic which can potentially replace a number of petrochemical based polymers in the next era. However, the use of PL is limited to certain applications due to its brittle characteristic and low heat distortion temperature, which indicate its poor mechanical and thermal properties, respectively. Many approaches adopted to improve heat distortion temperature of PL such as plasticization and nucleation [8,9], and blending PLL with its
⇑ Corresponding author at: Department of Chemistry, Faculty of Science, Chiang Mai University, Chiang Mai 50200, Thailand. E-mail address: [email protected] (K. Manokruang). http://dx.doi.org/10.1016/j.eurpolymj.2015.06.015 0014-3057/Ó 2015 Elsevier Ltd. All rights reserved.
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enantiomeric forms, i.e. poly(D-lactide) (PDL) and poly(L-lactide) (PLL) [10–13], mostly leaded to stereocomplexation by which the two enantiomeric forms of PL are simply mixed to form stereocomplexed crystallites (Sc-crystallites). Sc-crystallites, originated from stereoselective van der Waals forces, act as molecular-interlocking arrangements in the semicrystalline region. These strongly packed crystallites provide a considerable increase in the melting point, typically 50–70 °C higher than either PDL or PLL alone [14–17]. Although stereocomplexation offers a choice for increasing heat distortion temperature, the brittleness is still of concern because the increase in crystallinity generally results in low ductility [18]. So to simultaneously increase thermal stability and toughness of PL, the copolymers of its enantiomeric segments are adopted for stereocomplexation. Among the successful literatures reported [7,18–26], the copolymers of poly(ethylene glycol) (PEG) and PDL, when blending with PLL, were shown to provide good balance between crystallinity and toughness due to the plasticizing effect of the former and the specific affinity toward stereocomplexation of the latter [7,18,23,24]. It is well established that morphological features, e.g. the degree of crystallinity, the crystalline size, the chain conformation distribution of the amorphous phase and the crystallization rate, are responsible for the mechanical properties of the semicrystalline polymers, in general [23,24,27–29]. Specifically, the mechanisms for toughening of Sc-PL involving PEG were proposed to originate from a combination of various factors including lowering Tg of the amorphous phase system by PEG, the presence of PL continuous amorphous phase, the presence of a soft PEG amorphous phase within the PL amorphous region and the appearance of the a0 phase, i.e. the crystalline structure formed at the temperatures below 120 °C, of PL [23]. These morphological features have been widely investigated using several techniques such as differential scanning calorimetry (DSC) [15,18], polarized optical microscopy (POM) [7,23,24], wide-angle X-ray diffraction (WAXD) [15] and infrared (IR) and Raman spectroscopy [7,23,24]. Since the aforementioned Sc-PL prepared by blending PDL and PLL is well known to improve thermal properties but the brittleness of Sc-PL still remains a problem. The present work demonstrated the toughening of PL by stereocomplexation in nonequimolar mixtures of PLL and PDL–PEG–PDL copolymers. In addition to the literatures reported, this work, for the first time, showed that the Sc-crystallites, formed by the enantiomer pairs and, then, dispersed in a continuous amorphous phase of PL and the non-crystallizing PEG can be viewed using atomic force microscopy (AFM) technique. The morphology-toughness relationship of stereocomplexed-polylactide (Sc-PL) was elucidated. Also supported by DSC results, the crystallization of PEG was interrupted by the presence of PL and, instead, was incorporated into a continuous amorphous phase. AFM analysis also showed that the combination of the size of Sc-crystallites and the determined amorphous regime is attributable to the toughness of Sc-PL. 2. Experiments 2.1. Materials Commercial grade PLL (PLL 2003D, Mn = 1.6 105 g mol1 and PDI = 2.0) was supplied by Nature Works. D-lactic acid was purchased from Musashino Chemical Laboratory. PEG (Mn = 8000 and 20,000) and stannous octoate, Sn(Oct)2, were purchased from Sigma–Aldrich. All other reagents were analytical grade and used as received. 2.2. Synthesis of PDL–PEG–PDL copolymers The PDL–PEG–PDL copolymers were synthesized by ring-opening polymerization. The predetermined amounts of D-lactide
(DL) and PEG, in the presence of 1 wt% Sn(Oct)2, were added to the reaction flask equipped with a magnetic stirrer bar at the same time under vacuum condition. Intrinsically, PDL–PEG–PDL copolymers were synthesized using PEG molecular weight 8000 and applied different DL/EG molar ratio to yield different PDL segment lengths. The polymerization was carried out at 120 °C for 24 h. After polymerization, the products were dissolved in chloroform and precipitated in methanol. The copolymers were dried under vacuum before further use. 2.3. Preparation of PLL/copolymer blended-films
PLL and the copolymer were solution-blended at a concentration of 5 g/dl in chloroform for the predetermined PLL/copolymer compositions (100/0, 90/10, 80/20, 70/30, 60/40 and 50/50 wt%). The solutions were cast onto glass petri-dishes and left for solvent evaporation at room temperature for approximately 3 days in a vacuum oven. 2.4. Characterizations The synthesized copolymers were subjected to structural characterized by 1H NMR spectroscopy (Bruker 400 NMR spectrometer, Switzerland). The spectra were obtained from polymer solutions in deuterated chloroform (CDCl3) at room temperature with TMS as an internal standard. Average molecular weights and polydispersity indices of the copolymers were determined by gel permeation chromatography (GPC, Waters 2485 size-exclusion chromatograph) operating at 35 °C using tetrahydrofuran (THF) as a solvent.
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Thermal stability and decomposition profile of the samples were determined by thermogravimetric analysis (TGA, Perkin–Elmer). Each of the sample was heated over the temperature range of 50–500 °C with heating rate of 20 °C/min under a nitrogen atmosphere. Thermal transitions and percentage crystallinity of the blended-films were followed by DSC (DSC7, Perkin–Elmer). A test sample with a typical weight of 8–10 mg was encapsulated in a sealed aluminum pan. The sample was heated over a temperature range of 0–250 °C at heating rate of 10 °C/min. Sc-crystallites were determined according to the literature [30] as followed. The crystallinity of the homocrystallites (Xh) was calculated as:
Xh % ¼
DHh 100% p 93 J=g
and the crystallinity of the Sc-crystallites (Xsc) was calculated as:
Xsc % ¼
DHsc 100% p 142 J=g
where DHh and DHsc are the heat of fusion for homocrystallites and Sc-crystallites, respectively, p was the PL percent (wt%, including PLL and PDL in PLL/PDL–PEG–PDL) in the prepared film. The 93 J/g and 142 J/g were reported heat of fusion values with 100% crystallinity for homocrystallites and Sc-crystallites, respectively. According to ASTM D882, the samples were cut into 70 mm 10 mm 0.5 mm. Tensile testing was performed on at least three specimens operating with the universal testing machine at room temperature using a sample grip with a gauge length of 50 mm and the load cell of 0.1 N (preload = 0.1 N). The tensile speed of cross head was 50 mm/min. Each AFM sample was prepared by depositing one drop of a 0.1% w/v the polymer solution (in chloroform) on the glass substrate (1 1 cm). The solvent was left to evaporate at room temperature. Measurements were done in tapping mode by a digital Instrument’s Nanoscope IIIa MultiMode SPM atomic force microscope (the tip radius = 5–10 nm, spring constant = 20–100 N/m, drive frequency = 200–300 kHz, cantilever length = 125 lm and set-point ratio, i.e. oscillation amplitude/free air oscillation amplitude = 37.56 mV). From the AFM images, the average diameter of Sc-crystallites, the areas occupied by the Sc-crystallites and amorphous regime were determined using Image J software.
3. Results and discussions 3.1. Structural characterization Triblock copolymers, PDLm–PEGn–PDLm with various PDL segment lengths, were synthesized by ring-opening polymerization. PEG, as a macromonomer, was copolymerized with DL in the presence of 1 wt% Sn(Oct)2. The reaction was carried out at 120 °C for 24 h. Fig. 1(a) showed a schematic representation of the copolymer synthesis. The final polymer products were subjected to structural characterization employing 1H NMR Spectrometry. The spectra of PDLm–PEGn–PDLm copolymers were recorded at 400 MHz in deuterated chloroform at room temperature. Tetramethylsilane (TMS) was used as the internal standard at d = 0.00 ppm. Shown in Fig. 1(b), the signals from methine proton (–CH, d) and methyl proton (–CH3, a) were evidenced at d = 5.2 ppm and at 1.6 ppm, respectively, representing PDL segments in the copolymers. The characteristic signal of the ethylene oxide protons, (–OCH2CH2, b) appeared at d = 3.7 ppm. The NMR spectra clearly showed the progressive increase of methyl- and methine proton intensity when DL was more polymerized (Fig. 1(b)). The PDL composition of the copolymers, determined from the 1H NMR integral peaks, was found to be in the range of 53– 213. The inset spectrum in Fig. 1(b) showed the multiplicity caused by the methine proton (c) at the copolymer chain ends and the last methylene proton (f) unit on PEG, which confirmed that the reaction between PEG and DL was successfully achieved [31]. Molecular weights of the synthesized polymers were determined by GPC using THF as a solvent. Table 1 summarizes the synthesized copolymers and their molecular weights. 3.2. Thermal characterizations 3.2.1. Thermal stability of the copolymers Thermogravimetric analysis of PLL and PDL–PEG–PDL copolymers were performed at a heating rate of 20 °C min1 under nitrogen atmosphere. The thermal stability profiles of PLL and the copolymers are shown in Fig. 2. PLL homopolymer illustrated a one-state degradation process under nitrogen atmosphere. The initial decomposition (weight loss) temperature for PLL started at 239 °C until the degradation process was completed at 390 °C. The copolymers, on the other hands, showed a two-state degradation process as the degradation temperatures (Td) were approximately evidenced at 255 °C and 380 °C. The former corresponds to the degradation temperature at which PDL block was degraded and the latter corresponds to the temperature at which the degradation of PEG was taken place. For all the copolymers, the two degradation temperatures were almost comparable and independent of the composition. The total weight loss of the copolymers was prolonged by the addition of PEG in the main chain.
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Fig. 1. (a) Schematically represents the copolymer synthesis. The 1H NMR spectra (b) show the progressive increase of DL in the copolymers. The inset spectrum confirms that the reaction between PEG and DL was successfully achieved.
Table 1 The synthesized copolymers and their molecular weights. Polymer
DPPEG a
PLL2003D PDL53–PEG182–PDL53 PDL81–PEG182–PDL81 PDL126–PEG182–PDL126 PDL213–PEG182–PDL213 a b c
– 182 182 182 182
a
DPPDL – 53 81 126 213
b
Mn
b
– 1.6 104 1.9 104 2.6 104 3.9 104
Mn
PDIc
c 5
3.5 10 1.4 104 1.8 104 2.2 104 2.7 104
Provided by Sigma–Aldrich. Determined from 1H NMR. Determined from GPC.
Fig. 2. Thermogravimetric analysis profiles of the neat PLL and the synthesized copolymers.
1.94 1.53 1.47 1.47 1.53
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3.2.2. Thermal properties of the Sc-PL films The DSC analyses of PLL/copolymer blended-films were performed at a heating rate of 10 °C/min under nitrogen atmosphere. Thermal properties of the blended-films are summarized in Table 2. The thermograms of the films prepared from blending PLL with PDL53–PEG182–PDL53, PDL81–PEG182–PDL81, PDL126– PEG182–PDL126 and PDL213–PEG182–PDL213 are shown in Fig. 3(a)–(d), respectively. All melting temperatures (Tm) at about 150 °C corresponding to homocrystalline melting temperatures were clearly observed in both the neat PLL film and the blended films with added copolymer up to 30 wt%. Unlike PLL/PDL blended-films (see in supporting information), the addition of, at least, 40 wt% copolymer brought the complete stereocomplexation of the two PL enantiomers, i.e. PLL and PDL, in all blended-films as can be seen from a single melting peak observed at about 200 °C, corresponding to a characteristic melting temperature of Sc-crystallites [14,16]. The Sc-crystallinity (% Xsc) was, in addition, increased with increasing the copolymer composition in the blended-films. It should be noted that all the PLL/copolymer blended-films, with the addition of copolymers higher than 20 wt%, demonstrated the amount of (% Xsc) more than PLL/PDL blended-films (see in supporting information). These relatively high values of % Xsc were resulted from the plasticization effect of PEG in the copolymers that facilitated the mobility of the polymer chains and, thus, substantial alignment of the PLL and PDL segments [25,26]. The mobility of the polymer chains, however, was restricted by the increase of PDL segments in the copolymers. The longer PDL segments leaded to the progressive decrease of % Xsc of the films for the same blending compositions. Also, it should be noted that the melting peak of PEG block in the copolymer was decreased in its intensity when the PDL segment increased from 53 (in PDL53–PEG182–PDL53) to 81 (in PDL81–PEG182–PDL81), respectively shown in Fig. 3(a) to (b) bottom thermograms, and the PEG melting peak was completely disappeared when the PDL segment further increased, i.e. from 126 (in PDL126–PEG182–PDL126) to 213 (in PDL213–PEG182–PDL213) as respectively observed in Fig. 3(c) to (d) bottom thermograms. The effect was clearly exemplified in Fig. 3(e), showing the thermograms in second heating run where the thermal history of the copolymers was erased. The gradual decline and eventual disappearance of PEG melting temperature with increasing PDL segment length in the copolymers clearly suggested that the crystallization of PEG was restricted by the addition of PL. This effect was also emphasized when blending PLL with the copolymers. The thermograms of PLL/copolymer films mostly exhibited no PEG melting temperature. The crystallization of PEG, therefore, was interrupted and, thus, prohibited by either adding PDL or PLL in the system [25]. In addition, the melting temperature of PDL in the copolymers slightly shifted to higher temperature with increasing PDL segment length. 3.2.3. Morphology-mechanical properties relationship The tensile testing experiment of the neat PLL showed a typical feature of a glossy, rigid material with tensile stress = 52 MPa, shown in Fig. 4(a). As seen in Fig. 4(a), the tensile strength was initially decreased with the addition of 10 wt% copolymer in the films prepared by blending PLL with either PDL126–PEG182–PDL126 or PDL213–PEG182–PDL213 and, then, the tensile strength was increased and reached the maximum when 20 wt% of the respective copolymers were blended in the films. The film strength was decreased again when more copolymer were added (i.e. 30–50 wt%). For the films prepared by blending PLL and the other copolymers, tensile stress was decreased when increasing the copolymer composition in the films. In addition, Fig. 4(a) also showed that the films demonstrated the higher tensile stress values when the copolymer with longer PDL segment was blended. Fig. 4(b) shows the plot of elongation at break as a function of composition for all blended films. The neat PLL illustrated low elongation at break of about 5%, shown in Fig. 4(b). However, PLL tensile ductility was improved showing the increase of elongation at break by the addition of the copolymers. The addition of about 30 wt% copolymer leaded to the dramatic increase of elongation at break of the films prepared by blending PLL with PDL53–PEG182–PDL53, PDL126–PEG182–PDL126 and PDL213–PEG182–PDL213. The significant increase of elongation at break for PLL/PDL81–PEG182–PDL81 film, however, can be seen early at 20 wt% addition of the copolymer. The increase of tensile strain of the films reached the threshold at 102.0 ± 11.5%, 206.1 ± 14.7%, 155.2 ± 4.1% and 144.0 ± 16.9% when PLL was blended with PDL53–PEG182–PDL53, PDL81– PEG182–PDL81, PDL126–PEG182–PDL126 and PDL213–PEG182–PDL213, respectively, with 40 wt% loading of the copolymers. Except for PLL/PDL53–PEG182–PDL53 film that showed the minimum elongation at break, increasing PDL segment length in the copolymer resulted in the significant decrease of the elongation at break of the blended-films. Since complete stereocomplexation was reached when at least 40 wt% copolymer was blended with PLL, the 60/40 wt% blending of PLL/copolymer films were chosen to study the effect of PDL segment length on the morphology of blended-films. AFM phase images of the neat PLL, shown in Fig. 5(a), illustrated that phase morphology of PLL was likely composed of large fibril-aggregated structure tending to grow from the center of a spherulite. These large fibrils were attributed to the folding of high MW of PLL chains (MW 200,000). Its blown out picture in Fig. 5(b) shows that the neat PLL spherulite was composed of the long-stripes and planar lamellae aligned in the disorder arrangement [2,15]. Fig. 5(c) shows an AFM phase image of the pure PDL53–PEG182–PDL53, as a representative copolymer, in which 2 types of crystalline phase were observed. Appeared in bright circular shapes were assigned to PEG homocrystallites since the relatively soft PEG caused several degrees of the phase lag in the AFM image. The other dark and relatively small spherical shapes were assigned to the PDL homocrystallites due to a slight phase shift caused by the PDL hard phase. These assignments were in correlation with the DSC thermogram of pure PDL53–PEG182–PDL53 (Fig. 3(a), bottom thermogram), in which two separated melting peaks were shown for melting temperatures of PEG and PDL. AFM, for the first time, showed that the crystallization of PEG was interrupted by blending the copolymer with PLL as can be seen that all the bright circular shapes were disappeared in Fig. 5(d) in which only the crystalline phase of PL were visible. The blown out picture is to show that the crystallites were
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S. Tacha et al. / European Polymer Journal 69 (2015) 308–318 Table 2 Thermal properties of the blended-films. PLL (wt%)
100 90 80 70 60 50 0
Copolymer (wt%)
0 10 20 30 40 50 100
PDL53–PEG182–PDL53 Tm(h) (°C)
DHm(h) (J/g)
150.5 149.8 151.0 147.5 – – 163.8
26.2 16.4 14.3 6.1 – – 22.0
PDL81–PEG182–PDL81
DHm(sc) (J/g)
X(sc)
DHm(h) (J/g)
(%)
Tm(sc) (°C)
DHm(sc) (J/g)
X(sc)
(%)
Tm(h) (°C)
X(h)
(%)
Tm(sc) (°C)
28.2 18.6 17.1 7.7 – – 23.7
– 204.0 204.5 204.5 207.7 206.2 –
– 7.9 16.1 34.7 42.5 51.4 –
– 5.9 14.0 28.7 37.4 48.3 –
150.5 149.8 149.7 143.7 – – 164.7
26.2 18.0 12.0 4.5 – – 29.2
28.2 20.2 14.0 5.5 – – 31.4
– 205.2 206.7 204.8 204.5 206.3 –
– 8.7 19.1 33.1 41.2 50.3 –
– 6.4 14.6 26.5 34.5 44.3 –
X(h)
PDL126–PEG182–PDL126
100 90 80 70 60 50 0
0 10 20 30 40 50 100
(%)
PDL213–PEG182–PDL213
Tm(h) (°C)
DHm(h) (J/g)
X(h) (%)
Tm(sc) (°C)
DHm(sc) (J/g)
X(sc) (%)
Tm(h) (°C)
DHm(h) (J/g)
X(h) (%)
Tm(sc) (°C)
DHm(sc) (J/g)
X(sc) (%)
150.5 148.8 144.7 139.0 – – 168.2
26.2 17.7 9.3 3.5 – – 36.5
28.2 19.6 10.6 4.1 – – 39.2
– 207.5 209.0 209.2 209.2 208.3 –
– 8.9 22.4 30.4 40.0 46.8 –
– 6.5 16.8 23.5 32.0 38.8 –
150.5 150.3 149.3 145.5 – – 170.7
26.2 14.7 5.3 0.7 – – 38.5
28.2 16.1 5.9 0.8 – – 41.4
– 208.8 212.5 209.3 206.7 209.7 –
– 8.4 15.5 30.3 38.0 41.8 –
– 6.0 11.4 22.7 29.1 32.7 –
formed in small spherical shape. These crystallites were solely corresponded to Sc-crystallites of the PL enantiomers, i.e. PDL and PLL, forming the dispersed phase in the continuous phase of amorphous PL and the non-crystallizing PEG. This morphology was also proven by the appearance of single crystalline melting peak observed at the temperature at which Sc-crystalline melting temperature was identified on the DSC thermogram of such blended film (60/40 thermogram in Fig. 3(a)). This effect was emphasized by the increase of the size for Sc-crystallites when the copolymers with longer PDL segment were blended with PLL. Fig. 5(e)–(g) shows that the dark circular spots, in a circular connected pattern, in the AFM phase images were progressively increased in size with increasing PDL segment length from PDL81–PEG182–PDL81, PDL126–PEG182–PDL126 and PDL213–PEG182–PDL213, respectively, in the blended-films. The average diameter and their distribution of Sc-crystallites, analyzed by Image J software, ranged from 0.064 ± 0.02, 0.494 ± 1.02, 0.534 ± 1.31 and 0.690 ± 1.70 lm when PLL was blended with PDL53–PEG182–PDL53, PDL81–PEG182–PDL81, PDL126–PEG182–PDL126 and PDL213–PEG182–PDL213, respectively. The distinguished area occupied by Sc-crystallites was also analyzed by the software, which resulted in the remaining amorphous phase regime being subsequently determined, shown in Fig. 5(h). As it will be seen in further discussion, PLL/PDL81– PEG182–PDL81 film illustrated the largest amorphous region among the blended-films, which contributed to the dramatic increase of the film toughness. It is noted that in the conventional blend of PLL and PDL, Sc-crystallization proceeds much faster than the homo-crystallization in the composition at which both crystallizations are possible. The Sc-crystallites were initially formed and, then, the homo-crystallites were nucleated in the matrix of the preformed Sc-crystallites and grew until impinging to others. For the blending compositions where the nucleation of homo-crystallites were absence, only Sc-crystallites were nucleated and continued to grow, giving the size being larger than 20 lm [15]. The Sc-crystallization in PLL/PDL–PEG– PDL copolymer blends, on the other hand, was restricted within the block copolymer phase separation where Sc-crystallites can grow only to some extent until reaching the phase separated domain, resulting in relatively much smaller crystallites (0.064–0.699 lm). The changes in tensile strength and elongation at break in correlation with morphology were monitored as a function of increasing PDL segment length for the films prepared by blending PLL with 40 wt% copolymers. Shown in Fig. 4(c) demonstrates the changes of tensile stress in correlation with Sc-crystallite size while the change of elongation at break in correlation with % amorphous area are shown in Fig. 4(d). The changes in tensile strength and elongation at break were respectively governed by the Sc-crystallite size and % amorphous area. As can be seen in Fig. 4(c) and (d), the highest value of tensile strength was obtained in corresponding with the largest size of Sc-crystallites while the highest value of elongation at break was achieved in corresponding with the largest amorphous area. Fig. 4(e) shows the tensile stress–strain curves of the 60/40 wt% blended films of PLL/copolymers in comparisons with the neat PLL film. Toughness of the films, calculated from the area under the curve, was 2.3 104, 5.6 104, 4.8 104 and 4.5 104 kJ m3 when PLL was blended with PDL53–PEG182–PDL53, PDL81–PEG182–PDL81, PDL126–PEG182–PDL126 and PDL213–PEG182–PDL213, respectively. Among the blended films, the toughness of PLL was increased the most when blending with PDL81–PEG182–PDL81. This behavior may be attributed to the large amorphous area of this blended film, in corresponding to AFM image analysis in that PLL/PDL81–PEG182–PDL81 system demonstrated the largest amorphous area (i.e. 72%) so, as a consequence, stress induced-mobility of the polymer chains was well accommodated in the amorphous region. Also, the sufficient Sc-crystallite size leaded to the effective tensile strength. With these two morphological features, this blended-film underwent relatively large deformation before breaking down.
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Fig. 3. DSC thermograms of PLL/copolymer blended-films prepared from blending PLL with (a) PDL53–PEG182–PDL53, (b) PDL81–PEG182–PDL81, (c) PDL126– PEG182–PDL126 and (d) PDL213–PEG182–PDL213 in various compositions; (e) the second heating run thermograms of the copolymers.
It is noted that the order of increasing tensile strength was directly related to the increase of Sc-crystallite size. PLL/PDL213–PEG182–PDL213 demonstrated the highest values of tensile strength due to its formation of the largest size of Sc-crystallites (Fig. 4(a)). Elongation at break was also proven to rely on the % amorphous area. The increase of % amorphous area resulted in the increase in elongation at break of the films. Owing to its largest % amorphous area, PLL/PDL81–PEG182– PDL81 exhibited the maximum elongation a break (Fig. 4(b)). Although the AFM morphology of the blended-films were investigated only for 60/40 wt% PLL/copolymer films, the changes in stress and elongation at break behaved in a similar fashion at other blending compositions. The AFM study on other blending compositions is currently in progress to establish the composition dependence on the morphology of the films. We hypothesized that the ratio between PDL and PEG segment length in PDL81–PEG182–PDL81 offered the optimum balance between Sc-crystallite size and % amorphous area when it was blended with PLL and, hence, provided the maximum toughness among the prepared films. Such blended-film demonstrated the toughness of about 5.6 104 kJ m3. This value was approximately 400 times larger than the toughness of the neat PLL film. To test the hypothesis, PDL189–PEG455–PDL189 was also synthesized having the ratio between PDL and PEG segment length as similar as in PDL81–PEG182–PDL81. This is to
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Fig. 4. (a) The tensile strength and (b) elongation at break as a function of blending composition for the PLL/copolymer blended-films, (c) the change of tensile stress in correlation with Sc-crystallite size and (d) the change of elongation at break in correlation with % amorphous area plotted as a function of PDL segment length, and (e) the tensile stress–strain curves, reflecting the toughness, of the 60/40 wt% blended-films of PLL/copolymers in comparisons with the neat PLL film.
elucidate whether this specific ratio reflects the optimum balance between the size of Sc-crystallites and % amorphous area, which offered the enhanced toughness of PL film in a similar fashion to PLL/PDL81–PEG182–PDL81 blending system. The synthesis detail of PDL189–PEG455–PDL189 copolymer was shown in supporting information. Fig. 6(a) shows the DSC thermograms of the neat PDL189–PEG455–PDL189 and the 60/40 wt% blended-film of PLL/PDL189–PEG455–PDL189. Similar to PDL81–PEG182–PDL81, two separate melting peaks corresponding to melting temperatures of PEG and PDL segments in PDL189–PEG455–PDL189 were evidenced at 55 °C and 167 °C, respectively. These two melting peaks were disappeared when the copolymer was blended with PLL, which, then, exhibited only the melting temperature of Sc-PL at 206 °C. Stress–strain curves of the PLL/PDL189–PEG455–PDL189 films were illustrated in Fig. 6(b). PLL/PDL189–PEG455–PDL189 films exhibited the behavior of ductile materials with tensile strength, elongation at break and toughness being determined at 36 MPa, 195% and 5.8 104 kJ m3, respectively. AFM image in Fig. 6(c) shows that PLL/PDL189–PEG455–PDL189 film illustrated the spherical Sc-crystallites, as well as other blended films, with the average crystallite diameter of
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Fig. 5. AFM phase images of (a) PLL spherulite and (b) its blown out picture, (c) the neat PDL53–PEG182–PDL53, (d) PLL/PDL53–PEG182–PDL53, (e) PLL/PDL81– PEG182–PDL81, (f) PLL/PDL126–PEG182–PDL126 and (g) PLL/PDL213–PEG182–PDL213; (h) the average diameter and their distribution of Sc-crystallites including the distinguished area of Sc-crystallites and the amorphous phase.
0.699 ± 1.50 lm. When compared with PLL/PDL81–PEG182–PDL81, PLL/PDL189–PEG455–PDL189 exhibited relatively high tensile strength and low elongation at break due to its relatively large size of Sc-crystallites and less % amorphous area, respectively. The toughness of the film, however, was comparable to that of PLL/PDL81–PEG182–PDL81 on account of the similar ratio between PDL and PEG segment length. The results suggested that the specific ratio between PDL and PEG segment length in PDL81–PEG182–PDL81 and PDL189–PEG455–PDL189 copolymers provided the optimum balance between the size of Sc-crystallites and % amorphous area, when blending with PLL, which resulted in the largely-enhanced toughness of the blended films.
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Fig. 6. (a) The DSC thermograms of the neat PDL189–PEG455–PDL189 and the 60/40 wt% blended-film of PLL/PDL189–PEG455–PDL189, (b) stress–strain curves of the PLL/PDL189–PEG455–PDL189 indicating the toughness of the films and (c) AFM phase image of PLL/PDL189–PEG455–PDL189 film and the determined average crystallite diameter and % amorphous area.
It is also noted that the crystallite size of PLL/PDL189–PEG455–PDL189 film was comparable to that of PLL/PDL213–PEG182– PDL213 film. Therefore the tensile strength values of the two are indifferent. However, owing to its noticeably large amorphous regime, PLL/PDL189–PEG455–PDL189 film demonstrated the significantly higher elongation at break. 4. Conclusions The Sc-PL films were prepared by blending the PDL–PEG–PDL copolymers, varying PDL segment lengths, with the neat PLL. The DSC results showed that the complete sterocomplexation was reached at 60/40 wt% PLL/copolymer blending composition for almost blended films. In addition, the crystallization of PEG was interrupted and, thus, prohibited by either adding PDL or PLL in the system. AFM images showed that, for the first time, the Sc-crystallites were formed and, then, dispersed in the amorphous of PL and the non crystallizing PEG when the copolymers were blended with the neat PLL. The increase of PDL segment length in the copolymer leaded to the increase of the Sc-crystallite size, which, consequently, resulted in the increase of the tensile strength of the blended-films. Elongation at break of the films was found to rely on the determined amorphous area. To maximize the toughness of the films, the specific ratio of PDL and PEG segment length must be achieved to provide the optimum balance between Sc-crystallite size and % amorphous area. Acknowledgements We gratefully acknowledge the financial support from the PTT Research and Technology Institute (PTT-RTI). This work was also supported by the National Research University Project under Thailand’s Office of the Higher Education Commission. Appendix A. Supplementary material Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/j.eurpolymj.2015.06.015.
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Stereocomplexation of PLL/PDL–PEG–PDL blends: Effects of blend morphology on film toughness Sunan Tacha a, Tinnakorn Saelee a, Wootichai Khotasen a, Winita Punyodom a,b, Robert Molloy a,b, Patnarin Worajittiphon a,b, Puttinan Meepowpan a, Kiattikhun Manokruang a,b,⇑ a b
Department of Chemistry, Faculty of Science, Chiang Mai University, Chiang Mai 50200, Thailand Materials Science Research Center, Faculty of Science, Chiang Mai University, Chiang Mai 50200, Thailand
a r t i c l e
i n f o
Article history: Received 25 April 2015 Received in revised form 8 June 2015 Accepted 9 June 2015 Available online 10 June 2015 Keywords: Stereocomplex Poly(lactide) Toughness Atomic force microscopy (AFM)
a b s t r a c t Polylactide (PL) finds its use in a wide variety of applications ranging from medical devices to engineering plastics. However, the use of PL is limited to certain applications due to its brittle characteristic and low heat distortion temperature, which indicate its poor mechanical and thermal properties, respectively. This present work demonstrated the toughening of PL by stereocomplexation adopting poly(D-lactide)–poly(ethylene glycol)– poly(D-lactide) (PDL–PEG–PDL) copolymers with various PDL segment length. The DSC results showed that the complete stereocomplexation was reached when 40 wt% copolymer was blended with poly(L-lactide) (PLL). In addition, the crystallization of PEG was interrupted and, thus, prohibited by either adding PDL or PLL in the system. AFM images showed that, for the first time, the stereocomplex crystallites (Sc-crystallites) formed by the enantiomer pairs were dispersed in a continuous amorphous phase of PL and the non-crystallizing PEG. The increase of PDL segment length in the copolymer leaded to the increase of the Sc-crystallite size, which, consequently, resulted in the increase of the tensile strength of the blended-films. Elongation at break of the films was found to rely on the determined amorphous area. To maximize the toughness of the films, the specific ratio of PDL and PEG segment length in the copolymer must be achieved to provide the optimum balance between Sc-crystallite size and % amorphous area. Ó 2015 Elsevier Ltd. All rights reserved.
1. Introduction Polylactide (PL), one of the well-known aliphatic polyesters, has received much interest due to its enzymatic-/hydroly tic-degradability as well as biocompatibility [1–3]. PL finds its use in a wide variety of applications ranging from medical devices to engineering plastics [4–7]. Since its monomer can be derived from renewable agricultural sources, this polymer is considered as a ‘‘green’’ alternative thermoplastic which can potentially replace a number of petrochemical based polymers in the next era. However, the use of PL is limited to certain applications due to its brittle characteristic and low heat distortion temperature, which indicate its poor mechanical and thermal properties, respectively. Many approaches adopted to improve heat distortion temperature of PL such as plasticization and nucleation [8,9], and blending PLL with its
⇑ Corresponding author at: Department of Chemistry, Faculty of Science, Chiang Mai University, Chiang Mai 50200, Thailand. E-mail address: [email protected] (K. Manokruang). http://dx.doi.org/10.1016/j.eurpolymj.2015.06.015 0014-3057/Ó 2015 Elsevier Ltd. All rights reserved.
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enantiomeric forms, i.e. poly(D-lactide) (PDL) and poly(L-lactide) (PLL) [10–13], mostly leaded to stereocomplexation by which the two enantiomeric forms of PL are simply mixed to form stereocomplexed crystallites (Sc-crystallites). Sc-crystallites, originated from stereoselective van der Waals forces, act as molecular-interlocking arrangements in the semicrystalline region. These strongly packed crystallites provide a considerable increase in the melting point, typically 50–70 °C higher than either PDL or PLL alone [14–17]. Although stereocomplexation offers a choice for increasing heat distortion temperature, the brittleness is still of concern because the increase in crystallinity generally results in low ductility [18]. So to simultaneously increase thermal stability and toughness of PL, the copolymers of its enantiomeric segments are adopted for stereocomplexation. Among the successful literatures reported [7,18–26], the copolymers of poly(ethylene glycol) (PEG) and PDL, when blending with PLL, were shown to provide good balance between crystallinity and toughness due to the plasticizing effect of the former and the specific affinity toward stereocomplexation of the latter [7,18,23,24]. It is well established that morphological features, e.g. the degree of crystallinity, the crystalline size, the chain conformation distribution of the amorphous phase and the crystallization rate, are responsible for the mechanical properties of the semicrystalline polymers, in general [23,24,27–29]. Specifically, the mechanisms for toughening of Sc-PL involving PEG were proposed to originate from a combination of various factors including lowering Tg of the amorphous phase system by PEG, the presence of PL continuous amorphous phase, the presence of a soft PEG amorphous phase within the PL amorphous region and the appearance of the a0 phase, i.e. the crystalline structure formed at the temperatures below 120 °C, of PL [23]. These morphological features have been widely investigated using several techniques such as differential scanning calorimetry (DSC) [15,18], polarized optical microscopy (POM) [7,23,24], wide-angle X-ray diffraction (WAXD) [15] and infrared (IR) and Raman spectroscopy [7,23,24]. Since the aforementioned Sc-PL prepared by blending PDL and PLL is well known to improve thermal properties but the brittleness of Sc-PL still remains a problem. The present work demonstrated the toughening of PL by stereocomplexation in nonequimolar mixtures of PLL and PDL–PEG–PDL copolymers. In addition to the literatures reported, this work, for the first time, showed that the Sc-crystallites, formed by the enantiomer pairs and, then, dispersed in a continuous amorphous phase of PL and the non-crystallizing PEG can be viewed using atomic force microscopy (AFM) technique. The morphology-toughness relationship of stereocomplexed-polylactide (Sc-PL) was elucidated. Also supported by DSC results, the crystallization of PEG was interrupted by the presence of PL and, instead, was incorporated into a continuous amorphous phase. AFM analysis also showed that the combination of the size of Sc-crystallites and the determined amorphous regime is attributable to the toughness of Sc-PL. 2. Experiments 2.1. Materials Commercial grade PLL (PLL 2003D, Mn = 1.6 105 g mol1 and PDI = 2.0) was supplied by Nature Works. D-lactic acid was purchased from Musashino Chemical Laboratory. PEG (Mn = 8000 and 20,000) and stannous octoate, Sn(Oct)2, were purchased from Sigma–Aldrich. All other reagents were analytical grade and used as received. 2.2. Synthesis of PDL–PEG–PDL copolymers The PDL–PEG–PDL copolymers were synthesized by ring-opening polymerization. The predetermined amounts of D-lactide
(DL) and PEG, in the presence of 1 wt% Sn(Oct)2, were added to the reaction flask equipped with a magnetic stirrer bar at the same time under vacuum condition. Intrinsically, PDL–PEG–PDL copolymers were synthesized using PEG molecular weight 8000 and applied different DL/EG molar ratio to yield different PDL segment lengths. The polymerization was carried out at 120 °C for 24 h. After polymerization, the products were dissolved in chloroform and precipitated in methanol. The copolymers were dried under vacuum before further use. 2.3. Preparation of PLL/copolymer blended-films
PLL and the copolymer were solution-blended at a concentration of 5 g/dl in chloroform for the predetermined PLL/copolymer compositions (100/0, 90/10, 80/20, 70/30, 60/40 and 50/50 wt%). The solutions were cast onto glass petri-dishes and left for solvent evaporation at room temperature for approximately 3 days in a vacuum oven. 2.4. Characterizations The synthesized copolymers were subjected to structural characterized by 1H NMR spectroscopy (Bruker 400 NMR spectrometer, Switzerland). The spectra were obtained from polymer solutions in deuterated chloroform (CDCl3) at room temperature with TMS as an internal standard. Average molecular weights and polydispersity indices of the copolymers were determined by gel permeation chromatography (GPC, Waters 2485 size-exclusion chromatograph) operating at 35 °C using tetrahydrofuran (THF) as a solvent.
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Thermal stability and decomposition profile of the samples were determined by thermogravimetric analysis (TGA, Perkin–Elmer). Each of the sample was heated over the temperature range of 50–500 °C with heating rate of 20 °C/min under a nitrogen atmosphere. Thermal transitions and percentage crystallinity of the blended-films were followed by DSC (DSC7, Perkin–Elmer). A test sample with a typical weight of 8–10 mg was encapsulated in a sealed aluminum pan. The sample was heated over a temperature range of 0–250 °C at heating rate of 10 °C/min. Sc-crystallites were determined according to the literature [30] as followed. The crystallinity of the homocrystallites (Xh) was calculated as:
Xh % ¼
DHh 100% p 93 J=g
and the crystallinity of the Sc-crystallites (Xsc) was calculated as:
Xsc % ¼
DHsc 100% p 142 J=g
where DHh and DHsc are the heat of fusion for homocrystallites and Sc-crystallites, respectively, p was the PL percent (wt%, including PLL and PDL in PLL/PDL–PEG–PDL) in the prepared film. The 93 J/g and 142 J/g were reported heat of fusion values with 100% crystallinity for homocrystallites and Sc-crystallites, respectively. According to ASTM D882, the samples were cut into 70 mm 10 mm 0.5 mm. Tensile testing was performed on at least three specimens operating with the universal testing machine at room temperature using a sample grip with a gauge length of 50 mm and the load cell of 0.1 N (preload = 0.1 N). The tensile speed of cross head was 50 mm/min. Each AFM sample was prepared by depositing one drop of a 0.1% w/v the polymer solution (in chloroform) on the glass substrate (1 1 cm). The solvent was left to evaporate at room temperature. Measurements were done in tapping mode by a digital Instrument’s Nanoscope IIIa MultiMode SPM atomic force microscope (the tip radius = 5–10 nm, spring constant = 20–100 N/m, drive frequency = 200–300 kHz, cantilever length = 125 lm and set-point ratio, i.e. oscillation amplitude/free air oscillation amplitude = 37.56 mV). From the AFM images, the average diameter of Sc-crystallites, the areas occupied by the Sc-crystallites and amorphous regime were determined using Image J software.
3. Results and discussions 3.1. Structural characterization Triblock copolymers, PDLm–PEGn–PDLm with various PDL segment lengths, were synthesized by ring-opening polymerization. PEG, as a macromonomer, was copolymerized with DL in the presence of 1 wt% Sn(Oct)2. The reaction was carried out at 120 °C for 24 h. Fig. 1(a) showed a schematic representation of the copolymer synthesis. The final polymer products were subjected to structural characterization employing 1H NMR Spectrometry. The spectra of PDLm–PEGn–PDLm copolymers were recorded at 400 MHz in deuterated chloroform at room temperature. Tetramethylsilane (TMS) was used as the internal standard at d = 0.00 ppm. Shown in Fig. 1(b), the signals from methine proton (–CH, d) and methyl proton (–CH3, a) were evidenced at d = 5.2 ppm and at 1.6 ppm, respectively, representing PDL segments in the copolymers. The characteristic signal of the ethylene oxide protons, (–OCH2CH2, b) appeared at d = 3.7 ppm. The NMR spectra clearly showed the progressive increase of methyl- and methine proton intensity when DL was more polymerized (Fig. 1(b)). The PDL composition of the copolymers, determined from the 1H NMR integral peaks, was found to be in the range of 53– 213. The inset spectrum in Fig. 1(b) showed the multiplicity caused by the methine proton (c) at the copolymer chain ends and the last methylene proton (f) unit on PEG, which confirmed that the reaction between PEG and DL was successfully achieved [31]. Molecular weights of the synthesized polymers were determined by GPC using THF as a solvent. Table 1 summarizes the synthesized copolymers and their molecular weights. 3.2. Thermal characterizations 3.2.1. Thermal stability of the copolymers Thermogravimetric analysis of PLL and PDL–PEG–PDL copolymers were performed at a heating rate of 20 °C min1 under nitrogen atmosphere. The thermal stability profiles of PLL and the copolymers are shown in Fig. 2. PLL homopolymer illustrated a one-state degradation process under nitrogen atmosphere. The initial decomposition (weight loss) temperature for PLL started at 239 °C until the degradation process was completed at 390 °C. The copolymers, on the other hands, showed a two-state degradation process as the degradation temperatures (Td) were approximately evidenced at 255 °C and 380 °C. The former corresponds to the degradation temperature at which PDL block was degraded and the latter corresponds to the temperature at which the degradation of PEG was taken place. For all the copolymers, the two degradation temperatures were almost comparable and independent of the composition. The total weight loss of the copolymers was prolonged by the addition of PEG in the main chain.
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Fig. 1. (a) Schematically represents the copolymer synthesis. The 1H NMR spectra (b) show the progressive increase of DL in the copolymers. The inset spectrum confirms that the reaction between PEG and DL was successfully achieved.
Table 1 The synthesized copolymers and their molecular weights. Polymer
DPPEG a
PLL2003D PDL53–PEG182–PDL53 PDL81–PEG182–PDL81 PDL126–PEG182–PDL126 PDL213–PEG182–PDL213 a b c
– 182 182 182 182
a
DPPDL – 53 81 126 213
b
Mn
b
– 1.6 104 1.9 104 2.6 104 3.9 104
Mn
PDIc
c 5
3.5 10 1.4 104 1.8 104 2.2 104 2.7 104
Provided by Sigma–Aldrich. Determined from 1H NMR. Determined from GPC.
Fig. 2. Thermogravimetric analysis profiles of the neat PLL and the synthesized copolymers.
1.94 1.53 1.47 1.47 1.53
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3.2.2. Thermal properties of the Sc-PL films The DSC analyses of PLL/copolymer blended-films were performed at a heating rate of 10 °C/min under nitrogen atmosphere. Thermal properties of the blended-films are summarized in Table 2. The thermograms of the films prepared from blending PLL with PDL53–PEG182–PDL53, PDL81–PEG182–PDL81, PDL126– PEG182–PDL126 and PDL213–PEG182–PDL213 are shown in Fig. 3(a)–(d), respectively. All melting temperatures (Tm) at about 150 °C corresponding to homocrystalline melting temperatures were clearly observed in both the neat PLL film and the blended films with added copolymer up to 30 wt%. Unlike PLL/PDL blended-films (see in supporting information), the addition of, at least, 40 wt% copolymer brought the complete stereocomplexation of the two PL enantiomers, i.e. PLL and PDL, in all blended-films as can be seen from a single melting peak observed at about 200 °C, corresponding to a characteristic melting temperature of Sc-crystallites [14,16]. The Sc-crystallinity (% Xsc) was, in addition, increased with increasing the copolymer composition in the blended-films. It should be noted that all the PLL/copolymer blended-films, with the addition of copolymers higher than 20 wt%, demonstrated the amount of (% Xsc) more than PLL/PDL blended-films (see in supporting information). These relatively high values of % Xsc were resulted from the plasticization effect of PEG in the copolymers that facilitated the mobility of the polymer chains and, thus, substantial alignment of the PLL and PDL segments [25,26]. The mobility of the polymer chains, however, was restricted by the increase of PDL segments in the copolymers. The longer PDL segments leaded to the progressive decrease of % Xsc of the films for the same blending compositions. Also, it should be noted that the melting peak of PEG block in the copolymer was decreased in its intensity when the PDL segment increased from 53 (in PDL53–PEG182–PDL53) to 81 (in PDL81–PEG182–PDL81), respectively shown in Fig. 3(a) to (b) bottom thermograms, and the PEG melting peak was completely disappeared when the PDL segment further increased, i.e. from 126 (in PDL126–PEG182–PDL126) to 213 (in PDL213–PEG182–PDL213) as respectively observed in Fig. 3(c) to (d) bottom thermograms. The effect was clearly exemplified in Fig. 3(e), showing the thermograms in second heating run where the thermal history of the copolymers was erased. The gradual decline and eventual disappearance of PEG melting temperature with increasing PDL segment length in the copolymers clearly suggested that the crystallization of PEG was restricted by the addition of PL. This effect was also emphasized when blending PLL with the copolymers. The thermograms of PLL/copolymer films mostly exhibited no PEG melting temperature. The crystallization of PEG, therefore, was interrupted and, thus, prohibited by either adding PDL or PLL in the system [25]. In addition, the melting temperature of PDL in the copolymers slightly shifted to higher temperature with increasing PDL segment length. 3.2.3. Morphology-mechanical properties relationship The tensile testing experiment of the neat PLL showed a typical feature of a glossy, rigid material with tensile stress = 52 MPa, shown in Fig. 4(a). As seen in Fig. 4(a), the tensile strength was initially decreased with the addition of 10 wt% copolymer in the films prepared by blending PLL with either PDL126–PEG182–PDL126 or PDL213–PEG182–PDL213 and, then, the tensile strength was increased and reached the maximum when 20 wt% of the respective copolymers were blended in the films. The film strength was decreased again when more copolymer were added (i.e. 30–50 wt%). For the films prepared by blending PLL and the other copolymers, tensile stress was decreased when increasing the copolymer composition in the films. In addition, Fig. 4(a) also showed that the films demonstrated the higher tensile stress values when the copolymer with longer PDL segment was blended. Fig. 4(b) shows the plot of elongation at break as a function of composition for all blended films. The neat PLL illustrated low elongation at break of about 5%, shown in Fig. 4(b). However, PLL tensile ductility was improved showing the increase of elongation at break by the addition of the copolymers. The addition of about 30 wt% copolymer leaded to the dramatic increase of elongation at break of the films prepared by blending PLL with PDL53–PEG182–PDL53, PDL126–PEG182–PDL126 and PDL213–PEG182–PDL213. The significant increase of elongation at break for PLL/PDL81–PEG182–PDL81 film, however, can be seen early at 20 wt% addition of the copolymer. The increase of tensile strain of the films reached the threshold at 102.0 ± 11.5%, 206.1 ± 14.7%, 155.2 ± 4.1% and 144.0 ± 16.9% when PLL was blended with PDL53–PEG182–PDL53, PDL81– PEG182–PDL81, PDL126–PEG182–PDL126 and PDL213–PEG182–PDL213, respectively, with 40 wt% loading of the copolymers. Except for PLL/PDL53–PEG182–PDL53 film that showed the minimum elongation at break, increasing PDL segment length in the copolymer resulted in the significant decrease of the elongation at break of the blended-films. Since complete stereocomplexation was reached when at least 40 wt% copolymer was blended with PLL, the 60/40 wt% blending of PLL/copolymer films were chosen to study the effect of PDL segment length on the morphology of blended-films. AFM phase images of the neat PLL, shown in Fig. 5(a), illustrated that phase morphology of PLL was likely composed of large fibril-aggregated structure tending to grow from the center of a spherulite. These large fibrils were attributed to the folding of high MW of PLL chains (MW 200,000). Its blown out picture in Fig. 5(b) shows that the neat PLL spherulite was composed of the long-stripes and planar lamellae aligned in the disorder arrangement [2,15]. Fig. 5(c) shows an AFM phase image of the pure PDL53–PEG182–PDL53, as a representative copolymer, in which 2 types of crystalline phase were observed. Appeared in bright circular shapes were assigned to PEG homocrystallites since the relatively soft PEG caused several degrees of the phase lag in the AFM image. The other dark and relatively small spherical shapes were assigned to the PDL homocrystallites due to a slight phase shift caused by the PDL hard phase. These assignments were in correlation with the DSC thermogram of pure PDL53–PEG182–PDL53 (Fig. 3(a), bottom thermogram), in which two separated melting peaks were shown for melting temperatures of PEG and PDL. AFM, for the first time, showed that the crystallization of PEG was interrupted by blending the copolymer with PLL as can be seen that all the bright circular shapes were disappeared in Fig. 5(d) in which only the crystalline phase of PL were visible. The blown out picture is to show that the crystallites were
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S. Tacha et al. / European Polymer Journal 69 (2015) 308–318 Table 2 Thermal properties of the blended-films. PLL (wt%)
100 90 80 70 60 50 0
Copolymer (wt%)
0 10 20 30 40 50 100
PDL53–PEG182–PDL53 Tm(h) (°C)
DHm(h) (J/g)
150.5 149.8 151.0 147.5 – – 163.8
26.2 16.4 14.3 6.1 – – 22.0
PDL81–PEG182–PDL81
DHm(sc) (J/g)
X(sc)
DHm(h) (J/g)
(%)
Tm(sc) (°C)
DHm(sc) (J/g)
X(sc)
(%)
Tm(h) (°C)
X(h)
(%)
Tm(sc) (°C)
28.2 18.6 17.1 7.7 – – 23.7
– 204.0 204.5 204.5 207.7 206.2 –
– 7.9 16.1 34.7 42.5 51.4 –
– 5.9 14.0 28.7 37.4 48.3 –
150.5 149.8 149.7 143.7 – – 164.7
26.2 18.0 12.0 4.5 – – 29.2
28.2 20.2 14.0 5.5 – – 31.4
– 205.2 206.7 204.8 204.5 206.3 –
– 8.7 19.1 33.1 41.2 50.3 –
– 6.4 14.6 26.5 34.5 44.3 –
X(h)
PDL126–PEG182–PDL126
100 90 80 70 60 50 0
0 10 20 30 40 50 100
(%)
PDL213–PEG182–PDL213
Tm(h) (°C)
DHm(h) (J/g)
X(h) (%)
Tm(sc) (°C)
DHm(sc) (J/g)
X(sc) (%)
Tm(h) (°C)
DHm(h) (J/g)
X(h) (%)
Tm(sc) (°C)
DHm(sc) (J/g)
X(sc) (%)
150.5 148.8 144.7 139.0 – – 168.2
26.2 17.7 9.3 3.5 – – 36.5
28.2 19.6 10.6 4.1 – – 39.2
– 207.5 209.0 209.2 209.2 208.3 –
– 8.9 22.4 30.4 40.0 46.8 –
– 6.5 16.8 23.5 32.0 38.8 –
150.5 150.3 149.3 145.5 – – 170.7
26.2 14.7 5.3 0.7 – – 38.5
28.2 16.1 5.9 0.8 – – 41.4
– 208.8 212.5 209.3 206.7 209.7 –
– 8.4 15.5 30.3 38.0 41.8 –
– 6.0 11.4 22.7 29.1 32.7 –
formed in small spherical shape. These crystallites were solely corresponded to Sc-crystallites of the PL enantiomers, i.e. PDL and PLL, forming the dispersed phase in the continuous phase of amorphous PL and the non-crystallizing PEG. This morphology was also proven by the appearance of single crystalline melting peak observed at the temperature at which Sc-crystalline melting temperature was identified on the DSC thermogram of such blended film (60/40 thermogram in Fig. 3(a)). This effect was emphasized by the increase of the size for Sc-crystallites when the copolymers with longer PDL segment were blended with PLL. Fig. 5(e)–(g) shows that the dark circular spots, in a circular connected pattern, in the AFM phase images were progressively increased in size with increasing PDL segment length from PDL81–PEG182–PDL81, PDL126–PEG182–PDL126 and PDL213–PEG182–PDL213, respectively, in the blended-films. The average diameter and their distribution of Sc-crystallites, analyzed by Image J software, ranged from 0.064 ± 0.02, 0.494 ± 1.02, 0.534 ± 1.31 and 0.690 ± 1.70 lm when PLL was blended with PDL53–PEG182–PDL53, PDL81–PEG182–PDL81, PDL126–PEG182–PDL126 and PDL213–PEG182–PDL213, respectively. The distinguished area occupied by Sc-crystallites was also analyzed by the software, which resulted in the remaining amorphous phase regime being subsequently determined, shown in Fig. 5(h). As it will be seen in further discussion, PLL/PDL81– PEG182–PDL81 film illustrated the largest amorphous region among the blended-films, which contributed to the dramatic increase of the film toughness. It is noted that in the conventional blend of PLL and PDL, Sc-crystallization proceeds much faster than the homo-crystallization in the composition at which both crystallizations are possible. The Sc-crystallites were initially formed and, then, the homo-crystallites were nucleated in the matrix of the preformed Sc-crystallites and grew until impinging to others. For the blending compositions where the nucleation of homo-crystallites were absence, only Sc-crystallites were nucleated and continued to grow, giving the size being larger than 20 lm [15]. The Sc-crystallization in PLL/PDL–PEG– PDL copolymer blends, on the other hand, was restricted within the block copolymer phase separation where Sc-crystallites can grow only to some extent until reaching the phase separated domain, resulting in relatively much smaller crystallites (0.064–0.699 lm). The changes in tensile strength and elongation at break in correlation with morphology were monitored as a function of increasing PDL segment length for the films prepared by blending PLL with 40 wt% copolymers. Shown in Fig. 4(c) demonstrates the changes of tensile stress in correlation with Sc-crystallite size while the change of elongation at break in correlation with % amorphous area are shown in Fig. 4(d). The changes in tensile strength and elongation at break were respectively governed by the Sc-crystallite size and % amorphous area. As can be seen in Fig. 4(c) and (d), the highest value of tensile strength was obtained in corresponding with the largest size of Sc-crystallites while the highest value of elongation at break was achieved in corresponding with the largest amorphous area. Fig. 4(e) shows the tensile stress–strain curves of the 60/40 wt% blended films of PLL/copolymers in comparisons with the neat PLL film. Toughness of the films, calculated from the area under the curve, was 2.3 104, 5.6 104, 4.8 104 and 4.5 104 kJ m3 when PLL was blended with PDL53–PEG182–PDL53, PDL81–PEG182–PDL81, PDL126–PEG182–PDL126 and PDL213–PEG182–PDL213, respectively. Among the blended films, the toughness of PLL was increased the most when blending with PDL81–PEG182–PDL81. This behavior may be attributed to the large amorphous area of this blended film, in corresponding to AFM image analysis in that PLL/PDL81–PEG182–PDL81 system demonstrated the largest amorphous area (i.e. 72%) so, as a consequence, stress induced-mobility of the polymer chains was well accommodated in the amorphous region. Also, the sufficient Sc-crystallite size leaded to the effective tensile strength. With these two morphological features, this blended-film underwent relatively large deformation before breaking down.
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Fig. 3. DSC thermograms of PLL/copolymer blended-films prepared from blending PLL with (a) PDL53–PEG182–PDL53, (b) PDL81–PEG182–PDL81, (c) PDL126– PEG182–PDL126 and (d) PDL213–PEG182–PDL213 in various compositions; (e) the second heating run thermograms of the copolymers.
It is noted that the order of increasing tensile strength was directly related to the increase of Sc-crystallite size. PLL/PDL213–PEG182–PDL213 demonstrated the highest values of tensile strength due to its formation of the largest size of Sc-crystallites (Fig. 4(a)). Elongation at break was also proven to rely on the % amorphous area. The increase of % amorphous area resulted in the increase in elongation at break of the films. Owing to its largest % amorphous area, PLL/PDL81–PEG182– PDL81 exhibited the maximum elongation a break (Fig. 4(b)). Although the AFM morphology of the blended-films were investigated only for 60/40 wt% PLL/copolymer films, the changes in stress and elongation at break behaved in a similar fashion at other blending compositions. The AFM study on other blending compositions is currently in progress to establish the composition dependence on the morphology of the films. We hypothesized that the ratio between PDL and PEG segment length in PDL81–PEG182–PDL81 offered the optimum balance between Sc-crystallite size and % amorphous area when it was blended with PLL and, hence, provided the maximum toughness among the prepared films. Such blended-film demonstrated the toughness of about 5.6 104 kJ m3. This value was approximately 400 times larger than the toughness of the neat PLL film. To test the hypothesis, PDL189–PEG455–PDL189 was also synthesized having the ratio between PDL and PEG segment length as similar as in PDL81–PEG182–PDL81. This is to
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Fig. 4. (a) The tensile strength and (b) elongation at break as a function of blending composition for the PLL/copolymer blended-films, (c) the change of tensile stress in correlation with Sc-crystallite size and (d) the change of elongation at break in correlation with % amorphous area plotted as a function of PDL segment length, and (e) the tensile stress–strain curves, reflecting the toughness, of the 60/40 wt% blended-films of PLL/copolymers in comparisons with the neat PLL film.
elucidate whether this specific ratio reflects the optimum balance between the size of Sc-crystallites and % amorphous area, which offered the enhanced toughness of PL film in a similar fashion to PLL/PDL81–PEG182–PDL81 blending system. The synthesis detail of PDL189–PEG455–PDL189 copolymer was shown in supporting information. Fig. 6(a) shows the DSC thermograms of the neat PDL189–PEG455–PDL189 and the 60/40 wt% blended-film of PLL/PDL189–PEG455–PDL189. Similar to PDL81–PEG182–PDL81, two separate melting peaks corresponding to melting temperatures of PEG and PDL segments in PDL189–PEG455–PDL189 were evidenced at 55 °C and 167 °C, respectively. These two melting peaks were disappeared when the copolymer was blended with PLL, which, then, exhibited only the melting temperature of Sc-PL at 206 °C. Stress–strain curves of the PLL/PDL189–PEG455–PDL189 films were illustrated in Fig. 6(b). PLL/PDL189–PEG455–PDL189 films exhibited the behavior of ductile materials with tensile strength, elongation at break and toughness being determined at 36 MPa, 195% and 5.8 104 kJ m3, respectively. AFM image in Fig. 6(c) shows that PLL/PDL189–PEG455–PDL189 film illustrated the spherical Sc-crystallites, as well as other blended films, with the average crystallite diameter of
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Fig. 5. AFM phase images of (a) PLL spherulite and (b) its blown out picture, (c) the neat PDL53–PEG182–PDL53, (d) PLL/PDL53–PEG182–PDL53, (e) PLL/PDL81– PEG182–PDL81, (f) PLL/PDL126–PEG182–PDL126 and (g) PLL/PDL213–PEG182–PDL213; (h) the average diameter and their distribution of Sc-crystallites including the distinguished area of Sc-crystallites and the amorphous phase.
0.699 ± 1.50 lm. When compared with PLL/PDL81–PEG182–PDL81, PLL/PDL189–PEG455–PDL189 exhibited relatively high tensile strength and low elongation at break due to its relatively large size of Sc-crystallites and less % amorphous area, respectively. The toughness of the film, however, was comparable to that of PLL/PDL81–PEG182–PDL81 on account of the similar ratio between PDL and PEG segment length. The results suggested that the specific ratio between PDL and PEG segment length in PDL81–PEG182–PDL81 and PDL189–PEG455–PDL189 copolymers provided the optimum balance between the size of Sc-crystallites and % amorphous area, when blending with PLL, which resulted in the largely-enhanced toughness of the blended films.
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Fig. 6. (a) The DSC thermograms of the neat PDL189–PEG455–PDL189 and the 60/40 wt% blended-film of PLL/PDL189–PEG455–PDL189, (b) stress–strain curves of the PLL/PDL189–PEG455–PDL189 indicating the toughness of the films and (c) AFM phase image of PLL/PDL189–PEG455–PDL189 film and the determined average crystallite diameter and % amorphous area.
It is also noted that the crystallite size of PLL/PDL189–PEG455–PDL189 film was comparable to that of PLL/PDL213–PEG182– PDL213 film. Therefore the tensile strength values of the two are indifferent. However, owing to its noticeably large amorphous regime, PLL/PDL189–PEG455–PDL189 film demonstrated the significantly higher elongation at break. 4. Conclusions The Sc-PL films were prepared by blending the PDL–PEG–PDL copolymers, varying PDL segment lengths, with the neat PLL. The DSC results showed that the complete sterocomplexation was reached at 60/40 wt% PLL/copolymer blending composition for almost blended films. In addition, the crystallization of PEG was interrupted and, thus, prohibited by either adding PDL or PLL in the system. AFM images showed that, for the first time, the Sc-crystallites were formed and, then, dispersed in the amorphous of PL and the non crystallizing PEG when the copolymers were blended with the neat PLL. The increase of PDL segment length in the copolymer leaded to the increase of the Sc-crystallite size, which, consequently, resulted in the increase of the tensile strength of the blended-films. Elongation at break of the films was found to rely on the determined amorphous area. To maximize the toughness of the films, the specific ratio of PDL and PEG segment length must be achieved to provide the optimum balance between Sc-crystallite size and % amorphous area. Acknowledgements We gratefully acknowledge the financial support from the PTT Research and Technology Institute (PTT-RTI). This work was also supported by the National Research University Project under Thailand’s Office of the Higher Education Commission. Appendix A. Supplementary material Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/j.eurpolymj.2015.06.015.
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