Novel Compatibilized Nylon-Based Ternary Blends with Polypropylene and Poly(lactic acid): Fractionated Crystallization Phenomena and Mechanical Performance.

This paper presents an investigation into the behavior and performance of blends of Nylon 6 (PA6), polypropylene (PP), and poly(lactic acid) (PLA), compatibilized with maleic anhydride-grafted PP (PP-g-MA). The mechanical performance of ternary PA6/PP/PLA blends was superior to that of binary PA6/PP blends because of the addition of PLA. Through blending with PLA, the tensile and flexural strength and modulus were enhanced, maintaining performance similar to that of neat PA6. Tensile performance was further enhanced through reactive compatibilization of the blends with PP-g-MA due to the improved homogeneity of the materials. Impact behavior of the blends was found to be highly dependent on morphology, and the toughening behavior was observed at certain blending ratios. In PA6/PP blends, fractionated crystallization behavior was investigated through differential scanning calorimetry, in which both PA6 and PP droplets were crystallized at supercooled states. This effect was highly influenced by the presence of the compatibilizing agent and its effect on the morphology of the dispersed phase. As the droplet size of the dispersed phase was decreased to submicron levels, the efficiency of heterogeneous nucleation was limited. Crystallization of PLA in the blend was poor, but PP-g-MA was found to have an impact on its rate of crystallization.

Introduction

The durability and mechanical and tribological properties of polyamides, especially Nylon 6 (PA6) and Nylon 6,6 (PA6,6), are the driving reasons behind their frequent adoption in applications requiring high strength and longevity, especially within the automotive industry. Polyamides are frequently blended with polyolefins or elastomers to improve the additional aspects, especially their resistance to moisture and their toughness. Blending of PA6 with polypropylene (PP) is a well-researched topic because of their complementary properties. Because of the immiscibility of these polymers, a compatibilization route is necessary to produce blends with desirable qualities. Compatibilization is frequently accomplished through reactive extrusion with maleic anhydride-grafted PP (PP-g-MA), which eases the interfacial tension between PP and PA6 through the formation of a graft copolymer.[1−8] Maleated compatibilizing agents are very commonly used in polyolefin systems for improving miscibility in multipolymer blends, as well as for promoting adhesion to dispersed filler materials.[9] In PA6/PP blends, this reactive compatibilization scheme yields a PP-g-PA6 graft copolymer via a reaction in the melt state between terminal diamine groups of the PA6 with anhydride groups in PP-g-MA.[10] Although the incorporation of PP into PA6 can improve its moisture resistance and reduce the overall cost, this comes at the expense of reduced strength and stiffness. This study aims to offset the loss of tensile properties through the incorporation of a biopolymer, poly(lactic acid) (PLA), to produce ternary polymer blends. PLA is a renewable, biobased polymer with excellent tensile and flexural strength and modulus; through ternary blending with PLA, it is hypothesized that mechanical properties close to those of neat PA6 can be achieved. The morphology and rheological behavior of these blends will be presented in a concurrent manuscript.

Crystallization in blends of semicrystalline polymers is influenced by the presence of nucleating sites, interaction between the polymer phases, and blend morphology. In binary and ternary blends such as the ones investigated in this study, the size of the dispersed droplets and the interaction between the polymer phases are heavily dependent on the presence of a compatibilizer. A concurrent study demonstrated, by scanning electron microscopy (SEM) and atomic force microscopy, that the incorporation of PP-g-MA in PA6/PP and PA6/PP/PLA blends reduced the droplet size of PP to below one micron in diameter and promoted the interaction and adhesion at the PA6/PP interface. Rheological investigation further confirmed that PP-g-MA enhanced the interaction between PA6 and the dispersed phases. Interesting crystallization behavior has been observed in several studies of PA6-based blends. When analyzing the crystallization of PA6/PP blends compatibilized with PP-g-MA, Lee and Yang observed that the crystallizing temperature (Tc) of the dispersed PP decreased in the presence of the compatibilizer.[7] Holsti-Miettinen et al. observed a similar depression of the Tc of PP through compatibilization of PA6/PP blends with maleated elastomers but found no such effect in the presence of PP-g-MA.[11] The present work studies how the addition of PP-g-MA affects the crystallization processes in both binary and ternary blends, as well as the influence of the droplet size and polymer–polymer interactions in between PA6, PP, and PLA. A wide range of blending ratios was employed to provide an in-depth understanding of how the mechanical properties and crystallization mechanics evolve as a function of blend morphology and blending ratios.

Results and Discussion

Crystallization Behavior

Fractionated Crystallization of Dispersed Phases in Binary Blends

Differential scanning calorimetry (DSC) was used to analyze the melting and crystallization behavior of the binary blends of PA6 and PP. The full description of the blending ratios and nomenclature is given in the experimental section. For reference, Figure S1 (Supporting Information) shows both the cooling (a) and the second heating curves of each of the neat polymers (b).

For both PA6 and PP, only the melting transition appears on heating in the range investigated. The beginning of melting of PA6 is around 190 °C, and the melt phase consists of two peaks, corresponding to the melting of the γ-crystalline phase at 215 °C and the α-crystalline phase at 222 °C.[12,13] The beginning of PP melting is around 130 °C with a peak maximum at 164 °C.

The cooling curves at all blending ratios of PA6/PP are observed in Figure a. The observed behavior is expected for crystallization of two immiscible polymers, as the crystallization events are independent and proportional only to the fraction of each respective polymer. This was further confirmed by the nearly constant crystallinity at all blending ratios of the respective PA6 and PP phases as shown in Table . The crystallinity (Xc) of the PA6 fraction is around 24–25% in all blending ratios, whereas the PP fraction is consistently twice the value, around 48–50% crystalline. The lone exception is observed in B-90, in which the crystallinity of the PP fraction was slightly reduced. As explained by Arnal et al., as well as Santana and Müller, PP crystallization is highly dependent on morphology when it is in the dispersed phase.[14,15] PP crystallinity is reduced when it is finely dispersed because of the limitation of heterogeneous nucleation as the PP domain size is reduced. This is corroborated by the fine dispersion of PP droplets observed via SEM in B-90. In this blend, PP droplets were on average around 5 μm in diameters, with some droplets observed as small as 1 μm (Figure a). It is also reported by the manufacturer that the grade of PP used in this study contains a nucleating additive.

Figure 1

Cooling curves of (a) uncompatibilized and (b) compatibilized PA6/PP binary blends.

Table 1

Degree of Crystallinity of PA6 and PP in Binary Blends Measured from DSC Second Heating Curve

	without PP-g-MA		with 5 wt % PP-g-MA
PA6 wt % in blend	PP Xc (%)	PA6 Xc (%)	PP Xc (%)	PA6 Xc (%)
100		24.6
90	39.0	24.6	20.2	25.3
80	51.1	25.0	33.7	24.7
70	48.3	25.3	41.8	24.6
60	49.1	25.1	45.5	25.5
50	48.4	25.6	47.7	29.5
40	50.8	24.7	47.1	28.4
30	48.6	25.0	50.2	23.8
20	47.4	24.4	47.2	32.8
10	47.8	24.2	47.8	37.4
0	48.6

Figure 2

SEM images of impact fractured surfaces of (a) B-90 at 5000× magnification and (b) BC-80 at 12 000× magnification.

This concept applies further to the crystallization of the PP fraction in compatibilized binary blends. As observed in Table , the crystallinity of the PP fraction is nearly constant when PP is the dominant phase in the blend. However, as the PP fraction in these blends decreases below 50% and becomes the dispersed phase as phase inversion occurs, its crystallinity is steadily reduced, with a minimum of 20% crystallinity in BC-90. Furthermore, a fractionated crystallization of PP was observed in the cooling curves of these materials (Figure ). The exothermic crystallization is shifted progressively lower as the PP fraction decreases and is separated into two distinct exothermic events.

Figure 3

DSC cooling curves showing fractionated crystallization of the PP phase. Curves are vertically scaled proportionally to the blend PP fraction.

A similar fractionated crystallization in PA6/PP blends compatibilized by PP-g-MA has been observed in several studies.[7,16,17] A double crystallization of the peak of PP in PA6-compatibilized blend was also observed by Lee and Yang.[7] Ikkala et al. and Tang et al. both support the explanation that the fractionated crystallization of PP is caused by the limitation of nucleating heterogeneities in the dispersed PP droplets as the particle is reduced below a critical volume.[16,17] In the context of an immiscible system containing numerous small dispersed domains, Frensch and Jungnickel adapted eqs –3 on the premise that the proportion of those domains which contain exactly n nucleating heterogeneities follows a Poisson distribution[18]in which M is the concentration of nucleating heterogeneities in the dispersed phase and VD is the average domain volume of the dispersed phase. It follows that the proportion of these domains which contains at least one nucleating heterogeneity can be expressed as

Thus, we can expect that as the size of the dispersed PP domains decreases, the fraction of these domains that can crystallize through heterogeneous nucleation is reduced. These domains may undergo crystallization at supercooled temperatures because of nucleation from less efficient heterogeneities or even homogeneous nucleation of the polymer.[14−17] Supercooling is a phenomenon in which a polymer may remain in a liquid state below its typical crystallization temperature. The crystallization observed in this study correlates strongly with the greatly reduced domain size of dispersed PP with the addition of PP-g-MA as observed by SEM in Figure b. Upon addition of PP-g-MA, the PP was very finely dispersed in droplets less than 1 μm in diameter.

Supercooling of the dispersed phase is also observed when PA6 is dispersed in the PP matrix. From Figure , we can clearly observe that in compatibilized binary blends in which PA6 is the dominant phase, the PA6 crystallizes at ∼190 °C, just as observed in the uncompatibilized blends. However, upon phase inversion beginning in BC-50, the PA6 crystallization exotherm at this temperature completely disappears (Figure b). At the first glance, it would appear that PA6 is remaining in an amorphous state; however, the subsequent second heating curves revealed PA6 melting enthalpies between 60 and 90 J/g, depending on the polymer content. While no cold crystallization was observed on heating, the high melting enthalpies measured should present a significant crystallization peak. The cooling curves in Figure show a single exotherm at ∼123 °C, and no cold crystallization is observed in the heating curve upstream from the individual melting, implying that PA6 is crystallizing simultaneously with PP crystallization. A typical heating and cooling curve of BC-30 demonstrating this effect is included in the Supporting Information (Figure S2). PA6 crystallization not only occurs but also increases as the PA6 fraction decreases (Table ). This is explained by the supercooled crystallization of small dispersed domains of PA6. As is demonstrated via SEM in the concurrent study, the size of the PA6 domains in these blends is reduced below 1 μm because of the improved miscibility facilitated by PP-g-MA grafting.

Further investigation of the derivative heat flow in this region confirms that this exotherm is a single-step process (Figure ) and that the enthalpy of these events is equal to the sum of the PP and PA6 melting enthalpy in the subsequent heating curve. Thus, the dispersed PA6 fraction remains in a supercooled state until fast crystallization is induced at the droplet interface as the surrounding PP phase crystallizes. A series of papers by Tol et al. comprehensively investigated the phenomena of supercooled crystallization of PA6 dispersed in blends with polystyrene (PS).[19−22] They observed similar depression of the crystallization of PA6 when it was finely dispersed in PS, with maximum supercooling achieved in compatibilized blends with submicron dispersed domains of PA6. In this study, they concluded that PA6 crystallization can be induced at supercooled temperatures either through homonucleation or by the vitrification of the surrounding PS phase. In this case, the vitrification of the polymers studied occurs below 60 °C, and hence this factor was not considered.

Figure 4

Cooling curve showing derivative heat flow with respect to temperature in compatibilized binary blends (a) BC-10, (b) BC-20, (c) BC-30, (d) BC-40, and (e) BC-50.

Observing the second melting endotherm, there is a clear difference in the polymorphism of PA6 crystallized at supercooled temperatures (PA6 dispersed phase) and the PA6 main phase. Figure reveals that the more stable α-form dominates the PA6 crystallized at supercooled temperatures (in this case, BC-30), whereas in BC-70 both the α-form and γ-form are observed.[13] It is also clear that the onset melting temperature of PA6 was lower when PA6 was finely dispersed. This trend was consistent through all tested blends, with PA6 exhibiting both crystalline forms when the phase was continuous but mainly the γ-form when it was dispersed; similar melting exotherms of finely dispersed PA6 can also be observed in the work of Ikkala et al.[11] This is in contrast to the findings of Tol et al., who observed that supercooled PA6 droplets in PS crystallized mainly in the γ-form as the droplet size was decreased.[22] This may be explained by the different methods of PA6 nucleation; Tol et al. observed in this paper that PA6 crystallized mainly through homogenous nucleation, whereas in the present study, the crystallization of PP induced PA6 crystallization.

Figure 5

PA6 melting during the second DSC heating curve of BC-70 and BC-30.

Crystallization in Ternary Blends

For several reasons, analysis of PLA melting and crystallization within the ternary blends proved to be more difficult than that in the binary blends. The melting peak of PLA was observed at 168 °C, which coincided very closely with that of PP (164 °C), making it challenging to differentiate the two events. In the cooling curves, differentiation of PLA crystallization events was made additionally difficult by the relatively low enthalpy of crystallization of PLA (14 J/g) compared to that of PP (107 J/g) and PA6 (65 J/g). In blends containing such small fractions of PLA (as low as 5 wt %), accurate measurements of the enthalpy of PLA transitions were challenging. Although PA6 and PP showed a fast crystallization, PLA crystallized slowly, with a broad, shallow crystallization peak at 96 °C during the cooling run. In fact, the majority of the melting enthalpy observed in neat PLA was due to cold crystallization during the heating run. At a cooling rate of 10 °C/min, the PLA crystallized incompletely because of its relatively slow rate of crystallization, leading to cold crystallization around 98 °C during the second heating as seen in Figure S1b.[23]

The crystallization of PA6 was incredibly consistent in all ternary blends, as seen in Table . Regardless of the proportion of PLA and PP, the crystallinity of the PA6 phase was around 24–25% in all ternary blends, just as in neat PA6. Because PA6 is the major, continuous phase in these systems and it crystallizes at a much higher temperature than PP and PLA, its crystallization is completely independent of the dispersed phases in the system. Just as in the binary blends, although PA6 is the major phase in the ternary blends, its crystallization was also unaffected by the addition of PP-g-MA as a compatibilizer. However, this was not the case for the dispersed phases.

Table 2

Degree of Crystallinity of PA6, Enthalpy of Melting of PP + PLA, and Enthalpy of Cold Crystallization of PLA in Ternary Blends As Measured by DSC

material	PA6 Xc (%)	PP + PLA Hm (J/g) (2nd heating)	PLA Hcc (J/g) (1st heating)	PLA Hcc(J/g) (2nd heating)
PA6	24.5
PLA		39.3	27.7	22.3
PP		92.3
T-90	24.4	47.2	11.8a	15.4a
T-80	24.6	51.4	16.7	16.6
T-70	25.4	52.9	18.1	16.0
T-60	24.6	55.5	19.6	15.5
TC-90	25.5	23.2	13.2a	5.4a
TC-80	25.3	32.9	15.7	3.9
TC-70	25.5	40.8	18.8	3.2
TC-60	24.4	47.2	20.4	3.3

These values have greater uncertainty because of the relatively small fraction of PLA in these blends.

As was observed in the preceding section, the addition of PP-g-MA depressed the crystallization temperature of the PP phase because of the greatly decreased droplet size (Figure ). In all ternary blends, the addition of PP-g-MA depressed the PP crystallization by 15–20 °C and resulted in a double peak. Although it would seem intuitive to attribute the two peaks to each of PP and PLA, this behavior was also observed in the binary blends. The crystallization peaks observed in Figure are caused almost entirely by the crystallization of PP because of the substantially higher enthalpy of crystallization observed in neat PP compared to neat PLA. Rather than a second peak due to the crystallization of PLA, the second peak is attributed to the crystallization of small PP droplets by less efficient nucleating heterogeneities. Table shows the combined melting enthalpy of PP and PLA in the ternary blends; whereas the two melting events could not be differentiated because of their close overlap, and some conclusions may still be drawn from the combined value. It has already been shown that the majority of this melting enthalpy comes from the PP phase. In the uncompatibilized blends, the melting enthalpy is fairly consistent, decreasing slightly as the fraction of PP and PLA is reduced. As observed in the binary blends, fractionated crystallization due to the addition of PP-g-MA decreases the crystallinity of the PP phase, and this effect increases as the PP volume fraction decreases (see Figure ). These results correlate strongly with the vast reduction in the PP droplet size observed by microscopic methods in the compatibilized ternary blends.

Figure 6

Depression and fractionation of PP crystallization in compatibilized ternary blends.

Though the melting and crystallization of PLA is clouded by the PP phase in this system, the faster PP crystallization during cooling allowed the observation of the cold crystallization of PLA separately; this can be observed clearly in the second heating curve of T-80 in Figure a. In the uncompatibilized ternary blends and neat PLA, cold crystallization of PLA was observed during both the first and second heating runs. The crystallization of PLA in these blends was not completed during injection molding nor was it completed on cooling at 10 K/min, resulting in a cold crystallization peak during the subsequent heating cycles, which was probably inherent to the nuclei formed at a lower temperature favoring the crystallization. As explained previously, most of the PLA crystals, determined by their melting enthalpy, were formed during their cold crystallization. Thus, we can expect that the PLA phase in the molded testing samples is poorly crystalline because of the rapid cooling experienced during molding. An interesting effect on PLA crystallization transition was observed after the addition of PP-g-MA to the ternary blends. Like the uncompatibilized blends, the compatibilized blends demonstrated cold crystallization of PLA during the first heating cycle, showing the presence of an incompletely crystallized PLA phase after molding. However, cold crystallization during the second heating was reduced by up to 85% compared to the first heating, as shown in Table . It is probable that the presence of the PP-g-MA compatibilizer increased the crystallization rate of PLA, causing the majority of PLA crystallization to complete during the cooling run. This effect was also observed by Akrami et al., using the same grade of PLA in blends with thermoplastic starch, compatibilized with a maleic anhydride-based additive.[24] They found that the addition of the compatibilizer reduced both the temperature and enthalpy of cold crystallization of PLA in the blends; an increased rate of PLA crystallization facilitated more crystal growth during the cooling run.

To further investigate this behavior, blends of TC-80 were prepared by varying the levels of PP-g-MA, that is, 0, 2.5, 5, and 7.5 wt %. Although the same effect was observed on cold crystallization in each blend regardless of the PP-g-MA concentration (see Figure a), the effect of compatibilizer concentration on the crystallization of PP from the melt was observed, as shown in Figure b. With increasing concentration of PP-g-MA, the crystallization temperature of PP is reduced further. Although the depression of PP crystallization has been explained previously as an effect of the droplet size, the significant amount of grafting used here likely decreased the mobility of PP, which hindered the molecular chain movement required to allow its crystallization. Such a behavior would also induce a shift of the crystallization to a lower temperature.

Figure 7

Cooling (A) and second heating (B) of (a) T-80 and TC-80 with (b) 2.5, (c) 5.0, and (d) 7.5 percent PP-g-MA.

Mechanical Properties

Tensile and Flexural Behavior

The full mechanical properties of the blends investigated in this study are provided in a table in the supporting figures (Table S1). Starting with the uncompatibilized binary blends, it is evident that with increasing addition of PP to PA6, there is a consistent decrease in the tensile and flexural strength and modulus of the material. These decreased properties follow the same trend as the reductions observed in multiple studies of PA6/PP binary blends by Agrawal et al.,[3] Huber et al.,[1] and Pal and Kale.[25] In general, this behavior is typical and expected for an immiscible polymer blend such as the one studied. The PA6 used in this work has a relatively high tensile strength and modulus of 75.2 MPa and 2.67 GPa, respectively, compared to the tensile strength and modulus of the PP (38.9 MPa and 1.89 GPa). As the PP fraction of the blend is increased, we naturally expect the strength and stiffness of the blend to be reduced in comparison to neat PA6. As proven by the SEM analysis in the concurrent study, the binary blend transitions to a cocontinuous morphology at a blending ratio of 60 PA6/40 PP. This morphology change can be correlated with the yield elongation behavior observed for these blends. The yield elongation of the blends decreases from neat PA6 to B-90 and so on until cocontinuous morphology is reached in B-60, at which point the minimum yield elongation of the binary blends is observed. A yield elongation value similar to that of B-60 is observed in B-70 and B-50, before a steady increase from B-40 to B-10 as the PP loading increases toward 100%. This change in the yielding behavior is explained by the shift from a continuous PA6 phase with dispersed PP domains, to a cocontinuous system, and finally to a continuous PP phase with dispersed PA6 domains. This morphology evolution is in very close agreement with that observed by Willis et al. in binary PA6/PP blends prepared at the same blending ratios.[26] Their study similarly demonstrated that phase inversion occurred at a 60 PA6/40 PP blending ratio.

A major change in the yielding behavior as well as the tensile and flexural strength of the blends was observed when 5% PP-g-MA compatibilizer was added to the binary systems. In the compatibilized blends, the yield elongation remains nearly the same as that of neat PA6 from BC-90 to BC-50, after which the yield elongation increases toward that of neat PP as the PP fraction approaches 100%. The break elongation is also improved dramatically in all compatibilized binary blends compared to their uncompatibilized counterparts, with especially high ductility observed in the PP-dominated blends (BC-30, BC-20, and BC-10). The increased ductility of the compatibilized binary blend is a result of the improved adhesion and continuity between the PA6 and PP phases.[4,11,25] The tensile and flexural behavior was also increased with the addition of PP-g-MA to the binary blends. In BC-90 through BC-50, the tensile and flexural strength were dramatically improved compared to those of B-90 through B-50. Figure shows the flexural strength of both the compatibilized and uncompatibilized binary blends compared to the upper bound of the ideal strength of the blends, as determined by the rule of the mixture. We can see that the uncompatibilized blends substantially underperform compared to the ideal values; however, the flexural strength of compatibilized blends is very close to the the ideal values, whereas PA6 is the continuous phase. The improved strength in the compatibilized blends is a result of the reduced size of the dispersed phase and superior interaction at the interface due to the graft copolymer. Improved stress transfer from the continuous PA6 phase to the dispersed PP allows the blend to behave more like a homogenous material, resulting in strength closer to that of an ideal blend. There is a drop-off in strength in the compatibilized blends when the volume fraction of PA6 drops below 50% and PP becomes the continuous phase. Even when compatibilized, PA6 did not reinforce PP efficiently as the dispersed phase.

Figure 8

Flexural strengths of binary blends based on experimental data and ideal values determined by the rule of mixture.

The addition of PLA in the ternary blends resulted in an increase in the tensile strength and modulus of the blends. Starting with the blends T-90 through T-60, there is a vast increase in the tensile strength and stiffness compared to that of the equivalent binary blends (B-90 through B-60). The neat PLA has a tensile strength similar to that of neat PA6, and a superior Young’s modulus, flexural strength, and flexural modulus. The high strength and stiffness of the added PLA counteracts the low strength and stiffness of the PP, such that the uncompatibilized ternary blends outperform their equivalent compatibilized binary blends in these aspects. Relative to the equivalent binary blends, the ductility of the uncompatibilized ternary blends is improved. As discussed in the morphology section, the PA6, PP, and PLA phases are mutually immiscible. This leads the dispersed domains of PLA and PP to remain smaller than those in the binary blends, rather than coalescing into larger droplets. The smaller size of the dispersed domains in the uncompatibilized ternary blends results in a higher elongation at yield and break than in the equivalent binary blends, and closer to that of neat PA6.

As in the binary blends, the addition of PP-g-MA compatibilized the ternary blends and resulted in improved tensile and flexural strength compared to the uncompatibilized blends, as well as improved elongation at break. Because of the addition of the PLA phase, the stiffness of the compatibilized ternary blends meets or exceeds that of neat PA6, even in BC-60. The flexural strengths of the compatibilized ternary blends are compared to those of the compatibilized binary blends in Figure . The addition of PLA improves the flexural and tensile strength and modulus of the polymer blend, resulting in ternary blends with superior properties to the homologous binary blends. The strength of the uncompatibilized ternary is improved compared to the binary blends but drops rapidly once the PA6 fraction drops below 80% because of the increased size of the dispersed phase, which easily debond from the PA6 when the material is deformed. As is observed in the morphology investigation of the concurrent study, the PLA and PP domains in the compatibilized ternary blends are better distributed and have improved interaction with the continuous PA6 phase. This leads to improved stress transfer and limits debonding between the PA6 and the dispersed phases, thus improving the strength and ductility of the material. Furthermore, the strength and stiffness of each of the compatibilized ternary blends far exceeds that of the equivalent binary blends.

Figure 9

Flexural strengths of ternary blends compared to those of homologous binary blends.

Impact Behavior

Blending with PP and PLA affected the fracture morphology of PA6, as demonstrated in Figure . SEM analysis at a larger scale (500×) magnification revealed the fracture morphology of neat PA6 (Figure e). Because of the pullout and facile debonding of the dispersed phases in the uncompatibilized ternary blends (Figure a,c), the fracture morphology is very different from that of neat PA6. However, compatibilization of the binary and ternary blends leads to a fracture surface that was nearly indistinguishable from that of neat PA6 (Figure b,d). The impact behavior of the binary blends is complex, whereas at high PP fractions the impact strength trends toward that of neat PP and in blends with low PP fractions (B-90 through B-70), the impact strength exceeds that of neat PA6. A similar trend in the Izod impact strength of binary PA6/PP blends was observed by González-Montiel et al.[5] This increase in impact strength occurs only in the blends that exhibited the sea-island morphology with a continuous phase of PA6 and well-dispersed PP domains, which suggests that PP dispersion is the driving factor behind the toughening effect.[11] PA6 is well-known as a notch-sensitive polymer, having very low crack propagation energy.[27] It is possible that the softer dispersed PP domains within the PA6 hindered crack propagation through the blend via deflection of the crack tip.

Figure 10

SEM micrographs at 500× magnification of impact fracture surfaces: (a) B-80, (b) BC-80, (c) T-80, (d) TC-80, and (e) neat PA6.

The toughening effect observed in these binary blends was increased further with the addition of the PP-g-MA compatibilizer. The impact strength was maximized in BC-80, which demonstrated a 62% increase in strength compared to neat PA6. Improved adhesion between the PP and PA6 phases further enhanced the toughening effect that was already present in B-90, B-80, and B-70. The addition of PLA decreased the impact strength of the ternary blends because of the relatively brittle nature of PLA, though at high PA6 weight fractions the impact strength remained close to that of neat PA6. With the addition of the compatibilizer, the impact strength was increased very slightly, except in BC-90, which became more brittle. In this blend, the PP content was replaced entirely by PP-g-MA, negating the toughening effect observed as a result of the dispersed droplets of pure PP within the PA6. Up to 30% of the PA6 fraction can be replaced with PLA and PP before the impact strength deviates significantly from that of neat Nylon.

Conclusions

This study investigates the crystallization and mechanical performance of PA6/PP/PLA blends, with further investigation of morphology and rheological behavior established in a concurrent manuscript. In this section, blend morphology and compatibilization were demonstrated to have a significant impact on blend crystallization. In particular, the fine dispersion achieved through compatibilization caused fractionated crystallization of dispersed PP droplets and confined the crystallization of PA6 at supercooled temperatures, which was nucleated at the droplet interface by PP crystallization. The incorporation of PLA as a ternary blend component had the intended effect of enhancing the tensile strength, flexural strength, and stiffness of the blends. Through compatibilization of the blends with PP-g-MA, the mechanical properties were improved further because of the improved dispersion of the minor phases, adhesion at the polymer interfaces, and better stress transfer from the dispersed droplets to the continuous PA6 phase. Overall, mechanical properties very close to that of neat PA6 were achieved in compatibilized blends in which up to 30% of PA6 by weight was replaced by PP and PLA. This study presents an innovative method for improving classical polyamide–polyolefin blends through blending with a biopolymer. The work presented here opens further opportunities for research into reinforced ternary blend systems, increased bio-based content in polyamide blends, process optimization of reactive extrusion for ternary blends, and finally investigations into alternate compatibilization schemes for this polymer system.

Experimental Section

Materials

The Nylon 6 used in this work was Ultramid B27E, a low-viscosity extrusion grade polymer supplied by BASF (Germany), referred to as PA6. Ingeo Biopolymer 3251D, an injection-grade PLA from NatureWorks (USA), was used as the PLA phase. PP 1120H supplied by Pinnacle Polymers (USA) was used as the PP phase. Fusabond P353 from DuPont (USA) was selected as the maleated PP compatibilizing agent. The maleation grade of this compatibilizer has been determined to be in the range of 1.4–1.9%.[1,28] The materials are further detailed in the Supporting Information (Table S2).

Blend Preparation

All materials were dried overnight at 80 °C to eliminate the moisture content prior to processing. Blends were prepared via melt blending in a Leistritz Micro-27 (Germany) twin-screw extruder in corotation configuration. The extruder was operated at 100 rpm, with all blend components premixed and fed from a single feeder at a feed rate of 7 kg/h. The configuration of the 12 heating zones set at ∼250 °C is shown in Table S3 in the Supporting Information. This heating profile was selected to promote reactive extrusion at high temperature in the screw-mixing zones, whereas limiting the time spent at high temperature to reduce the degradation of the PLA phase. The reduced temperature at the die also limited the expansion of the extrudate and improved the processability of the material. Residence time was approximated at 90 s from timing the throughput of the machine with a stopwatch.

All blends designed in this study are detailed in Table with the following naming convention: a prefix of B for “binary”, T for “ternary”, with C for “compatibilized”, with the number after the dash corresponding to the weight fraction of PA6 in the blend. For example, BC-80 is a compatibilized blend of PA6 and PP, with a PA6/PP ratio of 80/20. Blends were selected to investigate the properties of compatibilized and uncompatibilized PA6/PLA/PP ternary blends in parallel with equivalent PA6-PP blends. In the first section, a full blending profile of binary blends of PA6/PP was developed to analyze the morphology development and mechanical properties, as well as the role of the compatibilizer in the binary blend. In the second section, ternary blends of PA6/PLA/PP were produced following parallel blending ratios to the binary blends. In all compatibilized blends, 5% PP-g-MA by weight was added, sacrificing an equivalent amount of the PP fraction. Because of the reduced melt strength of the uncompatibilized ternary blends, the ternary blends were only investigated to a blending ratio of 60:20:20 PA6/PLA/PP. Homologous blends were formulated to contain equivalent weight fractions of PA6. Each of the neat polymers was also extruded and injection-molded in the same fashion to give baseline data for comparison.

Table 3

Blending Ratios for Ternary and Binary Blends of PA6, PLA, and PP

			component (wt %)
		name	PA6	PLA	PP	PP-g-MA
	neat polymers	neat PA6	100	0	0	0
		neat PLA	0	100	0	0
		neat PP	0	0	100	0
binary	uncompatibilized	B-90	90	0	10	0
		B-80	80	0	20	0
		B-70	70	0	30	0
		B-60	60	0	40	0
		B-50	50	0	50	0
		B-40	40	0	60	0
		B-30	30	0	70	0
		B-20	20	0	80	0
		B-10	10	0	90	0
	compatibilized	BC-90	90	0	5	5
		BC-80	80	0	15	5
		BC-70	70	0	25	5
		BC-60	60	0	35	5
		BC-50	50	0	45	5
		BC-40	40	0	55	5
		BC-30	30	0	65	5
		BC-20	20	0	75	5
		BC-10	10	0	85	5
ternary	uncompatibilized	T-90	90	5	5	0
		T-80	80	10	10	0
		T-70	70	15	15	0
		T-60	60	20	20	0
	compatibilized	TC-90	90	5	0	5
		TC-80	80	10	5	5
		TC-70	70	15	10	5
		TC-60	60	20	15	5

Differential Scanning Calorimetry

Melting and crystallization behavior of the blends was studied using a TA Instruments (USA) DSC Q200 differential scanning calorimeter. Temperature and enthalpy calibration were done using an an indium standard. All tests were run on samples from injection-molded bars to see the crystalline state of the material in the as-molded condition. The samples were prepared in 40 μL aluminum pans of about 6–8 mg. The experiments were run under a N2 atmosphere (60 mL min–1). A heat/cool/heat method was used with the following profile: heating ramp of +10 °C/min from room temperature to 250 °C (+20 °C above the melting temperature of the most thermally resistant polymer-PA6) to remove thermal history; cooling ramp of −10 °C/min to −70 °C; and heating ramp of 10 °C/min to 250 °C. For crystallinity calculations, the melting enthalpy for 100% crystalline samples (Hf) was 190 J/g for PP[29] and for 230 J/g for PA6.[30] Crystallinity (Xc) calculations were normalized to the weight fractions (wi) of the individual polymers as seen in eq

Mechanical Properties

Tensile and flexural properties were measured using an Instron (USA) 3382 Universal Testing Machine. Tensile tests were conducted with a 50 mm/min crosshead speed following the ASTM D638-14 method for type IV specimens. Flexural tests were conducted with a 14 mm/min crosshead speed following the ASTM D790-15 method. For each blend, five specimens were tested to determine the tensile and flexural properties. Ideal blend properties were determined volumetrically based on the upper bound according to the rule of mixture.

Izod impact testing was conducted using a TMI (USA) Monitor impact tester. Samples were notched and prepared following the ASTM D256-10 method for Izod impact resistance. At least five samples were prepared and tested for each blend.

Microscopic Analysis of Fracture Surfaces

The morphology of impact fracture surfaces was investigated by SEM. The SEM used for this application was Phenom ProX (Netherlands) set to an accelerating voltage of 10 kV.