Toughening mechanism behind intriguing stress–strain curves in tensile tests of highly enhanced compatibilization of biodegradable poly(lactic acid)/poly(3-hydroxybutyrate-co-4-hydroxybutyrate) blends

Abstract

Highly enhanced compatibilization of biosourced and biodegradable polylactide (PLA) and poly(3-hydroxybutyrate-co-4-hydroxybutyrate) (P(3HB-co-4HB)) blends were successfully prepared by reactive melt compounding. Large shifts towards each other in terms of glass transition temperatures, a considerable reduction in the dispersed phase particle size and a significant increase in the interfacial adhesion between the PLA and P(3HB-co-4HB) phases were observed after compatibilization. In addition, chain branches occurred during the branching reaction decreased the crystallization ability of PLA, while crosslinks formed in the crosslinking reaction enhanced the crystallization ability of PLA on a large scale. Moreover, the blends exhibited a remarkable improvement of rheological properties of melt state when compared with that of blank PLA/P(3HB-co-4HB) blends. Upon increasing the content of the crosslinking agent, dicumyl peroxide (DCP), the blends showed increased yield tensile strength, modulus, and elongation at break. However, when DCP cooperated with triallyl isocyanurate (TAIC), the elongation at break decreased because the crosslinking network limited the mobility of the polymer chains to deform under a tensile load. Most notably, two typical and different kinds of growth of stress–strain curves were observed, and for the first time we demonstrated the toughening mechanism behind it in detail. Furthermore, SEM images of the fracture surfaces of the blends confirmed the toughening mechanism and that plastic deformation of the matrix and a debonding process were the two important ways of induced energy dissipation leading to toughened blends.

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Introduction

Polymer blends have received much attention due to combining attractive features of each component whilst reducing their deficient characteristics at the same time. The performance of polymer blends is determined not only by the properties of each component but also the morphology formed.^1–4 In fact, most polymer blends are immiscible because of high interfacial tension and poor adhesion. However the mechanical properties of a multiphase system are usually driven by the ability of the interface to transmit stress from one phase to another.⁵ Therefore, the compatibilization is very necessary for immiscible polymer blends. This can be obtained by introducing a third component to refine the droplet size of the dispersed minor phase, to lower the interfacial tension and stabilize it against coalescence, and to ensure strong interfacial adhesion between the phases.⁶ A considerable number of research papers have been published on introducing nanoscale fillers or pre-made copolymers into polymer blends.^7–19 Compared with a third compatibilizer showing limited efficiency due to only relying on physical interactions, a free radical initiator, i.e., dicumyl peroxide (DCP), was introduced into a polylactide/poly(ε-caprolactone) blend and polylactide/poly(butylene succinate) blend to induce in situ compatibilization on a large scale.^20–22

PLA has excellent high tensile strength and good biocompatibility, but also shows brittleness and difficulty in process. On the contrary, poly(3-hydroxybutyrate-co-4-hydroxybutyrate), P(3HB-co-4HB), with a high molar fraction of 4-hydroxybutyrate (4HB), is a promising bio-elastomer possessing high flexibility in the large polyhydroxyalkanoates (PHAs) family.^23,24 Biobased microbial PHAs, produced by Ralstonia eutropha, Alcaligenes latus, and Comamonas acidovorans,^25–27 have attracted universal attention due to their renewable resources, biodegradable properties, biocompatibility, and potential applications as environmental friendly polymers for agricultural, marine, and medical applications. Therefore, to blend PLA and P(3HB-co-4HB) becomes a reasonable choice to improve the flexibility and toughness of PLA for many other potential applications. In order to enhance the in situ compatibilization, crosslinking is also introduced in the work by adding triallyl isocyanurate (TAIC) to cooperate with DCP. Since TAIC, used in industry as a common crosslinking agent for polyolefin and vinyl polymers,²⁸ has also been reported to be an effective crosslinking agent for PLA due to double bonds in TAIC to improve the efficacy of free radicals.²⁹

In this study, biodegradable and biosourced PLA/P(3HB-co-4HB) blends were prepared by melt compounding using crosslinking agents to enhance in situ compatibilization. The aim of this work was to investigate the different effects of branching and crosslinking on the morphology and compatibilization of the blends. Most importantly the profound mechanical mechanism of the system was discussed in detail. Furthermore, the thermal and rheological properties of the resultant blends were investigated.

Experimental

Materials

The PLA (Grade 4032D) used in this work was a commercially available product from NatureWorks LLC (USA). It exhibited a weight-average molecular weight of 207 kg mol⁻¹ and a polydispersity of 1.74 as determined by gel permeation chromatography (GPC). The bacterial copolymer P(3HB-co-4HB) was provided by Tianjin Guoyun Biotech (China). The number-average molecular weight was 195 kg mol⁻¹ with a polydispersity of 1.86 as determined by GPC. The content of 4HB in the copolymer was 23.9 mol% as determined by ¹H NMR spectroscopy. The free radical initiator DCP was purchased from Beijing Chemical Company (China). TAIC was purchased from Aldrich. All of these materials were used as received.

Blend preparation

Before processing, the two polymers were dried in a vacuum oven at 70 °C for 24 h. A certain amount of DCP and/or TAIC were/was dissolved precisely in acetone. The solution was mixed with the two polymers and dried in a vacuum oven at 50 °C for 12 h. The mixtures were mixed in an internal mixer (Rheomix 600p, Haake, Karlsruhe, Germany) at 165 °C and 50 rpm for 6 min. The samples were hot-pressed at 185 °C for 3 min and then cold-pressed at room temperature to form 1 mm thick sheets. The resultant samples were designated as x/y-Da/Tb. Here x and y denote the weight percentage of PLA and P(3HB-co-4HB), and a and b denote the weight percentage of DCP and TAIC, respectively. For comparison, neat PLA and neat P(3HB-co-4HB) were treated using the same procedure.

Characterization

The gel content of the branched/crosslinked P(3HB-co-4HB) was determined gravimetrically with a Soxhlet extraction cycle using boiling chloroform as the solvent for 24 h. Approximately 0.2 g samples were cut into small pieces and wrapped in a pre-weighted, quantitative filter paper. After the extraction, the samples were vacuum-dried to a constant weight. The gel fraction was calculated as follows:

Gel fraction = (M_d/M_i) × 100% (1)

where M_i is the initial weight of the sample and M_d is the dry weight of the sample after extraction.

Dynamic mechanical analysis (DMA) was carried out using a dynamic mechanical analyzer SDTA861^e (Mettler Toledo) in tensile mode. Samples with gauge dimensions of 20 × 4 × 1 mm³ were used. The dynamic loss factor (tan δ) and the storage modulus (E′) were determined at a frequency of 1 Hz and a heating rate of 3 °C min⁻¹ as a function of temperature from −60 to 120 °C.

The phase morphology of the blends were investigated using field emission scanning electron microscopy (XL30 ESEM FEG, FEI Co., Eindhoven, The Netherlands) at an accelerating voltage of 10 kV. The samples were immersed in liquid nitrogen for about 5 min, and then broken. Because of the similar physical properties of PLA and P(3HB-co-4HB) as aliphatic polyester, direct observation of the cryo-fractured surfaces of the PLA/P(3HB-co-4HB) blends using SEM to obtain the obvious dispersed phase morphology is difficult. Therefore, the selective enzymatic degradation method was used. For the PLA matrix samples, removal of the P(3HB-co-4HB) component from the cryo-fractured surfaces of the blends, allows the remaining morphology to be observed. The selective enzymatic degradation of the blends was carried out in phosphate buffer (pH 7.4) containing lipase from Pseudomonas mendocina at 37 °C with shaking at 140 rpm. In addition, crosslinked P(3HB-co-4HB) with a gel fraction of about 40% was treated using the same procedure to insure even crosslinked P(3HB-co-4HB) could be degraded by the lipase (the lipase from Pseudomonas mendocina revealed that it prefers the enzymatic degradation of P(3HB-co-4HB) but does not attacks the PLA in the blends³⁰). When the P(3HB-co-4HB) component on the surface of samples was degraded, the samples were removed, washed with distilled water, and dried to a constant weight in vacuum. In addition, for the P(3HB-co-4HB) matrix samples, removal of the PLA component from the cryo-fractured surfaces by proteinase K was carried out using the same procedure.³¹ Crosslinked PLA with a gel fraction of about 40% was also treated using the same procedure to insure that the crosslinked PLA was degraded by proteinase K. The degraded cryo-fractured surfaces of all the samples were sputter-coated with a thin layer of gold and observed using SEM.

Differential scanning calorimetry (DSC) experiments were carried out using a TA Instruments DSC Q20 (USA) under an N₂ atmosphere. The specimens were sealed in aluminum crucibles and had a nominal weight of about 5 ± 0.3 mg. The samples were heated from 40 °C to 185 °C at a heating rate of 30 °C min⁻¹ (first heating), held for 2 min to erase the previous thermal history, then cooled to 30 °C at a rate of 5 °C min⁻¹ (first cooling). The samples were further heated to 185 °C again from 30 °C at a heating rate of 20 °C min⁻¹ (second heating).

Rheological properties were measured using a rotational rheometer (TA Series AR2000ex, TA Instrument, USA). The compression-molded samples were cut into the disks (25 mm in diameter and 1 mm in thickness). The measurements were carried out in dynamic (oscillatory) mode by means of 25 mm parallel geometry at 180 °C under an air atmosphere. Amplitude sweeps were performed in advance to ensure that the dynamic tests were in the linear viscoelastic range and a strain value of 1.25% was consequently chosen. The frequency ranged from 0.1 to 100 rad s⁻¹.

Uniaxial tensile tests were carried out on dumbbell shaped specimens (20 × 4 × 1 mm³) that were punched out from the pressed sheets. The measurements were performed using a tensile-testing machine (Instron-1121) according to GB/T1040-2006 (China) at room temperature at a crosshead speed of 10 mm min⁻¹. At least five specimens were tested for each sample and the average value reported.

Results and discussion

Gel analysis of the PLA/P(3HB-co-4HB) blends

Both PLA and P(3HB-co-4HB) can form branched and/or crosslinked structures with crosslinking agents as pure materials as shown in Fig. 1. At the interface of the PLA/P(3HB-co-4HB) blends, branching/crosslinking can occur via a combination of PLA and P(3HB-co-4HB) free radicals. It is notable that the combination reaction of free radicals not only occurs at the interface, but can also occur in the PLA and P(3HB-co-4HB) domains. As a consequence, complex reaction products could be obtained, including branched/crosslinked PLA, branched/crosslinked P(3HB-co-4HB), PLA-g-P(3HB-co-4HB) copolymers, and a PLA-crosslink-P(3HB-co-4HB) network. Furthermore, melt blending was accompanied by chain scissions due to the thermal instability of the two types of polymers and the instability of their free radicals, resulting in even more complicated products. Fig. 1 shows the gel fraction of these blends. Obviously, there was no gel formed until the addition of TAIC as a co-crosslinking agent, and the gel fraction of the PLA/P(3HB-co-4HB) blends increased to near 40% by only adding 0.1 wt% of DCP and 0.1 wt% of TAIC. At a certain amount of DCP, only branching and/or chain scissions reactions occurred, and from the later measurements of their rheological properties, the complex viscosities of these blends were higher than that of the blank blend, indicating the branching reaction predominated over chain scissions. While DCP cooperated with TAIC, the crosslinking reaction became dominant. Additionally, TAIC proved to be an effective co-crosslinking agent in combination with DCP. It also can be observed that P(3HB-co-4HB) caused less chain scissions than that of PLA in the presence of free radicals because of the relatively higher gel fraction gained in crosslinked P(3HB-co-4HB) when adding the same amounts of crosslinking agents.

Fig. 1 Gel fraction of PLA, P(3HB-co-4HB), and PLA/P(3HB-co-4HB) blends.

Miscibility and phase morphology

It is well known that when the two polymers are miscible in the amorphous phase they will show a single glass transition temperature (T_g), while the appearance of two T_gs corresponding to each individual component is characteristic of immiscible blends. In partially miscible blends the T_gs of each component shift towards each other. Fig. 2(a) and (b) shows tan δ vs. temperature for neat polymers and their blends. Each neat component exhibited a single relaxation peak corresponding to T_g. For all the blends, tan δ curves revealed two T_gs, the higher T_g corresponded to the PLA component, and the lower one corresponded to the P(3HB-co-4HB) component. Moreover, the T_gs of both the PLA and P(3HB-co-4HB) components showed almost no shift towards each other in the blank and branched blends, suggesting that PLA and P(3HB-co-4HB) were immiscible in these blends. However, one should notice that for the crosslinked blends, owning to the formation of the PLA-g-P(3HB-co-4HB) and PLA-crosslink-P(3HB-co-4HB) network, the T_gs of PLA and P(3HB-co-4HB) components showed large shifts toward each other, indicating enhanced compatibilization. The crosslinked blends became partially miscible. Additionally, the difference of glass transition (ΔT_g) between the PLA component and P(3HB-co-4HB) component in each blend was calculated and shown in Table 1; 30/70 blends showed smaller ΔT_gs than those of the 70/30 blends, indicating the 30/70 blends achieved a finer compatibility than that of the 70/30 blends while adding the same amounts of DCP and TAIC. As shown in Fig. 2(c) and (d), the storage modulus (E′) at room temperature for the PLA/P(3HB-co-4HB) blends are above neat PLA and P(3HB-co-4HB), indicating both a branching structure and crosslinking network can enhance the chain entanglement. Additionally, the improvement of E′s in the 70/30 blends are a little lower than those of the 30/70 blends. This was because the PLA component suffered heavier thermal decomposition at the processing temperature than that of P(3HB-co-4HB) as mentioned above. In Fig. 2(c), the E′s of the blends dropped sharply about 0 °C due to the T_g of P(3HB-co-4HB), and then decreased again around 60 °C due to the T_g of PLA, but began to increase about 100 °C due to the cold crystallization of PLA. Moreover, the cold crystallization peak (T_cc) of PLA shifted to a lower temperature (90 °C) for the 70/30 blend and branched blends, and moved to 80 °C when crosslinking happened. These results suggested that the incorporation of both the P(3HB-co-4HB) and crosslinking network enhanced the cold-crystallization ability of PLA and decreased the T_cc of PLA in the blends. In Fig. 2(d), there was only a clear sharp drop around 0 °C owing to the T_g of P(3HB-co-4HB), and then increased about 100 °C due to the cold crystallization of PLA and the same enhanced cold-crystallization of PLA were observed in the 30/70 blends.

Fig. 2 DMA traces of PLA/P(3HB-co-4HB) blends, tan δ versus temperature of (a) 70/30 blends, and (b) 30/70 blends. The E′ versus temperature of (c) 70/30 blends, and (d) 30/70 blends (the insets give details of the transitions).

Table 1 Thermal and crystalline properties of PLA/P(3HB-co-4HB) blends

Sample	DMA			DSC
	Tg,P(3HB-co-4HB) (°C)	Tg,PLA (°C)	Tg (°C)	Tg (°C)	Tcc (°C)	ΔHcc^a (J g^−1)	Tc (°C)	ΔHc^a (J g^−1)	Tm1 (°C)	Tm2 (°C)	ΔHm^a (J g^−1)	Xc^b (%)
a ΔHcc and ΔHm are corrected for the content of PLA in the blend.b Degree of crystallinity, calculated from the ratio of ΔHm and ΔH^0m (the melting enthalpy ΔH^0m of 100% crystalline PLA was taken as 93 J g^−1).
PLA	—	70.2	—	63.9	118.4	34.6	93.3	2.3	—	165.0	36.6	39.4
70/30	−3.5	67.6	71.1	59.4	—	—	118.4	37.6	—	166.1	43.6	46.9
70/30-D0.05	−3.5	67.5	71.0	62.5	116.4	32.4	—	—	162.7	166.6	38.0	40.9
70/30-D0.1	−3.4	67.1	70.5	62.4	118.2	31.1	—	—	163.1	—	34.8	37.4
70/30-D0.1T0.1	−1.1	65.4	66.5	62.7	—	—	131.2	44	—	165.5	48.8	52.5
70/30-D0.2T0.2	−0.9	65.2	66.1	63.5	—	—	130.9	42.8	—	165.9	47.7	51.3
30/70	0.1	68.9	68.8	61.9	114.9	28.3	—	—	164.3	168.1	37.0	39.8
30/70-D0.05	0.1	69.0	68.9	61.9	117.2	26.3	—	—	162.7	—	35.7	37.5
30/70-D0.1	0.5	69.0	68.5	61.4	118.8	24.7	—	—	162.9	—	31.3	33.6
30/70-D0.1T0.1	0.9	65.7	64.6	62.1	—	—	120.7	35.3	163.6	—	44.3	47.6
30/70-D0.2T0.2	0.8	64.1	63.3	—	—	—	128.5	34	163.3	—	42.0	45.2
P(3HB-co-4HB)	−1.6	—	—	—	—	—	—	—	—	—	—

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Fig. 3 presents the typical matrix-droplet morphology of the cryo-fractured surfaces of the PLA/P(3HB-co-4HB) blends after selective enzymatic degradation. The black pores appeared upon the removal of the dispersed phase on the cryo-fractured surfaces of the blends. It was found that the blends displayed a clear fine dispersion. In Fig. 3(a) and (f), the particle sizes of these pores were very large and non-uniform. However, in Fig. 3(b–e) and (g–m), a much finer and more uniform dispersion of the dispersive phase was obtained with an increase of DCP and TAIC, attributed to the in situ formation of PLA-g-P(3HB-co-4HB) copolymers and the PLA-crosslink-P(3HB-co-4HB) network, which acted as a compatibilizer between the PLA and P(3HB-co-4HB) domains. Moreover, crosslinked PLA and crosslinked P(3HB-co-4HB), all these reaction products not only decreased the viscosity ratio of the two components but also increased the physical and chemical entanglement in the system, which prevented the coalescence of dispersed phase domains during melt mixing. So a large improvement of compatibility and strong interfacial adhesion between the PLA and P(3HB-co-4HB) phases was achieved. In addition, a much finer and more uniform dispersion was obtained in the crosslinked blends than that of branched blends. What was more, the 30/70 blends displayed a more uniform and smaller average particle size in the dispersed phase than those of the 70/30 blends with the same amount of crosslinking agents, especially for 30/70-D0.05 vs. 70/30-D0.05 and 30/70-D0.1 vs. 70/30-D0.1, which was in agreement with the results obtained from the DMA measurements. To further explore the reason, the complex viscosities of neat PLA, P(3HB-co-4HB) and their blank blends are shown in Fig. 4. At the given temperature (180 °C), the complex viscosity (|η*|) of neat PLA is twice that of neat P(3HB-co-4HB). The |η*| of the 70/30 blend is a little higher than that of the 30/70 blend. Moreover, the lower viscosity of P(3HB-co-4HB) could be broken into smaller droplets and stabilized more easily than PLA. So more PLA-g-P(3HB-co-4HB) copolymers and/or PLA-crosslink-P(3HB-co-4HB) networks occurred at the interface during the reactive blending for the 30/70 blends. Therefore, more uniform and smaller particle sizes were achieved in the 30/70 blends.

Fig. 3 SEM micrographs of cryo-fractured surfaces of the PLA/P(3HB-co-4HB) blends after selective enzymatic removal of dispersed phase domains, (a) 70/30, (b) 70/30-D0.05, (c) 70/30-D0.1, (d) 70/30-D0.1T0.1, (e) 70/30-D0.2T0.2, (f) 30/70, (g) 30/70-D0.05, (h) 30/70-D0.1, (l) 30/70-D0.1T0.1, and (m) 30/70-D0.2T0.2.

Fig. 4 Plots of complex viscosity |η*| versus frequency.

Thermal behaviour

The DSC results for the thermal behaviour of the PLA/P(3HB-co-4HB) blends are shown in Fig. 5 and the parameters of their thermal properties listed in Table 1. Neat PLA showed a weak crystallization ability with a weak crystallization peak (T_c) (93.3 °C), a strong T_cc (118.4 °C), and a melting endothermic peak (T_m) (165.0 °C). While neat P(3HB-co-4HB) was a soft elastomeric material with a T_g of only −8.9 °C. Compared with neat PLA in Fig. 5(a) and (b), the 70/30 blend exhibited a strong sharp T_c without a T_cc, indicating an enhanced crystallization ability. It probably resulted from the interface between the phase-separation domains where favourable nucleation sites exist due to immiscibility.³² However, upon adding DCP, the T_c disappeared and two relatively weaker T_ccs and a lower T_ms occurred, indicating chain branches decreased the crystallization of PLA. However, for the crosslinked blends, the T_c appeared and moved to a higher temperature with an even sharper shape than that of the 70/30 blend, so did the T_ms, demonstrating that crosslinking enhanced the crystallization of PLA. This was attributed to the disordered short chain branches formed during the branching and/or chain scissions reactions, which reduced the regularity in the original linear polyester backbone of PLA and then inhibited the crystallization process. While a controlled amount of crosslinks formed during the crosslinking reaction, this did not restrict the mobility of polymer chains among the crosslinks to crystallize, in contrast they served as nucleation sites, which enhanced the crystallization ability of PLA at given cooling rate. It is true that branching points can also be regarded as nucleation sites to increase the nucleation process in long-chain branched PP for instance,³³ but all those short chain branches reduce the local regularity and inhibit the movable polymer chains to crystallize. Detailed discussions were reported in our previous work.³⁴ Accordingly, 70/30-D0.05 and 70/30-D0.1 decreased the crystallization ability of PLA.

Fig. 5 DSC curves of the PLA/P(3HB-co-4HB) blends, (a) first cooling of the 70/30 blends, (b) second heating of the 70/30 blends, (c) first cooling of the 30/70 blends and (d) second heating of the 30/70 blends.

As shown in Fig. 5(c) and (d), the thermal behaviour of the PLA/P(3HB-co-4HB) 30/70 blends displayed the same tendency, except for the 30/70 blend, whose T_c was too low to be detected owing to the low weight content of PLA. What was more, the crystallinity of PLA in the 30/70 blends were relatively lower than those of the 70/30 blends with the same content of crosslinking agents. This finding is related to the miscibility of the blends. The PLA and P(3HB-co-4HB) phases showed relatively poorer compatibility and weaker interfacial adhesion in the 70/30 blends compared to those in the 30/70 blends with the same contents of DCP and TAIC, which was confirmed by the DMA measurements and SEM results described beforehand. So more branched/crosslinked PLA, PLA-g-P(3HB-co-4HB) copolymers, and the PLA-crosslink-P(3HB-co-4HB) network occurred in the 30/70 blends, which reduced the PLA domains and the final crystallinity of the blends.

Rheological properties

The effect of frequency on the complex viscosity |η*|, the storage modulus G′ (solid symbols), the loss modulus G′′ (open symbols) and tan δ for the PLA/P(3HB-co-4HB) blends are shown in Fig. 6. It was obvious that the |η*| of the blank 70/30 and 30/70 blends exhibited a Newtonian liquid behaviour, in contrast, the |η*| of other blends showed non-Newtonian shear-thinning behaviour as shown in Fig. 6(a) and (d). As expected, the |η*| diverged at low shear rates with high gel fractions. Especially, both the |η*| and G′ increased with an increase of crosslinking agents. In the low frequency region, both 70/30-D0.2T0.2 and 30/70-D0.2T0.2 exhibited more solid-like behaviour than those of the blank blends. This was because the branching points and/or crosslinks formed a chain network that enhanced the interactions among the macromolecules. However, there were exceptions for 70/30-D0.1 and 30/70-D0.1, which showed lower |η*| and G′ values than those of 70/30-D0.05 and 30/70-D0.05 accordingly, the reason may be attributed to the competitive reactions between branching and chain scissions with the active free radicals. In addition, the branching reaction was dominant over chain scissions with low amounts of DCP (D0.05), while chain scissions exceeded this as more free radicals were initiated by higher contents of DCP (D0.1). Both the |η*| and G′ of the 70/30 blends were lower than those of the 30/70 blends with the same content of crosslinking agents. This was because chain entanglement was enhanced to a higher level in the 30/70 blends than that in the 70/30 blends. Moreover, the tan δ (tan δ = G′′/G′) value, the ratio of energy dissipated to energy stored per cycle, is shown in Fig. 6(c) and (f). The branched blends were predominantly viscous, tan δ > 1 over the entire frequency range, and the peaks of the curves shifted towards a low frequency upon increasing the DCP content due to the formation of the branched structure. In addition, the tan δ of 70/30-D0.1 was higher than that of 70/30-D0.05 because more chain scissions occurred as mentioned beforehand, so was 30/70-D0.1. Other crosslinked blends showed tan δ < 1 over the entire frequency range, indicating a solid-like behaviour.

Fig. 6 Plots of: (a) complex viscosity |η*|, (b) storage modulus G′ (solid) and loss modulus G′′ (open), and (c) tan δ versus frequency of the 70/30 blends. Plots of (d) complex viscosity |η*|, (e) storage modulus G′ (solid) and loss modulus G′′ (open), and (f) tan δ versus frequency of the 30/70 blends.

In Fig. 7, the phase angle δ is plotted versus the absolute value of the complex modulus |G*| for the PLA/P(3HB-co-4HB) blends. In the literature, this plot is known as the van Gurp–Palmen plot, which can be used to detect the rheological properties of branched/crosslinked polymers.¹⁷ The δs of the blank 70/30 and 30/70 blends were close to 80° and became smaller when only adding DCP, but all indicated flow behaviour, which was presented by a viscoelastic fluid. In contrast, the lightly crosslinked PLA/P(3HB-co-4HB) blends with a network structure resembled the behaviour of a fluid–solid transition with a corresponding phase angle near 45°. While 70/30-D0.2T0.2 and 30/70-D0.2T0.2 presented a phase angle of 20°, owing to heavier crosslinking, finer compatibility and stronger interfacial adhesion among the PLA and P(3HB-co-4HB) phases.

Fig. 7 van Gurp–Palmen plots of the phase angle (δ) versus complex modulus (|G*|) of (a) the 70/30 blends and (b) 30/70 blends.

Mechanical properties

The PLA used in this study is brittle with a low elongation at break (5%), while P(3HB-co-4HB) is ductile with an elongation at break of 2122% and yield strength of 5.3 MPa. The elongation at break of PLA was improved on a large scale by the incorporation of P(3HB-co-4HB). Typical stress–strain curves of PLA/P(3HB-co-4HB) blends are shown in Fig. 8, and the detailed data of mechanical properties listed in Table 2. Compared with other mechanical properties, the elongation at break (ε_b) is more sensitive to the compatibilization effect in immiscible binary blends. The blank 70/30 blend presents an improved ε_b of 186%. In the presence of 0.05 wt% of DCP, the ε_b increased to 239%, indicating ductile behaviour. This could be ascribed to its enhanced interfacial adhesion. With 0.1 wt% of DCP, the ε_b of the blend was increased to 317%, the yield tensile strength and modulus were decreased slightly. But with the addition of both DCP and TAIC, the crosslinking network formed, which restricted the mobility of the polymer chains to dissipate energy under a tensile load. The ε_b finally decreased to 251%, while the yield tensile strength and modulus were improved slightly. Thus, the tensile toughness of the 70/30 blends was significantly enhanced.

Fig. 8 Tensile stress–strain curves of (a) the 70/30 blends and (b) 30/70 blends.

Table 2 Mechanical properties of the PLA/P(3HB-co-4HB) blends

Sample	Yield tensile strength (MPa)	Modulus (MPa)	Elongation at break (%)	Sample	Yield tensile strength (MPa)	Modulus (MPa)	Elongation at break (%)
PLA	63.4 ± 0.5	1985 ± 17	5 ± 1	P(3HB-co-4HB)	5.3 ± 1.9	85 ± 25	2122 ± 30
XPLA	69.4 ± 0.6	1990 ± 20	58 ± 23	XP(3HB-co-4HB)	5.6 ± 1.1	92 ± 21	1863 ± 33
70/30	29.6 ± 0.3	1227 ± 19	186 ± 21	30/70	14.5 ± 0.4	239 ± 12	145 ± 25
70/30-D0.05	29.2 ± 0.2	785 ± 17	239 ± 25	30/70-D0.05	7.0 ± 0.3	204 ± 15	564 ± 30
70/30-D0.1	28.2 ± 0.3	801 ± 13	317 ± 19	30/70-D0.1	7.8 ± 0.2	174 ± 10	593 ± 22
70/30-D0.1T0.1	29.3 ± 0.3	999 ± 16	310 ± 28	30/70-D0.1T0.1	11.0 ± 0.2	297 ± 14	410 ± 28
70/30-D0.2T0.2	31.2 ± 0.2	1123 ± 15	251 ± 24	30/70-D0.2T0.2	11.6 ± 0.3	338 ± 13	289 ± 31

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The overall mechanical properties of the 30/70 blends were improved as well after in situ compatibilization, as shown in Table 2. All the elongation at break, yield tensile strength and modulus displayed the same trend with that found with the 70/30 blends. As expected, the elongation at break increased sharply, yield tensile strength and modulus of 30/70 blends decreased greatly due to the incorporation of increased amounts of the soft elastomeric component, P(3HB-co-4HB).

To explore the fracture behaviour of the PLA/P(3HB-co-4HB) blends in tensile tests, the morphology of the fracture surfaces of the specimens was observed using SEM. Fig. 10 shows the SEM micrographs of the tensile fractured surface and the surface parallel tensile direction near the broken point of the blends. For comparison, neat PLA, XPLA, neat P(3HB-co-4HB) and XP(3HB-co-4HB) are shown in Fig. 9 (PLA and P(3HB-co-4HB) were crosslinked separately by adding 0.1 wt% of DCP and 0.1 wt% of TAIC, the resulting samples were designated as XPLA and XP(3HB-co-4HB) with a gel fraction of about 40%). Neat PLA, showed a fairly smooth fracture surface and typical brittle fracture, the parallel surface was flat as shown in Fig. 9(b) and (b′). For XPLA shown in Fig. 9(c), the fracture surface showed a large amount of cavitations and large-scale plastic deformation or fibrillation, which was related to the enhanced chain entanglement caused by its crosslinking network. But the parallel surface was still flat as shown in Fig. 9(c′). On the other hand, in Fig. 9(d, e, d′ and e′) both neat P(3HB-co-4HB) and XP(3HB-co-4HB) exhibited large scale plastic deformation without a fibrous structure on the fracture surface and the parallel surface became rough. What was more, one can observe small particles in the P(3HB-co-4HB) matrix, which was in common with the findings of other researchers.^35,36 As for the blank 70/30 blend, the fracture surface changed from smooth to rough compared with that of neat PLA. There presented many little particles with clear interface and the PLA matrix started to deform with visible plastic deformation (Fig. 10(a)) and the parallel surface showed protuberant ridge lines and some non-uniform pores (Fig. 10(a′)), which implied a ductile fracture behaviour. This was related to the elastic P(3HB-co-4HB) particles acting as stress concentrators. The consequent stress concentration leads to the development of triaxial stress in the P(3HB-co-4HB) particles. Because of the lack of phase adhesion, debonding can easily take place at the particle–matrix interface in the perpendicular external stress direction.³² Thus, in Fig. 10(a′) the cavities raised and were more clearly observed on the parallel surface. The fracture surface of 70/30-D0.05 was rough as well and had obvious ridges due to plastic deformation as shown in Fig. 10(b). In addition, the P(3HB-co-4HB) particles seemed to increase in size due to an enhanced phase adhesion to form a thicker interface layer, which bedded the particles. In Fig. 10(b′) the ridge lines on the parallel surface became blurred with small pores due to the enhanced plasticity of the PLA matrix. For 70/30-D0.1 shown in Fig. 10(c), the fracture surface became more and more rough, which was related to the ease of plastic deformation of the PLA chains induced by the finer interface adhesion and stronger entanglement between the two polymer domains caused by branching. The ridge lines on the parallel surface turned into lots of gibbosities in Fig. 10(c′). Especially for the blends with crosslinking network formed in Fig. 10(d) and (e), the fracture surfaces showed a large amount of cavitations, large-scale plastic deformation and a fibrous structure occurred, which implied a dramatic toughening effect. Moreover in Fig. 10(d′) and (e′) the parallel surface turned flat, which was similar to that of XPLA shown in Fig. 9(b′). The cavitations and plastic deformation induced energy dissipation and therefore led to the improvement in tensile toughness observed for the PLA/P(3HB-co-4HB) blends.^37,38

Fig. 9 (a) A schematic diagram of the measurement locations (B), fractured surface, and (C), surface parallel tensile direction near the broken points and the SEM of the PLA/P(3HB-co-4HB) blends in tensile strength tests, (b, b′) PLA, (c, c′) XPLA, (d, d′) P(3HB-co-4HB) and (e, e′) XP(3HB-co-4HB) in tensile strength tests. (b–e, fractured surface, b′–e′, surface parallel tensile direction near the broken points).

Fig. 10 SEM micrographs of the PLA/P(3HB-co-4HB) blends during tensile strength tests, (a–m, fractured surface, a′–m′, surface parallel tensile direction near the broken points), (a, a′) 70/30, (b, b′) 70/30-D0.05, (c, c′) 70/30-D0.1, (d, d′) 70/30-D0.1T0.1, (e, e′) 70/30-D0.2T0.2, (f, f′) 30/70, (g, g′) 30/70-D0.05, (h, h′) 30/70-D0.1, (l, l′) 30/70-D0.1T0.1, and (m, m′) 30/70-D0.2T0.2.

As for the five 30/70 blends shown in Fig. 10(f–m), the particles were more apparently viewed after tensile tests than those of the 70/30 blends. In Fig. 10(f), the PLA particles were exposed and only P(3HB-co-4HB) deformed under stretching. The parallel surface was rough with some pores on it. In addition, by adding DCP, the PLA particles trimmed down (Fig. 10(g) and (h)) and huge plastic fibrillation of the P(3HB-co-4HB) matrix was achieved. The parallel surface showed many gibbosities as shown in Fig. 10(g′) and (h′) with a few little pores. When the crosslinking network was introduced in 30/70-D0.1T0.1, it restricted the flexible P(3HB-co-4HB) polymer chains to deform to a certain extent, the ability of plastic deformation was limited (Fig. 10(l)), and the parallel surface showed a uniform accordion-structure perpendicularly along the tensile direction as shown in Fig. 10(l′). As for 30/70-D0.2T0.2, shown in Fig. 10(m), the plastic deformation of the P(3HB-co-4HB) matrix was confined on a large-scale, and the parallel surface displayed clear striking ridge lines. The plastic deformation of the matrix and the debonding process were the two important ways for induced energy dissipation and led to a toughened, biodegradable polymer blend. The conclusion is that both the compatibility between the PLA and P(3HB-co-4HB) components and the entanglement among the polymer chains are both necessary for toughness. The important point is that fine toughness requires not only a high level of adhesion between the particles and matrix, but also active molecular mobility, which is a crucial factor for yield stress and plastic flow.

To further investigate the toughening mechanism of the PLA/P(3HB-co-4HB) blends, amplified images of the strain-hardening area of the stress–strain curves are shown in Fig. 8 for its stable increase of stress with growing strain. For comparison, neat PLA, XPLA, neat P(3HB-co-4HB) and XP(3HB-co-4HB) are shown in Fig. 11. In Fig. 11(a), neat PLA showed no distinct yield point with subsequent failure by neck instability, while XPLA with a gel fraction of about 40%, exhibited an enhanced yield stress and ε_b, whose stress–strain curve showed “pulse-growth” after yielding. For neat P(3HB-co-4HB) and XP(3HB-co-4HB) shown in Fig. 11(b), clear yielding and stable neck growing through cold drawing was observed, the stress–strain curve showed a “step-growth” with increasing strain, but the steps of XP(3HB-co-4HB) were squeezed. As for the five 70/30 blends, the stress–strain curves all showed pulse-growth as strain increased, and in the five 30/70 blends they became even more complicated, including both step-growth and a combination of pulse-growth and step-growth. This phenomenon was also observed in our previous work,³² unfortunately, due to a lack of detailed investigation, the mechanism behind this has puzzled us for a long time. Here we attempted to depict the pulse-growth and step-growth as follows. The stress–strain curve of a typical ductile material owning a very long strain-hardening region, like P(3HB-co-4HB) used in this work, shows step-growth. That is, as the strain grows under a tensile load, the stress arises accordingly, and when the strain continues to grow, the stress will stop at a certain point and keep at that value, implying short stress-holding behaviour during which a certain amount of flexible entangled polymer chains begin to flow or orientate to dissipate energy. In addition, when the strain keeps on growing, another new short stress-holding process happens with a higher stress value. The apparent phenomenon is the stable growing of the neck region from its two ends. In contrast, a stiff and brittle polymer, like PLA, shows no distinct yield point before its failure. After proper crosslinking or toughened upon the addition of small amounts of another soft material, such as P(3HB-co-4HB), the stress–strain curve of the brittle polymer will exhibit pulse-growth. That is, as the strain grows under a tensile load, the stress increases to a point and suddenly drops slightly with the continuing growth of strain because there are only a limited number of rigid polymer chains (compared with P(3HB-co-4HB) chains) to orientate. The new and continuously emerging rigid polymer chains with a very limited degree of orientation are only available to dissipate energy during the neck growing process. So when the strain keeps on growing, another new pulse-growth happens. But the apparent phenomenon is the same as the growing of a neck region. Compared with pulse-growth, in order to hold that stress, it needs larger number of flexible polymer chains to orientate in the step-growth. As described by other researchers,³⁹ in reality, the fracture mechanisms of polymeric materials are a combination of chain scission and slippage (pull-out), which is governed by entanglements and the polymers molecular weight. Therefore, an enhanced entanglement of polymer chains by introducing an appropriate amount of crosslinks in XPLA can restrict the chain scission or slippage to some extent, thus giving opportunities for a small amount of polymer chains to flow, compared with only rigid PLA chains with limited physical entanglement. In addition, since there is only 30% flexible polymer chains of P(3HB-co-4HB) in the 70/30 blends, the stress–strain curves of crosslinked XPLA and the 70/30 blends all showed pulse-growth due to their relatively rigid polymer chains. Meanwhile for the 30/70 blends, they included both step-growth with flexible entangled polymer chains, and a combination of pulse-growth and step-growth with the decreased flexibility of their polymer chains. Furthermore, we can infer that for those polymer chains, which are not only very flexible but also entangled together an appropriate amount, like rubber, the stress–strain curves will grow smoothly due to the easy flow of polymer chains in response to applied stress. Sometimes, one can introduce an appropriate amount of chemical entanglement to enhance the weak and low degree of limited physical entanglement.

Fig. 11 Tensile stress–strain curves of (a) PLA and XPLA, (b) P(3HB-co-4HB) and XP(3HB-co-4HB).

Additionally, the detailed average Δσ and Δε values are listed in Table 3 (in step-growth, Δσ is the difference of two stress values between the two steps; and in pulse-growth, Δσ is the difference of peak stress value minus the following valley one. All the Δε values were calculated according to the Δσ). In the 70/30 blends, the Δε values of 70/30 and 70/30-D0.2T0.2 were slightly smaller than those of the other three 70/30 blends and XPLA due to their poorer compatibility between the two polymer components and heavier crosslinking. In addition, the Δσ values of the blends were all lower than that of XPLA because of the introduction of soft P(3HB-co-4HB). In addition, for the 30/70 blends, the Δε values of all detected blends and XP(3HB-co-4HB) were near 2%, which is much lower than that of neat P(3HB-co-4HB) (11.6%) owning to the limited mobility of the polymer chains. For four enhanced compatibility 30/70 blends, the entanglement among the polymer chains was increased when compared with the blank 30/70 blend, so that the Δσ values increased slightly. Consequently, it is both the flexibility of the polymer chain itself, the degree of chain entanglement and their cooperation that vary the stress–strain curves.

Table 3 Average Δσ and Δε values for the PLA/P(3HB-co-4HB) blends

Sample	Δε (%)	Δσ (MPa)	Sample	Δε (%)	Δσ (MPa)
PLA	—	—	P(3HB-co-4HB)	11.645 ± 2.012	0.060 ± 0.013
XPLA	0.713 ± 0.011	0.571 ± 0.036	XP(3HB-co-4HB)	2.263 ± 0.057	0.013 ± 0.005
70/30	0.710 ± 0.012	0.224 ± 0.032	30/70	—	0.022 ± 0.006
70/30-D0.05	0.714 ± 0.010	0.179 ± 0.030	30/70-D0.05	—	0.024 ± 0.005
70/30-D0.1	0.714 ± 0.009	0.252 ± 0.026	30/70-D0.1	2.343 ± 0.077	0.026 ± 0.007
70/30-D0.1T0.1	0.713 ± 0.013	0.187 ± 0.023	30/70-D0.1T0.1	2.770 ± 0.065	0.023 ± 0.005
70/30-D0.2T0.2	0.710 ± 0.010	0.233 ± 0.025	30/70-D0.2T0.2	—	0.023 ± 0.007

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Conclusions

Highly enhanced compatibilization of biosourced and biodegradable PLA and P(3HB-co-4HB) blends were prepared by reactive melt compounding. The miscibility, phase morphology, thermal behaviour, rheological and mechanical properties of the blends were investigated in detail using DMA, SEM, DSC and tensile tests. DMA and SEM results indicated that the miscibility of the two immiscible components was enhanced to different extents due to the introduction of branching and crosslinking. A larger shift towards each other for T_gs, a large reduction of the dispersed phase particle size and more significant increase of the interfacial adhesion of PLA and P(3HB-co-4HB) were achieved after crosslinking than that of branching and the blank blends. In addition, the branches decreased the crystallization ability of PLA, while crosslinking enhanced it on a large-scale. Moreover, the blends exhibited a remarkable improvement in their rheological properties in the melt state when compared with that of the blank PLA/P(3HB-co-4HB) blends. The transition from liquid-like to the solid-like viscoelastic behaviours at low frequencies demonstrated the formation of a crosslinking network. Upon increasing the DCP content, the blends showed increased yield tensile strength, modulus, and elongation at break. However, as DCP cooperated with TAIC, the elongation at break decreased because the crosslinking network limited the mobility of the polymer chains. The magnified images of the strain-hardening in the stress–strain curves for the blends showed step-growth with flexible entangled polymer chains and pulse-growth with rigid polymer chains. Furthermore, the SEM images of the fracture surfaces of the blends after the tensile tests presented fracture behaviour that changed from brittle fracture behaviour for neat PLA to ductile fracture behaviour for the blends. The plastic deformation of the matrix and debonding process were the two important ways for induced energy dissipation and led to a toughened blend. The important point is that fine toughness requires not only a high level of adhesion between the particles and matrix, but also enough molecular mobility, which is a crucial factor for yield stress and plastic flow.

Acknowledgements

The authors are grateful for financial support from the National Science Foundation of China (51021003).

Notes and references

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