Nanofibrillar Green Composites of Polylactide/Polyhydroxyalkanoate Produced in Situ Due to Shear Induced Crystallization

This study addresses the new concept of in situ inducing fibrillar morphology (micro or nanofibrils) of a minority component based on the simultaneous occurrence of orientation and shear induced crystallization of polymer fibers directly at the stage of extrusion in a single step. This possibility is demonstrated by using two entirely bio-sourced polymers: polylactide (PLA) and polyhydroxyalkanoate (PHA) as components. The shear induced crystallization allowed crystallization of PHA nanofibers immediately after applying high shear rate and elongation strain, avoiding subsequent cooling to initiate crystallization. Shearing of PHA increased non-isothermal crystallization temperature by 50 °C and decreased the temperature range in which the transition from a molten state to a crystallized one occurs by 17 °C. SEM observations demonstrate the successful transformation of the dispersed PHA phase into nanofibrils with diameters of nearly 200 nm. The transition from the droplets of PHA to fibers causes the brittle-to-ductile transition of the PLA matrix at a low concentration of PHA and contributes to the simultaneous increase of its rigidity and strength.


Introduction
Recently, there has been a lot of attention on the design of so-called "ecocomposites" or "green composites" that are alternative to synthetic polymer composites. As the performance of green polymer composites depends on the physical, melt rheological characteristics of the components, their morphology, and the interfacial interaction strength, the progress is related both to the search for new pairs of matrices and fillers and to development of new methods of production of polymer composites.
The most widely applied methods of production of green composites are those based on the blending of polymers and ready-made inorganic or natural-organic fibers or particles [1][2][3][4] as well as more complex multi-stage methods addressing in situ formation of varied morphology of the filler: discrete or entangled fibers, shish-kebab structures (when the crystallization of the matrix polymer occurs on the surface of the formed fibrils), etc. [5][6][7][8]. In the last cases, as a rule, a biopolymer immiscible with the biopolymer matrix is used as the minor phase. The general approach to the formation of polymer fibers in situ within another polymer matrix is the preparation of a polymeric blend at the first stage when one of the polymers is dispersed in the form of droplets (micro-or nano) that are later transformed into fibers under the action of forces. As a rule, hot/cold drawing, rolling or electrospinning is applied here. The resultant composite is characterized by the orientation of both

Scanning Electronic Microscopy (SEM)
The morphology of samples, cryogenically fractured along the extrusion direction and coated with gold, was investigated with a JEOL JSM-5500 LV scanning electron microscope (Tokyo, Japan). Further observation was done on the samples which were ultramicrotomed with a diamond knife (Diatome, Switzerland) at −130 • C. Gatan MT200 (Microtest Tensile Stage, Suffolk, UK), connected to the microscope, facilitated in situ observations of the tensile test according to ASTM D638. Specimens were coated with carbon and deformed with a rate of 0.2 mm/min following the procedure described previously [11].

Rheo-Optical Measurements: Shear-Induced Crystallization Test
Shear-induced crystallization of the PHA was scrutinized in the Linkam CSS450 (Optical Shearing System, Surrey, UK) as per [11]. Films of PHA were positioned between two heated glass stages undergoing subsequent treatment under a nitrogen flow: (1) heating to 60 • C above the melting point at (30 • C /min), (2) holding for 5 min to eliminate possible orientation of chains, (3) cooling down with simultaneous shearing at rates 100-1000 s −1 .

Rheological Measurements
The rheological behavior of materials was examined using a strain-controlled rotational rheometer (ARES LS2, TA Instruments). Uniaxial extension tests of molten samples were performed using extensional viscosity fixture (EVF, TA Instruments) attached to the ARES rheometer (New Castle, DE, USA). The 18 × 10 × 0.7 mm 3 rectangular specimens were prepared by hot compression molding at 200 • C in the standard mold, provided by TA Instruments with the EVF, or were cut out from extruded tapes. The specimen was uniaxially extended at 100 • C with a constant Hencky strain rate,έ, of 1.0, 2.0, or 4.0 s −1 . The tensile stress growth coefficient, η + E (t,έ), as a function of time, t, at a given Hencky strain rate was measured. The tests for each material were repeated at least eight times, and the results were averaged. Steady-state shear measurements were executed to investigate the viscosity-shear rate relationship. Samples were tested using cone-plate geometry with a cone angle of 0.1 rad and a gap distance of 0.046 mm. Disk-shaped samples were cut out from 1 mm thick hot-pressed films. The shear tests were performed at temperatures 135, 155, and 180 • C and shear rates, γ, ranging from 1 to 300 s −1 . Figure 1 shows the effect of the shear rate γ on the conversion degree of PHA. It is seen that as the shear rate increases, the crystallization temperature, T c , shifts towards higher temperature and the transition from a molten state to a crystallized one proceeds in a narrower temperature range, ∆T. In particular, an increase in γ up to 300 s −1 results in a shift of T c toward higher temperatures up to 137 • C and a decrease in ∆T from 38 to 24 • C. The maximum effect is achieved at γ = 850 s −1 . In this case, T c is 156 • C and ∆T = 17 • C. The shift in T c was smaller for the PHA sheared at constant temperature and subsequently cooled down ( Figure 2). The observed effect is due to the fact that the implementation of simultaneous shear and cooling leads to a significant reduction in the time gap between cessation of shearing and the onset of crystallization. The shorter this time gap, the less possibility for the crystal nucleus to re-melt at temperatures above non-isothermal crystallization temperature and for polymer chains to recoil and relax. An increase in the shearing rate promotes an increase in the concentration of nuclei and accelerates the crystallization process. In that way, in situ generated PHA fibers would crystallize directly under a high shear rate without the necessity for cooling below the temperature of their non-isothermal crystallization. The relaxation process minimized in this way would cause the retention of the achieved elongation degree of droplet inclusions into fibers. to 137 °С and a decrease in ΔT from 38 to 24 °C. The maximum effect is achieved at γ = 850 s -1 . In this case, Тс is 156 °С and ΔT = 17 °C. The shift in Tс was smaller for the PHA sheared at constant temperature and subsequently cooled down ( Figure 2). The observed effect is due to the fact that the implementation of simultaneous shear and cooling leads to a significant reduction in the time gap between cessation of shearing and the onset of crystallization. The shorter this time gap, the less possibility for the crystal nucleus to re-melt at temperatures above non-isothermal crystallization temperature and for polymer chains to recoil and relax. An increase in the shearing rate promotes an increase in the concentration of nuclei and accelerates the crystallization process. In that way, in situ generated PHA fibers would crystallize directly under a high shear rate without the necessity for cooling below the temperature of their non-isothermal crystallization. The relaxation process minimized in this way would cause the retention of the achieved elongation degree of droplet inclusions into fibers.    to 137 °С and a decrease in ΔT from 38 to 24 °C. The maximum effect is achieved at γ = 850 s -1 . In this case, Тс is 156 °С and ΔT = 17 °C. The shift in Tс was smaller for the PHA sheared at constant temperature and subsequently cooled down ( Figure 2). The observed effect is due to the fact that the implementation of simultaneous shear and cooling leads to a significant reduction in the time gap between cessation of shearing and the onset of crystallization. The shorter this time gap, the less possibility for the crystal nucleus to re-melt at temperatures above non-isothermal crystallization temperature and for polymer chains to recoil and relax. An increase in the shearing rate promotes an increase in the concentration of nuclei and accelerates the crystallization process. In that way, in situ generated PHA fibers would crystallize directly under a high shear rate without the necessity for cooling below the temperature of their non-isothermal crystallization. The relaxation process minimized in this way would cause the retention of the achieved elongation degree of droplet inclusions into fibers.      Figure 3 shows the morphologies of the cryogenically fractured surfaces of the PLA/PHA blend and in situ generated PLA/PHA composites. The PLA/PHA blend presents typical matrix-droplet morphology. PHA spherical particles and dark holes left by them during fracture are observed on the surface of the blend (Figure 3a). The surface of these particles is smooth with clear borders, suggesting poor compatibility and weak interfacial adhesion between the PLA and PHA phases. Meanwhile, the low concentration of PHA (3 wt.%) makes their dispersion in PLA matrix well-uniform and causes the formation of nano-sized spherical particles of PHA. Most PHA dispersed phases are in the range of 50-450 nm and the averaged size is about 275 nm. Intensive blending (screw speed was 120 rpm) also contributes to the deformation and fracture of the PHA phase and minimizes agglomeration of the shear-split PHA phase. Compared to the PLA/PHA blend, SEM micrographs of the in situ generated PLA/PHA composite reveal the presence of fibril-matrix morphology ( Figure 3b). Almost all the PHA dispersed phase particles are elongated into fibrils due to the presence of a shear force field within the area between the extruder screw and the extruder walls. Detailed mixing mechanisms of immiscible polymers are provided elsewhere [23]. It is already well established that when dispersed minor polymer droplets, immersed in the immiscible matrix, are subjected to high shear rates, they tend to deform and extend into long thin threads. As a result, particle radius decreases while the capillary instabilities at the interface increase. When the capillary number approaches to the critical value, the elongated nanofibrils tend to disintegrate and break up. The most important mechanisms, depending on the viscosity ratio and the degree of elongation, could be the growth of uniform Rayleigh disturbances in the middle part of the thread, end-pinching at the ends, and necking. However, rapid solidification caused by shear induced crystallization prohibits these time-dependent transformations. Thus, continuously elongated in the melt, PHA nanofibrils are stabilized and fixed by crystallization before they may undergo any retraction or disintegration. This is reasonable because the viscosity ratio between the PLA/PHA blend and neat PLA at low temperatures is, according to [24], in the optimum range, which is (0.1:10), and remains within this range despite the shear thinning which occurs at high shear rates, that is the hydrodynamic force overcomes the cohesive strength of the dispersed PHA droplets (shown in the Supplementary Material, Figure S1). It is also well known that fibrous droplets could eventually take the form of spherical ones or break up into small particles by interfacial tension, but rapid solidification prohibits these transformations and results in fibrous dispersion.

Results and Discussion
The final nanofibers are 70-130 nm in diameter and form the physical network structure. The formation of the physical network of PHA in the PLA matrix is possible during compounding and not during extensional flow in the slit die zone. Apparently, the slit capillary serves as a dumping for the flow and to increase the residence time in the extruder and facilitate the formation of fibril entanglements. The ratio of the diameters of the fibrils to the droplets shows that the aspect ratio of PHA fibrils is at least 100. Since the exact length of the fibers is unclear, it can be assumed that the actual aspect ratios of PHA fibrils can be even larger because of weak interfacial adhesion between the PLA and PHA phases and a uniform distribution of PHA nanoparticles, which increase the probability of them coalescing during shearing, and thereby facilitate formation of very long fibrils. Compared to the PLA/PHA blend, SEM micrographs of the in situ generated PLA/PHA composite reveal the presence of fibril-matrix morphology (Figure 3b). Almost all the PHA dispersed phase particles are elongated into fibrils due to the presence of a shear force field within the area between the extruder screw and the extruder walls. Detailed mixing mechanisms of immiscible polymers are provided elsewhere [23]. It is already well established that when dispersed minor polymer droplets, immersed in the immiscible matrix, are subjected to high shear rates, they tend to deform and extend into long thin threads. As a result, particle radius decreases while the capillary instabilities at the interface increase. When the capillary number approaches to the critical value, the elongated nanofibrils tend to disintegrate and break up. The most important mechanisms, depending on the viscosity ratio and the degree of elongation, could be the growth of uniform Rayleigh disturbances in the middle part of the thread, end-pinching at the ends, and necking. However, rapid solidification caused by shear induced crystallization prohibits these time-dependent transformations. Thus, continuously elongated in the melt, PHA nanofibrils are stabilized and fixed by crystallization before they may undergo any retraction or disintegration. This is reasonable because the viscosity ratio between the PLA/PHA blend and neat PLA at low temperatures is, according to [24], in the optimum range, which is (0.1:10), and remains within this range despite the shear thinning which occurs at high shear rates, that is the hydrodynamic force overcomes the cohesive strength of the dispersed PHA droplets (shown in the Supplementary Material, Figure S1). It is also well known that fibrous droplets could eventually take the form of spherical ones or break up into small particles by interfacial tension, but rapid solidification prohibits these transformations and results in fibrous dispersion.
The final nanofibers are 70-130 nm in diameter and form the physical network structure. The formation of the physical network of PHA in the PLA matrix is possible during compounding and not during extensional flow in the slit die zone. Apparently, the slit capillary serves as a dumping for the flow and to increase the residence time in the extruder and facilitate the formation of fibril entanglements. The ratio of the diameters of the fibrils to the droplets shows that the aspect ratio of PHA fibrils is at least 100. Since the exact length of the fibers is unclear, it can be assumed that the actual aspect ratios of PHA fibrils can be even larger because of weak interfacial adhesion between the PLA and PHA phases and a uniform distribution of PHA nanoparticles, which increase the probability of them coalescing during shearing, and thereby facilitate formation of very long fibrils. The dispersed PHA fibrils show tight contact with the PLA matrix, leaving no distinct voids between the elongated fibers and the matrix. This fact indicates a good adhesion between the fibrils and the matrix.
The SEM micrograph analysis of the PLA/PHA composite in longitudinal and transverse sections shows the uniform distribution of the microfibrils, and nearly the same number of revealed fibrils in both sections indicates the possibility of forming a physical network structure (Figure 4).  According to [25], 2.5 wt.% of the minor phase corresponds to the gelation point at which physical network structure forms. To verify this assumption, the viscoelastic behavior of the PLA/PHA blend and in situ generated PLA/PHA composites were examined. Figure 5 presents a time-dependence of the tensile stress growth coefficient, η E + (t, έ), for molten PLA, PLA/PHA blend and composite recorded during uniaxial tension at various Hencky strain rates, έ. For neat PLA and the PLA/PHA blend, η E + (t, έ) gradually increases with time, t, but it is nearly independent of έ. The monotonic stress growth reflects a predominantly plastic flow. The tensile stress growth coefficient, η E + (t, έ), then reaches the end of plateau when the chains become highly extended. At and beyond this plateau, chains could start to slide past one another at the entanglement points causing yielding of the entanglement network and the appearance of a strain softening effect. As the number of load-bearing entanglements decreases, the tensile stress growth coefficient, η E + (t, έ), is reduced as shown in Figure 5. According to [25], 2.5 wt.% of the minor phase corresponds to the gelation point at which physical network structure forms. To verify this assumption, the viscoelastic behavior of the PLA/PHA blend and in situ generated PLA/PHA composites were examined. Figure 5 presents a time-dependence of the tensile stress growth coefficient, η + E (t,έ), for molten PLA, PLA/PHA blend and composite recorded during uniaxial tension at various Hencky strain rates,έ. For neat PLA and the PLA/PHA blend, η + E (t,έ) gradually increases with time, t, but it is nearly independent ofέ. The monotonic stress growth reflects a predominantly plastic flow. The tensile stress growth coefficient, η + E (t,έ), then reaches the end of plateau when the chains become highly extended. At and beyond this plateau, chains could start to slide past one another at the entanglement points causing yielding of the entanglement network and the appearance of a strain softening effect. As the number of load-bearing entanglements decreases, the tensile stress growth coefficient, η + E (t,έ), is reduced as shown in Figure 5.  According to [25], 2.5 wt.% of the minor phase corresponds to the gelation point at which physical network structure forms. To verify this assumption, the viscoelastic behavior of the PLA/PHA blend and in situ generated PLA/PHA composites were examined. Figure 5 presents a time-dependence of the tensile stress growth coefficient, η E + (t, έ), for molten PLA, PLA/PHA blend and composite recorded during uniaxial tension at various Hencky strain rates, έ. For neat PLA and the PLA/PHA blend, η E + (t, έ) gradually increases with time, t, but it is nearly independent of έ. The monotonic stress growth reflects a predominantly plastic flow. The tensile stress growth coefficient, η E + (t, έ), then reaches the end of plateau when the chains become highly extended. At and beyond this plateau, chains could start to slide past one another at the entanglement points causing yielding of the entanglement network and the appearance of a strain softening effect. As the number of load-bearing entanglements decreases, the tensile stress growth coefficient, η E + (t, έ), is reduced as shown in Figure 5. In contrast to neat PLA and the PLA/PHA blend, exhibiting no strain hardening itself, the in situ generated PLA/PHA composite containing a PHA fiber network structure demonstrates a significant increase of η + E (t,έ) above the linear viscoelastic region during uniaxial extensional deformation. This is because the PHA fiber network, containing physical entanglements between the fibers, is deformed during uniaxial extension causing such an upward deviation of η + E (t,έ) for the PLA/PHA composite, known as strain hardening.
An important condition of effective conversion of the droplets of the minor polymer into fibers is the immiscibility of polymer constituents. According to the DSC analysis of the PLA/PHA extruded nanocomposite, melting of 3 wt.% of PHA is evidenced by a small endothermic peak at around 170 • C (shown in the Supplementary Material, Figure S2). The glass transition of the PLA matrix is seen at around 60 • C with an enthalpy relaxation signal caused by PLA orientation. The separate melting peak of PHA crystals proves that PHA is not miscible with PLA.
A similar conclusion about the immiscibility of PLA and PHA can be drawn based on E // data from DMTA measurements of PLA/PHA extruded tapes (shown in the Supplementary Material, Figure S3). The E // peak at the PLA glass transition is not shifted for the PLA/PHA nanocomposite and the glass transition of PHA is not discernible at -17 • C because of the low concentration of PHA (3 wt.%).
The morphology of the dispersed phase plays an important role in achieving improved mechanical properties of polymer blends or composites. In the case of the PLA/PHA blend, the spherical PHA particles with a diameter ranging between 50 and 450 nm were found to be distributed in the PLA matrix, which indicates a completely immiscible blend. Additionally, the PHA particles exhibited a uniform dispersion in the PLA matrix. The morphology of immiscible phases is also observed in the case of the PLA/PHA composite, with fibers with a diameter of 70-130 nm and an aspect ratio of not less than 100. As a result, for the PLA/PHA blend, the Young modulus, stress at break, and elongation at break exceeded the related parameters of neat PLA by 5%, 28%, and 6% (2.14 GPa, 55.0 MPa, and 7.4%), respectively. In the case of the in situ generated PLA/PHA composite chatacterized by fibrillar morphology, the Young modulus, stress at break, and elongation at break exceeded the related parameters of neat PLA by 15%, 30%, and 400% (2.35 GPa, 56.0 MPa, and 28.1%), respectively. This could be explained by the presence of PHA nanoparticles in the PLA matrix, which restrict the mobility of PLA chains and thus improve the stiffness of the PLA matrix. Moreover, the absence of signs of aggregation of the PHA nanoparticles in the PLA matrix i.e., the homogeneous dispersion of the nanoparticles, leads to a more regular and stronger structure of the blend, thereby improving the interfacial integrity of PLA and PHA to better transfer the stress under external force [26]. However, the much larger reinforcing effect, in the case of the PLA/PHA composite, is caused by the more extensive load-bearing capacity of the long PHA nanofibrils forming the network as well as the existence of larger interfacial areas between the nanofibers and the matrix. The main contribution to a significant increase in elastic and strength properties is associated with the formation of the PHA nanofibril network. This clearly demonstrates the dependences of true stress as a function of the Hencky strain after subtracting the data for neat PLA from respective stress-elongation curves for the nanocomposite ( Figure S4 in Supplementary Materials). It is revealed that the nanofibril's response is weak for small strains. However, exceeding the strain of 0.5, a strong strain hardening occurs which could be attributed to the straining of the network of PHA entangled nanofibers.
The formation of PHA nanofibers may also be responsible for an essential increase in the ductility of the PLA matrix. These PHA nanofibers, along with PLA nanofibers spanning the craze surfaces, dissipate a large amount of energy until they break, providing a continuous crack nano-bridge toughening right before the crack front. Figures 6 and 7 show the SEM images of crazes for the PLA/PHA blend and composite. For neat PLA, the crazes are thin, 1-2 µm, and do not thicken until failure [11]. Crazes in the PLA/PHA blend are almost the same size, 1-4 µm ( Figure 6). The slightly larger craze size is due to the fact that the PHA nanoparticles may toughen the matrix and facilitate the PLA nanofibers' formation during the development of crazes. However, for the PLA/PHA composite, crazes thicken with increasing strain up to a few tens of µm (Figure 7). until failure [11]. Сrazes in the PLA/PHA blend are almost the same size, 1-4 μm (Figure 6). The slightly larger craze size is due to the fact that the PHA nanoparticles may toughen the matrix and facilitate the PLA nanofibers' formation during the development of crazes. However, for the PLA/PHA composite, crazes thicken with increasing strain up to a few tens of μm (Figure 7). It has been shown in our previous paper [11] that neat PLA demonstrates a quite brittle fracture with severe localization of strain before yielding. The addition of 3 wt.% PHA to PLA increases the number of crazes. Despite some crack nano-bridge toughening by PLA nanofibers, which makes it possible to transfer loads to the rest of the material and avoid overloading heavily loaded crazes, the deformation mechanism remains crazing. Nevertheless, a certain amount of crazes provides some load-bearing capacity and leads to the improvement of the ductility of PLA/PHA blends. For the PLA/PHA composite, a neck appears after yielding. In the initial stages of deformation (e ≤ 0.20), intensive crazing is observed, but then at higher values of deformation (e > 0.22), crazing is replaced by the plastic flow of the material (Figure 8). until failure [11]. Сrazes in the PLA/PHA blend are almost the same size, 1-4 μm ( Figure 6). The slightly larger craze size is due to the fact that the PHA nanoparticles may toughen the matrix and facilitate the PLA nanofibers' formation during the development of crazes. However, for the PLA/PHA composite, crazes thicken with increasing strain up to a few tens of μm (Figure 7). It has been shown in our previous paper [11] that neat PLA demonstrates a quite brittle fracture with severe localization of strain before yielding. The addition of 3 wt.% PHA to PLA increases the number of crazes. Despite some crack nano-bridge toughening by PLA nanofibers, which makes it possible to transfer loads to the rest of the material and avoid overloading heavily loaded crazes, the deformation mechanism remains crazing. Nevertheless, a certain amount of crazes provides some load-bearing capacity and leads to the improvement of the ductility of PLA/PHA blends. For the PLA/PHA composite, a neck appears after yielding. In the initial stages of deformation (e ≤ 0.20), intensive crazing is observed, but then at higher values of deformation (e > 0.22), crazing is replaced by the plastic flow of the material (Figure 8). It has been shown in our previous paper [11] that neat PLA demonstrates a quite brittle fracture with severe localization of strain before yielding. The addition of 3 wt.% PHA to PLA increases the number of crazes. Despite some crack nano-bridge toughening by PLA nanofibers, which makes it possible to transfer loads to the rest of the material and avoid overloading heavily loaded crazes, the deformation mechanism remains crazing. Nevertheless, a certain amount of crazes provides some load-bearing capacity and leads to the improvement of the ductility of PLA/PHA blends. For the PLA/PHA composite, a neck appears after yielding. In the initial stages of deformation (e ≤ 0.20), intensive crazing is observed, but then at higher values of deformation (e > 0.22), crazing is replaced by the plastic flow of the material (Figure 8). It is known that neat PLA exhibits a strong strain softening, which stimulates strain localization and causes the build-up of local tri-axial stress. Since the local strain of neat PLA could not be delocalized, the local tri-axial stresses will induce void nucleation and crazes in the matrix leading to a brittle fracture behavior [27][28]. In the PLA/PHA blend, as well as the PLA/PHA composite, these tri-axial stresses could be released via intensive crazing i.e., the local strain in the PLA matrix would be delocalized. At a certain density of these crazes, the stress state in the PLA matrix could be converted from tri-axial to uni-axial. The new stress state is favorable for the initiation of shear bands leading to shear yielding of the matrix, as can be seen in Figure 9. Such a change in deformation mode from crazing to shear yielding was observed by us for the PLA/PBS system [11].  It is known that neat PLA exhibits a strong strain softening, which stimulates strain localization and causes the build-up of local tri-axial stress. Since the local strain of neat PLA could not be delocalized, the local tri-axial stresses will induce void nucleation and crazes in the matrix leading to a brittle fracture behavior [27,28]. In the PLA/PHA blend, as well as the PLA/PHA composite, these tri-axial stresses could be released via intensive crazing i.e., the local strain in the PLA matrix would be delocalized. At a certain density of these crazes, the stress state in the PLA matrix could be converted from tri-axial to uni-axial. The new stress state is favorable for the initiation of shear bands leading to shear yielding of the matrix, as can be seen in Figure 9. Such a change in deformation mode from crazing to shear yielding was observed by us for the PLA/PBS system [11].

Conclusions
The concept of the in situ generation of polymer-polymer nanocomposites was applied to a PLA/PHA blend. Green biopolymer-biopolymer nanocomposites combining biodegradability, biocompatibility, good strength of PLA, and good toughness of PHA were obtained. It was shown that a high shear rate allows the formation of nano-sized spherical particles of PHA at the stage of mixing, and then simultaneous fibrillation and crystallization at the extrusion stage at a higher temperature, without the necessity for succeeding cooling below the temperature of their non-isothermal crystallization. Shearing of PHA with 850 s −1 increased the non-isothermal crystallization temperature by 50 • C and decreased the temperature range in which the transition from a molten state to a crystallized one occurs by 17 • C. Using the slit capillary increased the residence time in the extruder and facilitated the formation of a fibril network at a low concentration of PHA. The formed PHA fibers had a diameter of nearly 200 nm and acted not only as reinforcements for the polymeric matrix, but also as toughening elements by providing continuous crack nano-bridging right before the crack front, resulting in an increase of both strength and plasticity of PLA. In particular, the Young modulus, stress at break, and elongation at break exceeded the related parameters of neat PLA by 15, 30, and 400% (2.35 GPa, 56.0 MPa, and 28.1%), respectively.

Supplementary Materials:
The following are available online at http://www.mdpi.com/2073-4360/11/11/1811/s1, Figure S1: Shear rate dependences of the complex viscosity for the matrix, PLA (a) viscosity ratio of blend to matrix (b) measured at 135, 155, 180 • C; Figure S2: Melting endotherm of PLA/PHA composite; Figure S3: Loss modulus of PLA, PHA and PLA/PHA composite in a tensile mode; Figure S4: True stress-Hencky strain curves of PHA nanofibers network for PLA/PHA nanocomposite determined at 100 • C and Hencky strain rates of 0.1, 0.2, and 0.4 sec-1. Plots obtained by subtracting stress-elongation curve for neat PLA from respective stress-elongation curves for the nanocomposite; Table S1: Mechanical properties of PLA, PLA/PHA blend and composite.