Skip to content
Publicly Available Published by De Gruyter December 23, 2016

Preparation and characterization of polyethylene glycol/poly(L-lactic acid) blends

  • Ioanna-Georgia Athanasoulia and Petroula A. Tarantili EMAIL logo


The effect of incorporation of poly(ethylene glycol) (PEG) on thermomechanical and hydrophilicity properties of poly(L-lactic acid) (PLLA) was investigated. PEG/PLLA blends, containing 10, 20, 30 and 40 wt% PEG, were prepared by melt-extrusion in a co-rotating twin-screw extruder. By DSC analysis, it was observed that the Tg of PLLA phase in PEG/PLLA blends decreased accompanied by a significant decrease in Tcc and increase in their melting enthalpy. Therefore, the addition of PEG enhances the crystallization ability of PLLA phase due to its lubricating effect which increased mobility of PLLA chains. From TGA it was observed that low concentrations of PEG (10 & 20 wt%) increase the Tonset of thermal degradation, probably due to improved heat resistance of the crystalline phase. At higher PEG content, the Tonset decreases, as the lubricating effect becomes the controlling mechanism for the initiation of degradation process. Decrease in tensile strength and modulus was recorded especially in PLLA blends with PEG content higher than 20 wt%. The elongation at break decreases reaching a maximum at 20 wt% PEG and then dropped again. To investigate the effect of PEG on the wetting ability of PLLA, water contact angle measurements were performed. The results indicate that the introduction of PEG lowers the contact angle values in PEG/PLLA film surfaces, as compared to pure PLLA, suggesting improved hydrophilic properties.


Poly(ethylene glycol) (PEG) and poly(L-lactic acid) (PLLA) are very important semicrystalline polymers. PEG shows hydrophilicity and biocompatibility. PLLA is a biodegradable polyester and has attracted increasing attention due to its potential applications as biomedical and environmentally friendly material. Nevertheless, the applications have been restricted due to the brittleness, hydrophobicity, low melt strength, poor heat resistance, slow degradation rate and crystallization kinetics of PLA [1, 2]. The above limitations can be improved through several approaches including copolymerization, blending, plasticization and incorporation of fillers [1, 3, 4]. As compared to copolymer synthesis, the blending technique is a simple and more easily affordable way. Polymer blends offer the advantage of a large spectrum of mechanical properties and degradation rates by simply varying the component ratio of the other phase of PLA.

Plasticizers have been employed extensively to improve processability, flexibility and impact toughness of glassy polymers. The plasticizer should not be prone to migration because this would cause contamination of the materials in contact with the plasticized PLA. It is well-known that for instance the monomer, lactide itself, is considered as one of the best plasticizers for PLA [5, 6], but it has the disadvantage of migrating very rapidly at the polymer surface. One of the most efficient plasticizers for PLA is PEG, a crystallizable, thermoplastic polymer which is produced by anionic ring opening polymerization of ethylene oxide. It possesses excellent biocompatibility, nontoxicity, hydrophilicity, lubrication capability, bondability, dispersibility and high mobility having large exclusion volumes in water. In addition, PEG has good solubility in H2O and many organic solvents [7]. In view of its aforementioned advantages, PEG has been widely used in pharmaceutical and biomedical fields, cosmetics, chemical industry and food processing.

It is well known that PEG has very good miscibility with PLA because the terminal hydroxyl groups in PEG molecules can react with the carboxyl groups in PLA molecules [1]. Several authors revealed that PEG acts as an effective polymeric plasticizer to facilitate the crystallization rates of the PLA [1, 8, 9]. With the addition of PEG, the cold crystallization temperatures of PLA molecules decreased remarkably as the PEG accelerates the crystallization ability of PLA. Sungsanit et al. [6] used PEG (Mw=1000 g/mol) to plasticize linear PLA (L-PLA). They observed that the crystallinity, elongation at break, and impact strength of L-PLA/PEG blends increased, while the glass transition temperature, tensile strength and modulus lowered with increasing PEG content. Moreover, they also noticed that the PEG phase separated from the L-PLA/PEG blends when PEG content exceeded 10 wt%. Baiardo et al. [10] plasticized PLA using PEGs with different molecular weights (Mw=400–10 000 g/mol). They found that the solubility limit and plasticizing efficiency of PEG in PLA decreased with the increase of PEG molecular weight. Lower molecular weight PEG could more significantly reduce Tg. Furthermore, they also noted that the strain at break of PLA/PEG blends increased sharply, while the tensile strength and elastic modulus dropped rapidly with increasing PEG content when the Tg of the blends was close to room temperature. Jacobsen and Fritz [5] investigated the properties of plasticized PLA, with 2.5–10 wt% of PEG (molecular weight (Mw=1.5×103 gmol−1) and observed decreased tensile strength and modulus with an increase in percentage of elongation at break. Also, increased impact resistance to about five times more as compared to that of a pure PLA observed with the addition of 10 wt% of PEG. Kulinski and Piorkowska [11] reported that incorporation of 10 wt% of PEG, within amorphous plasticized PLA results in considerable deformation to about 550%, while semicrystalline PLA exhibits non-uniform plasticization of the amorphous phase and showed less ability to the plastic deformation. Further due to the difference in crystallization temperature of PLA matrix and PEG plasticizer, the PLA/PEG blends are miscible in the amorphous phase. The mechanical properties of PLA/PEG blends are becoming lost over time which results in crystallization and phase separation [3]. Several authors have revealed that PEG can act as an effective polymeric plasticizer to facilitate the crystallization rates of the PLA [11]. With the addition of PEG, the cold crystallization temperatures of PLA molecules decreased remarkably as the PEG accelerates the crystallization ability of PLA [1].

One common feature of PEG appears to be the water-solubility. It is soluble also in many organic solvents including aromatic hydrocarbons (not aliphatics). They are used to make emulsifying agents and detergents, and as plasticizers, humectants, and water-soluble textile lubricants. This molecule is the focus of much interest in the biotechnical and biomedical communities. The molecular weight commonly used in biomedical and biotechnology applications ranges from a few hundred to approximately 20 000 g/mol [8]. For tissue engineering applications cell attachment and proliferation are closely related with the wettability of the scaffolds surface. However, the cytocompatibility of PLA is not desirable because of their hydrophobicity [12].

In this work, the effect of PEG on the crystals structure, thermal, mechanical and surface properties of a semi-crystalline PLLA matrix was studied. PLLA blends with different PEG concentrations (10, 20, 30 and 40 wt%) were melt-mixed in a co-rotating twin-screw extruder and characterized by XRD, DSC, TGA and Contact Angle Measurements experiments.



The PLLA used in this study was purchased from Nature Works Co. (USA). More specifically, the grade Ingeo PLLA 4043D is a semicrystalline material in pellet form, with a D-isomer content of approximately 2%. Poly(ethylene glycol) (PEG) powder was purchased from Merck (MO, USA) with molecular weight 10 000 g/mol.

Preparation of PLLA based blends

PEG/PLLA blends with compositions 100/0, 10/90, 20/80, 30/70 and 40/60 w/w were prepared by extrusion melt blending in a Haake co-rotating twin-screw extruder with L/D=25 and 16 mm diameter (Haake Polylab System, Haake Rheomex PTW16, Thermo Electron Corporation). The extruder was heated at five zones along the cylinder and the die. The temperature profiles of the barrel from the hopper to the die were 180–180–185–185–190–195°C for pure PLLA and 150–150–155–155–160–165°C for the PEG/PLLA blends and the screw speed 30 rpm. PLLA and PEG were dried in a vacuum oven at 80°C and 40°C, respectively for 4 h, in order to eliminate hydrolytic degradation reactions. Prior to processing, pre-weighed quantities of PLLA and PEG were dry-mixed. After melt extrusion, the obtained material, in the form of continuous strands, was granulated into regular, cylindrical pellets using a Brabender knife pelletizer.

X-ray diffraction (XRD)

X-ray diffraction analyses of PEG, PLLA and PEG/PLLA blends were made using a BRUKER D8-ADVANCE (twin/twin) diffractometer (40 kV, 40 mA) with a Cu X-ray tube (λ=1.5418 Å). The recording speed was 0.02°/s in the range of angles of 2θ=(5–50)°. Samples for X-ray analysis were obtained from compression-molded films.

Differential scanning calorimetry (DSC)

DSC measurements were run in a Mettler Toledo model DSC 1 differential scanning calorimeter with pure indium as a calibration standard. A total of 8–10 mg of each sample was heated from –40°C to 190°C and equilibrated at 190°C for 3 min to completely eliminate previous thermal history. The samples were then cooled to –40°C at a cooling rate of 10°C/min and kept at –40°C for 3 min to expose all samples to the same thermal treatment. Finally, they were reheated to 190°C at a rate of 10°C/min.

All runs were conducted under nitrogen flow to avoid thermo-oxidative degradation.

Thermogravimetric analysis

TGA measurements were performed in a Mettler Toledo (TGA-DTA model) thermal gravimetric analyzer. A total of 8–10 mg was heated from 25 to 600°C, at a heating rate of 10°C/min, in a nitrogen atmosphere. Three measurements were carried out for each sample.

Tensile properties measurement

Tensile tests were carried out by using an Instron tensometer (4466 model), equipped with a load cell of maximum capacity of 1.0 kN. The compression-molded tensile specimens were cut into dog-bone shape following ASTM D638 standard [13]. A crosshead speed of 10 mm/min was used at room temperature. Tensile strength, Young’s modulus and elongation at break were evaluated from the stress-strain data. Each sample included six tested replicates to obtain a reliable mean and standard deviation. The dimensions of the samples were: distance between clamps 64 mm, width of narrow section 5.5 mm, and thickness 2.1 mm.

Tensile strength (σs ) and Young’s modulus (E) were determined according to the following equations (Eq. (1 and 2)) [14]:


where F: force exerted on an object under tension, L0: original length, A: cross section area, ΔL: length of the object changes.

Contact angle measurements

Contact angle measurements were performed to evaluate the materials’ wettability. The sessile drop method [15] was used to measure the contact angle by depositing a deionized water droplet (2 μL) at room temperature on the surface of each compression-molded polymer film with the following dimensions (2×2×0.1 cm3). The angle of the water droplet was captured using a CCD digital video camera. The software for the image processing and the calculation of the contact angle value was performed using image analysis software NI-IMAQ Vision along with LabVIEW and Matlab. Six measurements were performed in each sample.

Scanning electron microscopy (SEM)

The morphology of PEG/PLLA blends was examined with scanning electron microscopy (SEM) in a JEOL 2000 Microscope (Tokyo, Japan) at an accelerating voltage of 12.5 kV. The samples were submerged in liquid nitrogen and broken down. The fractured surfaces were coated with a thin layer of gold prior to observation.

Results and discussion

X-ray diffraction (XRD)

X-ray diffraction (XRD) was applied to study the effect of PEG10 000, on the crystalline structure of PLLA, and the obtained spectrograms of pure PLLA, PEG and PEG/PLLA blends are shown in Fig. 1.

Fig. 1: X-ray diffraction patterns of pure PLLA and PEG/PLLA blends.
Fig. 1:

X-ray diffraction patterns of pure PLLA and PEG/PLLA blends.

For melt-extruded pure PLLA and blends with low PEG content (10, 20 wt%) reflection of PLLA crystals cannot be detected [16, 17]. When PEG was incorporated into the PLLA matrix at 30 wt%, one broad diffraction peak with low intensity appeared at 16.6°, either representing platelets with different dimensions [18] or α′-crystalline phase with lower packing density [19] due to incomplete crystallization. The intensity of this peak, in systems with 40 wt% PEG, increased and shifted to 16.8°. In addition, three more peaks appeared: at 2θ=14.9°, 16.8°, 19.1° and 22.3°, corresponding to the lattice planes (010), (200)/(110), (203) and (015), respectively, which are characteristic of the α-crystalline phase of PLLA [8, 20, 21, 22]. The intensity of these peaks indicates the enhanced ability of plasticized PLLA to induce α-crystalline structure. Additionally, at higher PEG loadings (40 wt%), new characteristic diffraction peaks of the PEG phase appeared at 2θ=19.2° and 23.2°, attributed to the lattice planes of (115) and (016) [23].

Thermal properties analysis

The thermal transitions of PLLA and PEG/PLLA blends, prepared by melt extrusion, were assessed by DSC experiments. The DSC thermograms during the first heating cooling and second heating cycles were presented in Figs. 24. From the second heating cycle (Fig. 4), the glass transition (Tg), cold crystallization (Tcc) and melting (Tm) temperatures of PEG and PLLA phases for each PEG/PLLA blend composition were evaluated and the results are presented in Tables 1 and 2.

Fig. 2: DSC thermograms of pure PLLA and PEG/PLLA (w/w) blends, during the first heating cycle.
Fig. 2:

DSC thermograms of pure PLLA and PEG/PLLA (w/w) blends, during the first heating cycle.

Fig. 3: DSC thermograms of pure PLLA and PEG/PLLA (w/w) blends, during the cooling cycle.
Fig. 3:

DSC thermograms of pure PLLA and PEG/PLLA (w/w) blends, during the cooling cycle.

Fig. 4: DSC thermograms of pure PLLA and PEG/PLLA (w/w) blends, during the second heating cycle.
Fig. 4:

DSC thermograms of pure PLLA and PEG/PLLA (w/w) blends, during the second heating cycle.

Table 1:

DSC results for PLLA phase during cooling and 2nd heating cycle of PEG/PLLA blends.

Table 2:

DSC results for PEG phase during cooling and 2nd heating cycle of PEG/PLLA blends.

PEG/PLLA (w/w)TcPEG (°C)TmPEG (°C)ΔHcPEG (J/g)ΔHmPEG (J/g)

The existence of α′ and α crystalline forms can be verified by XRD and DSC analysis [24]. In the first heating cycle, a small endothermic peak was observed immediately before the main melting peak for pure PLLA and 10 wt% PEG/PLLA blend, but not in blends with higher PEG content (Fig. 2). This peak, in the lower temperature area, is related with melting of the α′-crystals and its recrystallization into the α crystalline form, while the second peak at higher temperature corresponds to the melting of the α form. The two processes, i.e. melting of the α′ form and recrystallization into the α form, can be considered as the α′-α crystalline phase transition [22]. This bimodal melting peak of PLLA became a single melting peak at higher PEG content (>20 wt%).

Comparing the two polymers, crystallization of PLLA takes place at higher temperatures than for PEG (Tables 1 and 2). At higher PEG content (30 & 40 wt%) an enhancement in the crystallization ability of PLLA (Fig. 3) was observed. In general, plasticizers interact with the polymer chains on the molecular level and speed up its visco-elastic response [25]. The addition of a plasticizer increased the polymer chain mobility and accelerated the crystallization rate by reducing the energy required during crystallization for the chain folding process [25, 26]. On the other hand, it was reported that during the formation of PLLA crystals, PEG molecules could be slowing down the formation of PLLA crystalline structures because they interfere with spherulites, growth since they are present in fold surfaces of crystalline lamellae [26]. In addition, some amounts of PEG could be trapped in the intra-spherulitic region of PLLA leading to hindering of PLLA crystallization [8].

From Fig. 3 it is observed that crystallization of PLLA phase during cooling from melt (at a cooling rate of 10°C min−1) took place only in blends with higher amount of PEG (30 and 40 wt%), whereas cold crystallization was observed in all the examined PEG/PLLA (w/w) blends. These blends in which crystallization of PLLA phase during cooling was detected, presented significant lower crystallization during the heating cycle. Incorporation of PEG facilitates the cold crystallization process, decreasing significantly the values of Tcc of PLLA (Fig. 4, Table 1). The above depression of Tcc, especially at 30 and 40 wt% PEG content, is an indication about compatibility between PEG and PLLA phases. Similar results of Tcc decrease in parallel with a shift of Tg were also observed by Wang et al. [27], who suggested that PEG induced an easier rearrangement of polymer chains to crystalline structure. The increase of melting enthalpy of PLLA (Table 1) confirms that the enhanced chain mobility, due to the lubricating effect of PEG molecules, improves the crystallization ability of PLLA in their blends [9, 28, 29].

Crystallization and melting of PEG was not detectable in samples with low PEG concentrations (10 and 20 wt%) (Figs. 3 and 4). Since in a homogeneous polymer blend the heat flow is easily dissipated, the heat intake of PEG melting is absorbed by the neighboring PLLA chains and, as a consequence, the endothermic peak disappears in the 57–65°C region. However, the melting peaks of PEG became obvious at PEG loadings of 30 and 40 wt% in PLLA blends (Fig. 4). The melting temperature of PEG phase in the above blends decreases, whereas that of PLLA remains almost constant with PEG content above 10 wt% (Table 1) suggesting that this property was not greatly affected by the addition of PEG. According to Sheth et al., who studied the thermal properties of PLLA blends with PEG (Mw=20 000 g/mol), when each component exceeds 20 wt%, the blends were able to crystallize into two semi-miscible crystalline phases that are dispersed in an amorphous matrix, the PEG-dominated phase melts immediately after the glass transition of the amorphous matrix. For other compositions, only the major component is able to crystallize. These results indicate that the blends are semi-miscible, exhibiting some chain interpenetration [9].

Table 1 shows that the incorporation of PEG (10 wt%) in melt extruded PEG/PLLA blends, decreased Tg of PLLA phase, which can be explained by the fact that melted PEG chains penetrated into the non-crystalline region of PLLA increasing molecular chain mobility and allowing chain rearrangement. At 20 wt% PEG content, the glass transition of the 20/80 w/w PEG/PLLA blend occurs over a very broad temperature range (Fig. 4) and, therefore, it would be safe to conclude that the compatibility does not extend all the way to the molecular level [9]. The miscibility of the components of polymer blends might be concluded by the appearance of a single Tg, normally found between the Tg’s of pure components in the related thermograms. Under the experimental conditions of this work, the Tg of PEG was not determined within the range of the related DSC scans. However, the observed significant depression of Tg values for the phase of PLLA in the blends (10 & 20 wt%) is an evidence about the partial miscibility between PLLA and PEG.

At higher PEG content (30 & 40 wt%) the Tg’s were less clearly defined, owing to overlap of the PLLA glass transition with the PEG melting peak.

Depression of Tm values of PLLA phase was observed only in blends with 10 wt% PEG, which is an evidence of some partial miscibility between the two semicrystalline polymers at that composition [30].

TGA analysis

The thermal stability of pristine PLLA and PEG/PLLA blends was investigated by thermogravimetric analysis (TGA). The curves showing the weight loss and its derivative (DTG) versus temperature for PEG/PLLA blends, which provide information about the mechanism and extent of degradation of the examined materials, are presented in Fig. 5. Based on the above graphs the onset temperature (Tonset) and the temperature of maximum thermal degradation rate (Tpeak) were evaluated (Table 3).

Fig. 5: The TGA (a) and DTGA (b) curves for pure PLLA and PEG/PLLA (w/w) blends.
Fig. 5:

The TGA (a) and DTGA (b) curves for pure PLLA and PEG/PLLA (w/w) blends.

Table 3:

TGA results of PEG/PLLA blends.

PEG/PLLA (w/w)Tonset (°C)Tpeak (°C)Tpeak (°C)

From Fig. 5 it is observed that thermal decomposition of PEG took place at higher temperatures than for PLLA. Incorporation of PEG at 10 and 20 wt% content, resulted in an increase of the initial thermal degradation temperature of PLLA phase. This effect is related with the increased crystallinity of the polyester due to the presence of PEG, as it was confirmed by DSC experiments. In blends with higher PEG concentrations i.e. 30 and 40 wt%, the thermal decomposition was analysed into two peaks and degradation of PLLA phase shifted to lower temperatures (Table 3, Fig. 5). Therefore, it seems that at higher PEG content, the lubrication of PLLA molecules became the main mechanism which facilitates its thermal decomposition. In agreement with the above observation, Mohapatra et al. mentioned that plasticization of PLLA by PEG tends to reduce the total crystallinity of the polymer and this reduced crystallinity can make the polymer easier to degrade [28]. Regarding thermal decomposition of PEG phase, it remains almost unaffected by the presence of PLLA molecules [31].

Tensile testing

The tensile properties of pure PLLA and PLLA blends with 10, 20, 30 and 40 wt% of PEG were examined and the corresponding stress-strain curves are shown in Fig. 6. The incorporation of PEG in PLLA matrix significantly decreases the tensile strength and modulus of elasticity of the prepared blends, for PEG content above 20% (Table 4, Fig. 6).

Fig. 6: Stress-strain curves of PEG/PLLA blends.
Fig. 6:

Stress-strain curves of PEG/PLLA blends.

Table 4:

Tensile test results of PEG/PLLA blends.

PEG/PLLA (w/w)Tensile strength (MPa)Young’s Modulus (MPa)Strain at break (%)

Generally, the addition of plasticizers reduces the tensile modulus of PLLA, being a stiff, brittle and with limited extendibility material due to extensive intermolecular forces [14]. PEG/ PLLA blends are able to increase the free volume between polymeric chains. In doing so, the ease of movement of polymeric chains with respect to each other is significantly enhanced [29]. This is because the action of plasticizer molecules weakens polymer–polymer intermolecular interactions and imparts flexibility to the PLLA, leading to a decrease of tensile modulus with the increase of PEG content [29].

Regarding the strain at break, a significant increase was observed in blends with PEG concentrations of 20 and 30 wt% reaching a maximum of 30% for PLLA blends with 20 wt% PEG [14, 32]. Increase in elongation at break was also reported by Mohapatra et al. [28] for specimens with PEG6000 content from 10 to 30 wt%. At 40 wt% PEG the elongation at break drops again causing the blended material to regain the brittleness of pure PLLA as the strength of the material was significantly decreased. Similar study of the tensile properties of PEG/ PLLA blends by Li et al. [29] revealed a decline in strain at break with PEG content of 15 wt%. This effect was attributed to the lack of cohesion between the polymer matrix and PEG, which may lead to phase separation and increases brittleness of the blend. Hu et al. [33] also prepared plasticized PLLA using PEG8000 and observed a similar trend, with an increase in fracture strain and a reduction in modulus and stress. Sheth et al. [9] reported that above 50% PEG in the PLLA /PEG blend, the tensile modulus increased due to the increased crystallinity of PEG plasticizer.

The presence of higher PEG concentrations in the blends resulted in a decrease of the recorded percentage of elongation at break (40 wt% PEG) since PEG is prone to migration at higher concentrations to the polymer’s surface, causing a stiffening of the blends. At high PEG contents, the semicrystalline PLLA matrix of the system exhibited a non-uniform plasticization and less ability to the plastic deformation. This effect can also be attributed to the lack of cohesion between the polymer matrix and PEG that contributed in phase separation and induced brittle characteristics in the blend.

Contact angle measurements

Water contact angle measurements were performed for direct characterization of the blends hydrophilicity, and the obtained results are presented in Table 5. By increasing the PEG content in the PLLA blend, the contact angle values decreased, up to 30 wt% PEG composition. The value obtained for the 30/70 w/w PEG/PLLA blend is an analog evidence of PEG aggregation on the surface. The water droplet permeation into the hydrophilic polymer surface could affect the contact angle measurements especially at high PEG content. Therefore, time repeatability is critical for the accuracy of the obtained results [12]. From measurements in PEG/PLLA films with higher PEG concentrations (40 wt%) it was observed some dissolution of the PEG located on the surface into the water droplet and, therefore, the PEG concentration left in the matrix is not quite helpful on its hydrophilicity [12].

Table 5:

Contact angle measurements of PLLA and PEG/PLLA blends.

PEG/PLLA (w/w)Contact angle (°)


Micrographs of cryogenically fractured surfaces of PEG/PLLA are shown in Fig. 7. A smooth surface with less deformation as the amount of PEG in the blend increases can be seen. No clear evidence about the different phases was observed, which suggests good compatibility between the two polymers in the examined compositions.

Fig. 7: SEM micrographs of (a) 10/90, (b) 20/80 and (c) 30/70 w/w PEG/PLLA blends.
Fig. 7:

SEM micrographs of (a) 10/90, (b) 20/80 and (c) 30/70 w/w PEG/PLLA blends.

Concluding remarks

DSC analysis in extruded PEG/PLLA blends revealed no obvious crystallization during the cooling cycle of pure PLLA and blends with low PEG content (10 & 20 wt%), whereas exothermic crystallization appeared in systems with higher PEG content (30 & 40 wt%). The Tg of PLLA in PEG/PLLA blends was found to strongly decrease, with a significant decrease of Tcc at the same time. Therefore, it seems that the presence of PEG enhances the crystallization ability of PLLA phase in their blends, whereas that of PEG phase remains unaffected. The melting peak of pure PLLA and 10/90 (w/w) PEG/PLLA during the first heating cycle presented progressive melting of α′ and α crystalline form PLLA which became more symmetric with increasing PEG content. According to TGA experimental data, the incorporation of PEG at low content (<20 wt%) resulted in an increase of the initial thermal degradation temperature of PLLA and this effect can be related with the increased crystallinity of the polyester phase because of the action of PEG. At higher PEG contents (30 & 40 wt%), the thermal decomposition of PLLA phase shifted to lower temperatures due to the lubricating effect of PEG molecules, which makes the PLLA chains more vulnerable to degradation. Regarding the thermal decomposition of PEG phase, it remains almost unaffected in the blends with PLLA. The tensile strength and modulus were significantly decreased, especially for blends with PEG content higher than 30% wt%, whereas the strain at break increased reaching a maximum at 20 wt% PEG content. At 40 wt% PEG content it decreases again probably to phase separation and inhomogeneity phenomena. Regarding the surface wettability, the addition of a highly hydrophilic polymer, such as PEG, was proven advantageous for the modification of PLLA’s inherently hydrophobic nature, as a significant decrease in the contact angle values of the blends was observed, especially in, samples with 10–30 wt% PEG content.

The exploration of the effect of PEG’s concentration on the thermomechanical properties of the investigated blends, in combination with their hydrophobic-hydrophilic balance, is of main importance to find the recommended PEG content and to design their production process. Through the appropriate choice of the composition of the second component in the PLLA blend, a tailor made material with specific characteristics which comply with the desired application can be obtained.

Article note:

A collection of invited papers based on presentations at the 16th International Conference on Polymers and Organic Chemistry (POC-16), Hersonissos (near Heraklion), Crete, Greece, 13–16 June 2016.


We would like to acknowledge the Research Committee of National Technical Univ. of Athens for the scholarship of Ms. Athanasoulia PhD. Special thanks go to Dr. D. Korres for assistance in DSC and TGA experiments, Dr. N. Panagiotou for assistance in X-Ray Diffraction experiments, Dr. P. Schinas for assistance in SEM as well as to the Assistant Professor A. Papathanasiou and PhD student N. Chamakos for their help in contact angle measurements.


[1] W. C. Lai, W. Liau, T. T. Lin. Polymer45, 3073 (2004).10.1016/j.polymer.2004.03.003Search in Google Scholar

[2] L. Wang, X. Jing, H. Cheng, X. Hu, L. Yang, Y. Huang. Ind. Eng. Chem. Res.51, 10088 (2012).10.1021/ie300526uSearch in Google Scholar

[3] K. Nakane, Y. Hata, K. Morita, T. Ogihara, N. Ogata. J. Appl. Polym. Sci.94, 965 (2004).10.1002/app.20959Search in Google Scholar

[4] D. Battegazzore, S. Bocchini, A. Frache. eXPRESS Polym. Letters5, 849 (2011).10.3144/expresspolymlett.2011.84Search in Google Scholar

[5] S. Jacobsen, H. G. Fritz. Polym. Eng. Sci.39, 303 (1999).Search in Google Scholar

[6] K. Sungsanit, N. Kao, S. N. Bhattacharya. Polym. Eng. Sci.52, 108 (2012).10.1002/pen.22052Search in Google Scholar

[7] M. S. Thompson, T. P. Vadala, M. L. Vadala, Y. Lin, J. S. Riffle. Polymer49, 345 (2008).10.1016/j.polymer.2007.10.029Search in Google Scholar

[8] T. Nazari, H. Garmabi. J. Appl. Polym. Sci.133, 1 (2016).Search in Google Scholar

[9] M. Sheth, R. Kumar RA, V. Dave, R. Gross, S. P. Mc Carthy. J. Appl. Polym. Sci.66, 1495 (1997).10.1002/(SICI)1097-4628(19971121)66:8<1495::AID-APP10>3.0.CO;2-3Search in Google Scholar

[10] M. Baiardo, G. Frisoni, M. Scandola. J. Appl. Polym. Sci.90, 1731 (2003).10.1002/app.12549Search in Google Scholar

[11] Z. Kulinski, E. Piorkowska. Polymer46, 10290 (2005).10.1016/j.polymer.2005.07.101Search in Google Scholar

[12] X. Zhu, T. Zhong, R. Huang, A. Wan. J. Biomaterials Sci., Polym. Ed.26, 1287 (2015).Search in Google Scholar

[13] ASTM D638-14, Standard Test Method for Tensile Properties of Plastics, ASTM International, West Conshohocken (2003).Search in Google Scholar

[14] F. T. M. Noori, N. A. Ali. I. J. Applic. Innov. Eng. Manag. (IJAIEM) 3, 459 (2014).Search in Google Scholar

[15] J. Drelich. Surf. Innov.1, 248 (2013).10.1680/si.13.00010Search in Google Scholar

[16] H. Zhou, T. Green, Y. Joo. Polymer47, 7497 (2006).10.1016/j.polymer.2006.08.042Search in Google Scholar

[17] G. Perego, G. Cella. J. Appl. Polym. Sci.59, 37 (1996).10.1002/(SICI)1097-4628(19960103)59:1<37::AID-APP6>3.0.CO;2-NSearch in Google Scholar

[18] B. W. Chieng, N. A. Ibrahim, W. M. Z. W. Yunus, M. Z. Hussein. Polymers6, 93 (2014).10.3390/polym6010093Search in Google Scholar

[19] S. Y. Huang, H. F. Li, S. C. Jiang, X. S. Chen. Polymer52, 3478 (2011).10.1016/j.polymer.2011.05.044Search in Google Scholar

[20] H. Xu, G. J. Zhong, Q. Fu, J. Lei, W. Jiang, B. S. Hsiao, Z. M. Li. Appl. Mater. Interf.4, 6774 (2012).10.1021/am3019756Search in Google Scholar

[21] J. You, W. Yu, C. Zhou. Ind. Eng. Chem. Res.53, 1097 (2014).10.1021/ie402358hSearch in Google Scholar

[22] T. Tábi, I. E. Sajó, F. Szabó, A. S. Luyt, J. G. Kovács. eXPRESS Polym. Letters4, 659 (2010).10.3144/expresspolymlett.2010.80Search in Google Scholar

[23] Y. L. Li, X. X. Li, F. M. Xiang, T. Huang, Y. Wang. Polym. Adv. Technol.22, 1959 (2011).10.1002/pat.1702Search in Google Scholar

[24] P. Pan, B. Zhu, W. Kai, T. Dong, Y. Inoue. Macromolecules41, 4296 (2008).10.1021/ma800343gSearch in Google Scholar

[25] R. Song, R. Xue, L. He, Y. Liu, Q. Xiao. ChineseJ. Polym. Sci.26, 621 (2008).10.1142/S0256767908003357Search in Google Scholar

[26] Y. Song, D. Wang, N. Jiang, Z. Gan. Sustain. Chem. Eng.3, 1492 (2015).10.1021/acssuschemeng.5b00214Search in Google Scholar

[27] B. Y. Wang, S. Z. Fu, P. Ni, J. R. Peng, L. Zheng, F. Luo, H. Liu, Z. Y. Qian. J. Biomed. Mater. Res. A100, 441 (2012).10.1002/jbm.a.33264Search in Google Scholar

[28] A. K. Mohapatra, S. Mohanty, S. K. Nayak. Polym. Comp.35, 283 (2014).10.1002/pc.22660Search in Google Scholar

[29] F. J. Li, J. Z. Liang, S. D. Zhang. J. Polym. Environm.23, 407 (2015).10.1007/s10924-015-0718-7Search in Google Scholar

[30] F. J. Li, L. C. Tan, S. D. Zhang, B. Zhu. J. Appl. Polym. Sci.133, 1 (2016).10.1002/app.44339Search in Google Scholar

[31] J. P. Lee, G. C. Pimentel. J. Chem. Phys.75, 4241 (1981).10.1063/1.442652Search in Google Scholar

[32] B. S. Park, J. C. Song, D. H. Park, K. B. Yoon. J. Appl. Polym. Sci.123, 2360 (2012).10.1002/app.34823Search in Google Scholar

[33] Y. Hu, M. Rogunova, V. Topolkaraev, A. Hiltner. Polymer44, 5701 (2003).10.1016/S0032-3861(03)00614-1Search in Google Scholar

Published Online: 2016-12-23
Published in Print: 2017-1-1

©2016 IUPAC & De Gruyter. This work is licensed under a Creative Commons Attribution-NonCommercial-NoDerivatives 4.0 International License. For more information, please visit:

Downloaded on 21.9.2023 from
Scroll to top button