Fabrication of super tough poly(lactic acid)/ethylene-co-vinyl-acetate blends via a melt recirculation approach: static-short term mechanical and morphological interpretation

Abstract

The mechanical properties such as tensile strength, tensile modulus, elongation-at-break and impact strength of poly(lactic acid) (PLA)/ethylene-co-vinyl-acetate copolymer (EVA, vinyl acetate content 50 weight percent) blends were evaluated at EVA volume fractions ranging from 0–0.35. The tensile properties were compared using several theoretical models. The blends lost little of their tensile strength and modulus while elongation-at-break was simultaneously enhanced. Efficient dispersion of EVA in PLA using a micro compounder in which there is provision for melt recirculation significantly improved the Izod impact strength making the blends super tough. The phase miscibility, two phase morphology, fibrillation and interparticle distance were studied using scanning electron microscopy (SEM). The blend is a two phase system where the particle size is enhanced upon an increase in the concentration of the blending copolymer. The normalized values of the relative elongation-at-break and Izod impact strength were enhanced significantly in accordance with the crystallinity, 33 fold (53.73 kJ mm⁻²) at a 0.35 volume fraction of EVA, which indicated softening of the system with enhanced toughness.

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1. Introduction

Increasing concerns surrounding environmental pollution and sustainability issues associated with petroleum based non-bio-degradable plastics have motivated researchers to develop bio-degradable polymers.^1,2 Poly(lactic acid) (PLA), which degrades biologically, is derived from renewable sources and has drawn significant attention in recent years.^2–4 PLA exists in L- and D-forms which are optical isomers. PLA containing high concentrations of L-isomer is highly crystalline. However, the crystallinity and bio-degradability of PLA depend on the content of the D-form isomer.⁵ On the other hand, the processability of PLA is quite inferior to that of the general polyolefins. Despite numerous advantages such as high strength, high stiffness, good bio-compatibility, high transparency and excellent biodegradability,⁵ PLA suffers from major disadvantages notably brittleness (low strain at break and high modulus), a low heat distortion temperature (HDT, <60 °C), poor impact strength and a low rate of crystallization^5–7 which limit its applicability.

In order to overcome the drawbacks of PLA, various strategies have been employed, such as copolymerization, plasticization and blending with other polymers or rubbers.⁸ Among all these approaches blending of PLA with soft and tough polymers is the most effective and convenient way to toughen PLA.^7,9,10 Blending of PLA with immiscible polymers or partially miscible polymers leads to a significant improvement in the impact strength of PLA without trading off on stiffness.¹¹ It is well known that in immiscible polymer blends a degree of compatibilization can be achieved by addition of a pre-made polymer with intermediate surface energy or by in situ reaction of polymers during melt blending.^12–15

Blending of PLA with other polymers has been the focus of significant attention from researchers and engineers.^8,16,17 Previous studies on PLA blends have mainly focused on rheological properties and miscibility.¹⁶ Many studies have been reported concerning the blending of PLA to improve its performance. The blending polymers include poly(butylene succinate) (PBS),¹⁸ polyurethane (PU),^19–21 polyethylene (PE),^12,22,23 poly(vinyl acetate) (PVAc),^24,25 poly(methyl methacrylate) (PMMA),²⁶ poly(3-hydroxybutyrate) (PHB),²⁷ polycaprolactone (PCL),²⁸ poly(butylene adipate-co-terephthalate) (PBAT),¹³ acrylonitrile-butadine-styrene (ABS),²⁹ glycidyl methacrylate (GMA),³⁰ poly(ethylene-co-octene) (TPO),³¹ and poly(β-hydroxybutyrate-co-β-hydroxyvalerate) (PHBV).¹³ Most of these blends are immiscible resulting in complete phase separation showing a limited improvement in the toughness. However, some systems are reported to have miscibility or partial miscibility, for example blends of PLA with PMMA and PVAc.

It has been shown that PLA is miscible with PVAc.^24,25 The copolymer EVA contains vinyl acetate monomer. Thus the interaction between PLA and EVA may depend on the variation of VA content in the latter. Therefore, compatibility between PLA and EVA can be achieved by tuning the VA content without the need of an additional compatibilizer. Ma et al.⁶ reported that the maximum toughness of PLA was achieved when the vinyl acetate content in EVA was maintained between 50 to 60 wt (%). Moreover, in this range, the compatibility between PLA and EVA is such that sufficient phase separation is achieved with moderate phase adhesion required for effective rubber toughening.

In this work, an attempt has been made to prepare super tough PLA and to broadly study the effect of the flexibility of EVA containing 50 wt (%) of VA on the mechanical properties of PLA. To understand the phase interaction between PLA and EVA, tensile data have been analyzed employing predictive models. In order to evaluate the state of dispersion of the elastomer EVA in the matrix SEM studies have been undertaken. The notched Izod impact behaviour has been correlated with blend morphology.

2. Experimental

2.1. Materials

Injection moulding grade PLA {(Ingeo 3052 D, D-isomer content = 4.2%, melt flow index 14 g/10 min, density = 1.24 g cm⁻³ (210 °C and 2.16 kg load))} was purchased from NatureWorks LLC,⁷ USA. EVA copolymer (VA content 50 wt%, density 1 g cm⁻³, T_g −29 °C) was procured from Lanxess Pvt. Ltd., Germany.

2.2. Blend preparation

Both PLA and EVA copolymers were vacuum-dried at 50 °C for 12 h before use. Binary blends of PLA/EVA at varying concentrations of EVA of 0–30 wt% (volume fraction of EVA Φ_d = 0–0.35) were melt blended using a co-rotating micro extruder Haake Minijet Lab at a screw rpm of 100. The processing temperature was maintained at 190 °C. The volume fraction Φ_d of the blending polymer was calculated according to eqn (1):³²

Φ_d = (W₁/ρ₁)/[(W₁/ρ₁) + (W₂/ρ₂)] (1)

where W is the weight fractions and ρ the density (g cm⁻³) of the constituents. Subscripts 1 and 2 denote the dispersed and the continuous phase, respectively. The PLA granules were also extruded under same conditions to ensure an identical shear and thermal history as that of the blends. The details of the blend formulations and the corresponding Φ_d values are presented in Table 1.

Table 1 Compositions and values of DSC crystallization parameters of PLA in PLA/EVA blends

PLA (wt%)	EVA (wt%)	Volume fraction (Φd)	Tg (°C)	Tcc (°C)	ΔHcc (J g^−1)	Tm (°C)	ΔHm (J g^−1)	Xc (%)
100	0	0	60.9	104.8	31.1	154.9	29.3	32
95	5	0.06	61.1	125.8	24.7	151.6	27.0	31
90	10	0.12	61.2	126.8	19.2	152.1	22.9	28
80	20	0.24	61.4	125.4	18.7	151.7	20.8	27
70	30	0.35	61.9	126.4	10.3	151.3	15.2	23

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2.3. Preparation of test specimens

The specimens for the evaluation of the mechanical properties were micro injection molded using a Haake Minijet-II micro injection molding machine. The barrel temperature was 200 °C. Injection pressure and mold temperature were 650 bar and 65 °C, respectively.

2.3.1. Mechanical testing. Tensile tests of dumb-bell shaped specimens were performed on a Zwick universal tester, Model Z010 (Germany), at a cross-head speed of 10 mm min⁻¹ and cross-head separation of 15 mm according to the ASTM D 638 (type 5) test procedure.³³ Izod impact strength was measured using notched specimens on a falling hammer type instrument, Model 504 plastic Impact (USA). A notch of 2 mm depth with an angle of 45° was made on the impact specimens following the ASTM D 256 specifications.³² At least five samples at each composition were tested and the average values are reported. All the tests were performed at an ambient temperature of 30 ± 2 °C.

2.3.2. Differential scanning calorimetry (DSC). The crystallization behaviour of PLA and that of the blends was studied using DSC using a TA Instrument, Model Q 200, in an atmosphere of liquid nitrogen. Samples were powder/flakes scraped from the injection molded tensile specimens and vacuum dried at 80 °C for 4 h. The samples were first heated from 30 °C to 200 °C at a heating rate of 5 °C min⁻¹ and detained at 200 °C for 5 min to eradicate the effect of the previous thermal history. The melt was then cooled to 30 °C and again heated to 200 °C at 5 °C min⁻¹ cooling and heating rates. Heat of fusion values (ΔH) were employed to evaluate the crystallinity (%) of PLA in the blends using eqn (2):

Degree of crystallinity, χ (%) = [(ΔH_m/ΔH_o)/W] × 100 (2)

where, ΔH_m is the enthalpy of the fusion of PLA in the blends, ΔH_o the enthalpy of fusion of the 100 (%) crystalline PLA, 93 J g⁻¹,¹⁶ and W the weight fraction of PLA in the sample.

2.3.3. Scanning electron microscopy (SEM). The phase morphology of the impact fractured samples was observed using a scanning electron microscope, Model EVO 50, at 20 kV to evaluate the dispersion of the discrete phase in PLA. Before scanning, sample surfaces were sputter coated with a thin layer of silver. The SEM micrographs were analyzed using the image analysis software, Image J, to calculate the diameter of the dispersed phase. The weight average particle size, d_w of EVA was calculated from a mean of minimum 200 particles using eqn (3):³⁴

d_w = ∑n_id_i²/∑n_id_i (3)

where n is the number of particles with diameter d. The interparticle spacing, i.e. the ligament thickness, τ was determined from d_w and Φ_d following eqn (4):³²

τ = d_w[(π/6Φ_d)^1/3 − 1] (4)

In eqn (4), EVA particles were assumed to be uniform spheres arranged in a cubic lattice.

3. Results and discussion

3.1. Degree of crystallinity

Fig. 1 shows the DSC heating scans of the samples and the corresponding parameters of PLA are depicted in Table 1. It is observed that the presence of EVA does not influence the glass transition temperature (T_g) of PLA. From the DSC measurements, the endothermic peak temperature and peak height give information about crystal size and size distributions. Moreover the number of peaks gives information about the different crystal structures (bimodal), and the normalized enthalpy gives information about the degree of crystallinity.³⁵ For the second heating, exotherm peaks appear which are a well-known feature of PLA crystallization during DSC measurements.

Fig. 1 DSC traces of PLA and PLA/EVA blends recorded during the second heating at 5 °C min⁻¹.

The heating curve of neat PLA shows very distinct exothermic peaks and double melting peaks for cold crystallization temperature (T_cc) and melting temperature (T_m) respectively. With the increase of EVA concentration the cold crystallization peak weakened and significantly shifted to higher temperatures which indicate that the cold crystallization of PLA becomes more difficult and less PLA can transform into a crystalline state. This can be explained on the basis that during the heating process EVA melts prior to PLA and possibly promotes chain mobility in the interface between PLA and EVA, and subsequently plays a plasticizing role in promoting the cold crystallization of PLA.¹⁶ Consequently, the crystallinity (%) of PLA decreases with an increase in EVA wt (%). However, the decrease in the crystallinity of PLA may be due to the physical presence of the EVA at increasing concentrations which disrupts the continuity of the PLA matrix. Moreover, this is possibly due to the enhanced phase adhesion between PLA and EVA, which prevents the migration of the PLA chain segments out from the EVA phase, and hence limits the crystallization of PLA.^10,16

Since the crystallinity plays an important role in the properties of PLA/EVA blends, this parameter will be considered in the subsequent sections in the analysis of the mechanical properties.

3.2. Mechanical properties

3.2.1. Tensile stress–strain curves. The tensile stress–strain curves of the PLA and PLA/EVA blends are shown in Fig. 2. Neat PLA fails once it passes the yield point. The stress–strain curve of EVA is that of a typical rubbery polymer without a distinct yield point up to 1300% elongation-at-break, as shown in the Fig. 2 inset. Neat PLA possesses a strong strain softening which is not stabilized by a strain hardening during breaking, and subsequently strain softening initiates strain localization causing local tri-axial stress concentration.⁸ Because of the absence of local strain delocalization, local tri-axial stresses induce void nucleation and crazes leading to the brittle failure behaviour of PLA.⁸ Incorporation of EVA in PLA causes a lower yield stress with a broadening of yield peak because of the release of tri-axial stress or enhanced delocalization of strain introduced by the flexible EVA copolymer.^6,8 Elongation-at-break was enhanced with increased Φ_d and at the maximum Φ_d the parameter was 225%. PLA/EVA blends exhibited strain softening followed by neck formation and strain hardening which indicates the ductility of the blends.

Fig. 2 Variation of stress and strain with Φ_d in PLA/EVA blends. Inset: variation of stress and strain for pure EVA.

The lowering of the yield stress of the PLA/EVA blends with an accompanying elongation increase compared with that of pure PLA may be ascribed to the stress concentration effect of the rubbery EVA co-polymer. The area under the stress–strain curve of PLA increased significantly upon incorporation of EVA at the entire range of Φ_d which indicated that the toughness of the blends is higher than that of PLA. This enhanced toughness is attributed to increased energy absorption and enhanced ductility brought about by the elastomer EVA as also reported in other rubber modified polymers.^32,34,36

The tensile properties e.g. the tensile modulus, tensile strength and elongation-at-break were evaluated from the stress–strain curves and are presented in Fig. 3–5 as the ratio of the property of the blends (subscript b) to that of the neat PLA matrix (subscript m) as a function of the volume fraction, Φ_d.

3.2.2. Tensile modulus. The tensile modulus variations with crystallinity of PLA in the PLA/EVA blends are presented in Table 2. The modulus value decreases substantially with a decrease in crystallinity (%) which indicates that PLA is significantly softened by the presence of the flexible EVA co-polymer. The softening is caused by a decrease in crystallinity as well as the flexibility of EVA.

Table 2 Tensile property results of the PLA/EVA blends

Φd	Xc (%)	Tensile modulus (MPa)	Tensile strength (MPa)	Elongation-at-break (%)
0	32	439 ± 16.56	75 ± 0.51	16 ± 1.25
0.06	31	416 ± 26.21	65 ± 2.92	21 ± 1.76
0.12	28	327 ± 27.3	60 ± 0.55	30 ± 1.99
0.24	27	255 ± 39.8	47 ± 2.16	146 ± 10.76
0.35	23	205 ± 12.37	35 ± 1.09	164 ± 5.89

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Fig. 3 presents variations of the relative tensile modulus, E_b/E_m, of the PLA/EVA blends as a function of Φ_d. The relative modulus shows a substantial decrease with an increase in Φ_d implying that PLA is significantly softened by the EVA copolymer. The blend structure was evaluated by comparing experimental data with simple theoretical predictions according to the “rule of mixture”^32,37 eqn (5), as well as the “foam model”^32,34 eqn (6):

(E_b/E_m) = [E_d/E_m − 1]Φ_d + 1 (5)

(E_b/E_m) = [1 − Φ_d^2/3] (6)

here, E_b is the modulus of the blends, E_m the modulus of the matrix polymer PLA, and E_d the modulus value for EVA. In these calculations, the modulus values of the PLA (E_m = 439 MPa) and the blends (E_b) were determined from the initial slopes of the stress vs. strain curves. The modulus value E_d of the EVA copolymer was 0.19 MPa, which was a 1% secant modulus at 10 mm min⁻¹ cross-head speed. In the foam model, since the modulus ratio E_d/E_m was negligible, the dispersed phase was considered a non-interacting void or pore.

Fig. 3 Plot of variations of E_b/E_m of PLA/EVA blends against Φ_d and their predictive behaviour according to the foam model and rule of mixture. Inset: dependence of normalized tensile modulus, [(E_b/X_b)/(E_m/X_m)], of PLA/EVA blends against Φ_d.

The relative modulus data were in close agreement with the rule of mixture at Φ_d = 0.06, however the data at higher Φ_d were lower than the theory and agreed well with the foam model, (Fig. 3). It indicates that at Φ_d = 0.06 some kind of phase interaction may be operative, whereas at higher concentrations the elastomeric phase may function as a diluting or flexibilizing agent.

In order to evaluate the flexibilizing effect of EVA on the matrix, the effect of the crystallinity of PLA was eliminated by normalizing the relative modulus data by the crystallinity of PLA in the blends (X_b) and the matrix (X_m) respectively. Fig. 3 inset, presents the plot of the normalized relative tensile modulus, [(E_b/X_b)/(E_m/X_m)] against Φ_d. The value decreased drastically with an increase in Φ_d and at the highest Φ_d of 0.35 the decrease was 0.7 fold. This indicates that the elastomer is indeed a flexibilizing dispersed phase facilitating nonspecific phase interaction with the continuous PLA phase. Similar observations were reported in PP/EVA blends also.³⁸

3.2.3. Tensile strength. The variation of tensile strength with the crystallinity of PLA in the PLA/EVA blends is given in Table 2. The tensile strength of the PLA/EVA blends decreases continuously with the decrease in crystallinity of PLA, which may be due to interference in the nucleation and growth of PLA crystals by the presence of the flexible blending polymer EVA.

Fig. 4 presents variations in the relative tensile strength (ratio of the relative tensile strength of PLA/EVA blends to that of the PLA, σ_b/σ_m) vs. Φ_d. The tensile strength of PLA continuously decreased with increasing Φ_d. This weakening of blend structure may be attributed to the softening effect of EVA copolymer which subsequently decreases the effective cross sectional area of the continuous PLA phase. Similar results were reported in other toughened polymer systems also where elastomer is the discrete phase.^32,37 The decrease in the tensile strength is the consequence of a decrease in the crystallinity, and a similar trend was observed in the tensile modulus also (Table 2). The relative tensile strength data were analyzed following simple predictive models, the “Porosity model”^32,37,39 eqn (7), and the “Nicolais Narkis model”^34,37,40 eqn (8), to evaluate the discontinuity or weakness introduced by the dispersed phase EVA:

σ_b/σ_m = [exp(−αΦ_d)] (7)

σ_b/σ_m = [1 − KΦ_d^2/3] (8)

Fig. 4 Plot of variations of σ_b/σ_m of PLA/EVA blends (▪), porosity model (—), [eqn (7)] with α = 2.06, and Nicolais Narkis model (…), [eqn (8)] with K = 0.91, against Φ_d. Inset: variations of normalized relative tensile strength [(σ_b/X_b)/(σ_m/X_m)] of PLA/EVA blends against Φ_d.

Similar models have been employed in other two phase systems of polymer blends/composites.³⁷ These models assume no adhesion between the two phases³⁷ and tensile strength depends on either the area fraction or the volume fraction of the dispersed phase i.e. Φ_d^2/3 or Φ_d, respectively. Detailed descriptions about the theories and the significance of the parameters are available elsewhere.³⁹ In eqn (8), a low value of the interaction parameter K, also known as the weightage factor, denotes a higher phase interaction. Although K = 1.21 denotes poor adhesion, K = 0 describes significant adhesion so that the strength of the polymer matrix does not decrease. Similarly, the value of parameter α describes stress concentration, where the higher the value of α the higher the stress concentration.^32,37,41,42

Table 3 depicts the values of α and K for each Φ_d value obtained from comparisons of experimental tensile strength data with eqn (7) and (8). According to the Nicolais Narkis model⁴⁰ eqn (8), the values of K were less than one with a mean value of 0.91, which denotes a good degree of adhesion and subsequently a smaller extent of weakness in the blend structure.^43,44 Gupta and Purwar⁴⁵ reported similar results earlier in PP/SEBS/HDPE and PP/SEBS/PS blends. The porosity model, eqn (7) and Table 3, exhibit a significant degree of stress concentration with the value of a = 2.06. This value is close to the value of 2.04 reported in iPP/CSM blend.⁴⁶

Table 3 Values of the stress concentration factor α, eqn (7), and adhesion parameter K, eqn (8), in PLA/EVA blends

Φd	K	α
0	—	—
0.06	0.77	2.32
0.12	0.82	1.86
0.24	0.97	1.93
0.35	1.07	2.16
Mean value	0.91	2.06

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The variations of the normalized relative tensile strength data, (σ_b/X_b)/(σ_m/X_m), against Φ_d are presented in the Fig. 4 inset, to evaluate the effects of the phase interaction and flexibility of EVA on PLA. Here X_b and X_m represent the crystallinity of PLA in the PLA/EVA blends and that of neat PLA, respectively. The normalized tensile strength of PLA decreased continuously on blending with the EVA co-polymer (increasing Φ_d). The data were less than unity at Φ_d = 0.06–0.35. This weakening of the blend structure may be because of the decrease in the effective load bearing cross-sectional area of the matrix polymer in the presence of the co-polymer, similar to other elastomer-modified systems.⁴⁷ The decrease in tensile strength is also aided by the flexibility of EVA. Thus, there is no apparent phase adhesion observed, the flexibility of the blending polymer predominates and consequently the tensile strength decreases in the presence of the EVA co-polymer.

3.2.4. Elongation-at-break. The variation of elongation-at-break with the crystallinity (%) of PLA/EVA blends is presented in Table 2. The increase in elongation-at-break with the decrease in crystallinity may be attributed to an enhanced amorphization of PLA along with the flexibility of the blending polymer which eventually make the PLA more ductile.

Fig. 5 exhibits the variation of the relative elongation-at-break (ε_b/ε_m) against Φ_d. The ε_b/ε_m of the PLA/EVA blends increased with Φ_d. The increase is attributed to the enhanced ductility of the system due to an increase in the amorphization of PLA along with the flexibility of EVA. The continuous increase in the value of ε_b/ε_m with increasing Φ_d indicates a significant increase in the ductility of PLA caused by the EVA copolymer.

Fig. 5 Variation of relative elongation-at-break (ε_b/ε_m) of PLA/EVA blends against increasing Φ_d. Inset: variation of relative normalized elongation-at-break (ε_b/X_b)/(ε_m/X_m) of PLA/EVA blends with increasing Φ_d.

Initially, the data increase to a small extent up to Φ_d = 0.12; however, beyond Φ_d = 0.12 the increment was quite significant, at Φ_d = 0.24 and Φ_d = 0.35, the values increased by 9.13 times and 10.25 times respectively. Tensile moduli data also showed significant matrix softening in the presence of the EVA copolymer.

It may be noted here that the enhancement of the ductility of PLA may be due to flexibility of the discrete phase and enhanced amorphization of PLA. The modulus decrease of PLA in the presence of EVA also indicated the matrix softening. The softening of the PLA matrix indicates toughening of the polymer under the impact mode of load application. Moreover, the toughened polymer will consume additional energy to break and the fracture mechanism may also change similar to a ductile material.

To examine the effect of the flexibility of the EVA co-polymer on the enhancement of the ductility of PLA, the effect of the crystallinity of PLA was eliminated. Fig. 5 inset presents the relative normalized elongation-at-break, (ε_b/X_b)/(ε_m/X_m), against Φ_d. The data continuously increased in the entire range of Φ_d confirming an enhanced ductility due to matrix softening in the presence of the elastomer phase. Initially the data increased by a small extent up to Φ_d = 0.12, however, beyond Φ_d = 0.12, the increase in the elongation was very significant.

3.2.5. Impact strength. The notched Izod impact strength of the PLA/EVA blends was enhanced with the decrease in crystallinity (%) as shown in Table 4. Initially, up to Φ_d = 0.12, the data increased to a small extent, however, beyond Φ_d = 0.12, a sharp increase in the impact strength was observed. The notched Izod impact strength of the blends is low when the EVA percent is less than 10 wt% (Φ_d = 0.12) which may be due to the fact that small EVA particles are difficult to cavitate under the impact conditions. The increase in the impact values with a decrease in crystallinity may be due to the combined effect of the enhanced amorphization of PLA and flexibility of EVA.³²

Table 4 Values of impact strength (I_b), domain size (d_w), and interparticle distance (τ), of PLA/EVA blends^a

Φd	dw (μm)	τ (μm)	Ib (J m^−1)
a Note: values in the parentheses indicate impact strength in kJ m^−2.
0	—	—	17.1 (2.27)
0.06	0.51	0.54	18.1 (2.34)
0.12	0.55	0.35	25.5 (3.39)
0.24	1.10	0.33	136 (18.19)
0.35	1.91	0.27	403 (53.73)

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To eliminate the effect of crystallinity on the variation of impact strength, the relative Izod impact strength value of the PLA/EVA blends were normalized and are presented in Fig. 6 against Φ_d. Initially, up to Φ_d = 0.12, the normalized impact data increased to a small extent, however, beyond Φ_d = 0.12, the data increased drastically. The normalized impact strength at the maximum Φ_d is ∼33 times that of PLA, making the blend super tough, which indicates that the elastomer EVA toughened PLA substantially.⁷ This increase is related to the enhanced ductility of the blend due to increased flexibility caused by EVA which creates an increased extent of shear yielding. The variation of the impact strength against Φ_d data exhibits behavior similar to that of the elongation data implying an enhanced ductility in the system.^36,48,49

Fig. 6 Variation of normalized impact strength (I_b/X_b)/(I_m/X_m), of PLA/EVA blends against Φ_d.

3.3. Fracture surface morphology and mechanism of toughening

The cross-section images of the impact fractured surface morphology of the PLA/EVA blends are shown in Fig. 7a–e. PLA exhibits sharp ridges characteristic of a brittle fracture (Fig. 7a), without indicating shear yielding. The blends exhibited a two phase morphology e.g. a sea-island morphology, where EVA is the dispersed phase in the PLA matrix. Furthermore, in the sea-island morphologies EVA exhibits varying particle diameters with an increase in Φ_d. At low concentration the EVA droplets are small in size and mostly spherical shaped with a small percentage of them being slightly elongated, (Fig. 7b and c). With a further increase in Φ_d the droplets are mostly elongated and droplet size is enhanced as well. The elongated shape of the dispersed phase with increasing Φ_d can be explained on the basis that during the injection moulding process of the blends at a high shear rate, the softer elastomeric phase shows elongational flow into the relatively harder PLA phase. Due to the enhanced ease of deformation of the elongated elastomer particles into the neighboring stressed area, the shear yield stress of the matrix decreases.^50,51

Fig. 7 Scanning electron micrograph of PLA/EVA blends at 2000× magnification at varying Φ_d: (a) 0, (b) 0.06, (c) 0.12, (d) 0.24 and (e) 0.35. (f) Contains photographs of notched Izod impact fractured samples. Samples Φ_d = 0, 0.06 and 0.12, show complete breakage while Φ_d = 0.24 and Φ_d = 0.35 show partial breakage with extensive stress whitening.

The dispersed phase particle size increased with Φ_d, (Fig. 7b–d), which may be attributed to the increasing dynamic coalescence,^34,47 of the elastomer particles facilitated by the lower viscosity of PLA at the processing temperature. The blends show signs of elongation and tearing of the elastomer phase over a larger area, the extent of which increases with Φ_d, (Fig. 7b–e). The images of Φ_d = 0, 0.06 and 0.12 (Fig. 7f) indicated a brittle failure, since complete breakage takes place in these impact tested samples. Moreover, the samples Φ_d = 0.24 and 0.35 exhibited partial breakage of the impact tested sample, an indication of improved toughness and ductility, which can be a result of various fracture mechanisms such as crazing/micro-cracks, fibrillation, shear yielding etc.^52,53 Therefore, the improved toughness and high impact strength result from the special combination of energy absorption mechanisms that arrested the crack tip. A significant extent of whitening was observed in the SEM micrographs of the blends which was enhanced with an increase in Φ_d which may be due to the yielding of the elastomer phase at the interphase region as well as in the droplets which also dissipate energy. The stress whitening around the fractured surfaces for Φ_d = 0.24 and 0.35 illustrated plastic deformation of the matrix PLA indicating a shear yielding reported in similar other works.^6,52

Because of the low crack initiation as well as the low crack propagation energy PLA possesses a low Izod impact strength.⁵² Thus, the polymer is considered as brittle as in the case of the type-II polymer according to Wu’s definition.^37,54 Toughening of this type of polymer blends has been shown to be due to the matrix yielding introduced as a result of the elastomer’s interactions with the surrounding stress field at a critical value of ligament thickness (interparticle distance, τ). The parameter τ was calculated (Table 4) from the weight average particle diameter, d_w, and volume fraction, Φ_d (eqn (4)). The τ value ranged from 0.54 μm to 0.27 μm, as Φ_d varied from 0.06 to 0.35.

PLA is a semicrystalline polymer and the crystallinity decreases with the increase in EVA concentration. The crystallinity of PLA and ligament thickness, τ, will determine the impact strength of the PLA/EVA blends. The variations of normalized relative impact strength (I_b/X_b)/(I_m/X_m) against τ is shown in Fig. 8. As shown earlier, at maximum Φ_d = 0.35, the normalized impact strength was enhanced to ∼33 times (53.73 kJ m⁻²) that of PLA. In the studied range of EVA concentrations, the result indicates that the impact strength of the blends increases with a decrease in ligament thickness, τ. At τ = 0.33 and τ = 0.27 the impact strength increased to ∼8 times and ∼33 times, respectively. This increase may be due to a combination of fibrillation, crazing and extensive shear yielding caused by the EVA co-polymer. The flexibility of EVA, enhanced amorphization of PLA and creation of stress concentration may be assigned to the crazing, fibrillation and extensive shear yielding which ultimately led to the formation of the super tough PLA.

Fig. 8 Variation of normalized impact strength (I_b/X_b)/(I_m/X_m) of PLA/EVA blends against τ.

4. Conclusions

Incorporation of the EVA co-polymer into PLA decreases the crystallinity and substantially enhances its flexibility. The tensile modulus and strength decrease while toughness and ductility increase significantly with an increase in EVA content. The normalized tensile modulus values decreased significantly with an increase in Φ_d and at the highest Φ_d of 0.35 it decreased to 0.5 times the original, demonstrating that the elastomer is indeed a flexibilizing dispersed phase facilitating nonspecific phase interaction with the continuous PLA phase. The normalized tensile strength of PLA decreases continuously on blending with the EVA co-polymer (increasing Φ_d), indicating a decrease in the effective load bearing cross-sectional area of the matrix polymer. The normalized relative notched Izod impact strength was enhanced with increasing Φ_d, at maximum Φ_d = 0.35, the value was enhanced ∼33 times (53.73 kJ mm⁻²) making the blend super tough. A two-phase morphology was observed using SEM, where the elastomer exists as the dispersed phase. Spherical shaped EVA particles are uniformly dispersed in the PLA matrix. The droplets become elongated and the size increased at higher EVA contents. Shear yielding of elastomer in the elastomer phase as well as at the PLA–EVA inter-phase led to stress whitening.

Acknowledgements

The authors would like to acknowledge Indian Institute of Technology Delhi and University Grant Commission for providing research facilities and financial assistance to one of the authors (Rajendra Kumar).

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