2.1 Material

The isotactic poly(butene-1) (PB) is a commercial material with trade name PB0100M, which was kindly provided by Lyondell-Basell. The melt flow rate (MFR) is 0.4 g 10 min^–1 (190 °C, 2.16 kg).

Polypropylene (iPP), with an MFR of 0.3 g 10 min^–1 (at 230 °C, load 2.16 kg), was provided by Borealis Polyolefine GmbH under the tradename of BE50.

Poly(butylene succinate) (PBS) (PBE 003) was supplied by Natureplast. The MFR is ≈5 g 10 min^–1 (190 °C, 2.16 kg).

Poly(vinylidene fluoride) (PVDF), with tradename SOLEF 1012 was kindly provided by Solvay Specialty Polymers. The MFR is 0.5 g 10 min^–1 (at 230 °C, load 2.16 kg).

The HDPE (trade name FB3450, provided by Borealis Polyolefine GmbH) used in this study has an MFR of 0.3 g 10 min^–1 (at 190 °C, load 2.16 kg), and a density of 0.945 g cm^–3.

Polycaprolactone (PCL), tradename CAPA 6500, was supplied by Solvay and has an MFR of 7 g 10 min^–1 under a weight of 2.16 kg at 160 °C.

The blends were prepared by melt-mixing at 200 °C for 10 min at 100 rpm, using a Brabender-type internal mixer. The compositions of all the prepared samples were 20 wt.% dispersed phase and 80 wt.% continuous phase. For the sake of comparison, neat PB and neat PCL have been melt mixed in the Brabender using the same thermomechanical conditions. In the nomenclature of the samples, the dispersed phase was always indicated first. For instance, PB/iPP is the blend containing 20 wt.% PB and 80 wt.% iPP or PCL/HDPE is the blend containing 20 wt.% PCL and 80 wt.% HDPE.

2.2 Calorimetry (DSC)

The employed differential scanning calorimeter (DSC) was a TA Instrument DSC Q20 equipped with a refrigerating cooling system RCS40. It was calibrated with Indium and it worked under a 50 mL min^–1 flow of dry Nitrogen.

2.3 Scanning Electron Microscope (SEM)

The morphology of the cryogenically fractured surface of the different blends was investigated using a field-emission scanning electron microscope (Supra 40 VP model, Zeiss, Germany) at an accelerating voltage of 5 kV. The specimens were first cooled with liquid nitrogen and then fractured. A Polaron E5100 sputter coater was used to coat the sample with thin carbon layers.

3.1 Self-Nucleation of the Matrix and Non-Isothermal Crystallization of the Dispersed Phase in PB/Px Blends

Figure 2 presents micrographs of the cryogenically fractured surface of PB/iPP and PB/PVDF blends. The blends exhibit a sea-island morphology in which the minor PB phase is present in the form of droplets or micro-domains dispersed in the semicrystalline iPP and PVDF matrices. The morphology of each blend confirms the immiscibility of the studied systems. Figure S1 (Supporting information) reports the droplet size distributions of the two blends, the size of the dispersed droplets changes with the matrix type. Based on the original calculation method proposed in Chandrasekar's work,^[25] the sizes of dispersed micro-domains are reported in Table S1, Supporting Information. An increase in the diameter of the droplets is observed when changing the matrix from iPP to PVDF, which should be attributed to the differences in the melt-viscosity ratio and/or interfacial tension between the two components.^[26]

Results of the DSC standard cooling and heating scans for neat and blended components are shown in Figure 3. The peak crystallization and melting temperatures of neat PB are 72.7 and 116.5 °C, corresponding to the crystallization and melting of Form II. The larger endothermic and exothermic peaks are attributed to the crystallization and melting of iPP and PVDF, instead. A minor decrease in crystallization temperature (T_c) for PB in both blends with respect to the neat polymer was observed in Figure 3a. In fact, this phenomenon was already observed by Wang et al.^[27] who proposed that the nucleation of the finely dispersed PB droplets does not occur homogeneously, but at the interface with the iPP matrix, when the content of PB decreases below 25 wt.%. However, as the content of PB increases to 30 wt.% or more, the crystallization temperature of blended PB could be above that of neat PB, the reason for this being the migration of active heterogeneities from the iPP matrix to the PB phase during melt mixing. This explanation for blended systems are based on the previous works of Pracella et al.^[28-30] It should be noted that the slight decrease in T_c for PB (≈1 °C), although close to the precision of the DSC measurement, can be thus attributed to a loss of nucleating impurities of the dispersed phase toward the matrix.

In the subsequent DSC heating scans of Figure 3b, a new melting peak at ≈92 °C appears in the PB/iPP blend, indicative of the formation of Form I’.^[31] At a higher temperature, an exothermic signal is observed, indicating a recrystallization process, i.e., Form I’ crystals melt and recrystallize into Form II crystals. For PB/PVDF blend, the minor decrease of T_m for PB (≈6 °C) in comparison with the neat polymer should be still compatible with the melting of Form II, which crystallized at lower temperatures with respect to the neat PB.

Figure 4 shows DSC cooling and heating scans obtained after self-nucleating the iPP phase at different T_s values for the PB/iPP blend. This protocol can be used to study the effect of matrix crystallization on the dispersed phase. Possible changes in the thermal behavior of PB can be easily visualized.

The employed seeding temperatures allowed us to observe all of the typical self-nucleation domains of the iPP phase, as suggested by Müller et al.^{[16, 32]} By applying T_s in the range 175–200 °C to completely erase the crystalline thermal history, both crystallization and melting traces are unchanged, corresponding to Domain I. In this Domain, crystallization is dominated by heterogeneous nucleation produced by impurities or pre-existing nucleating heterogeneities. By lowering T_s to the range of 172–170 °C, a gradual increase in T_c values upon decreasing T_s is obtained, a behavior that corresponds to Domain II, indicating that a melt memory process is present. This additional nucleation occurs from the self-nuclei present in the non-isotropic melt.^{[24, 33]} The beginning of Domain III is detected when a melting shoulder appears to the right of the main melting peak in the heating scan, as produced by the fusion of the unmolten crystals annealed during the holding time in step 3 of the SN thermal protocol (see Figure 4b).

In Figure 4c, the behavior of dispersed PB crystallization within the specific thermal protocol (SN) for iPP matrix in PB/iPP blend is reported. A slight increase in T_c of PB phase in the blend in the cooling run is found as well, indicating an increase in crystallization kinetics. However, no significant variations in the melting behavior of the PB phase are noticeable during the final heating scans. In fact, the phenomenon has been found in similar blend systems,^{[10, 34]} in which the interaction of the two phases at the interface is proposed. Actually, epitaxy of PB with iPP has been reported in thin films.^[21] A geometrical relationship between the lamellar thickness of the matrix and dispersed phases are required for epitaxial growth of one phase on top of the other.^{[19, 20, 35-39]} In this case, surface-induced epitaxial crystallization has probably occurred in the PB/iPP blend, as suggested by the increased crystallization temperature of the PB dispersed phase when the T_c of the iPP matrix is increased via the self-nucleation protocol.

The effect of the self-nucleation on the PB and PVDF phases in the PB/PVDF blend is shown in Figure 5. As we described for the iPP phase in the PB/iPP blend, different SN domains for PVDF can also be highlighted, and are represented in different colors (red for Domain I, blue for Domain II, and green for Domain III) in Figure 5a,b. A slight shift toward higher temperatures of the crystallization peak of the PB phase in the blend can also be detected for T_s temperatures corresponding to Domain II, i.e., at 175 °C (See Figure 5c). This finding also suggests that PB nucleation occurs at the interface with the PVDF matrix. However, the existence of a specific epitaxial relationship between the two polymers is not demonstrated in the literature.

The T_c values of the matrix phases are reported in Figure 6a as a function of the matrix T_s values for both studied blends. The borders between the three characteristic domains are indicated as vertical lines. It is clear that no variation of T_c was found for the iPP and PVDF phases while in Domain I. On the contrary, a significant increase of T_c of both matrices are recorded in Domain II, meaning the crystallization kinetics of the matrix is greatly enhanced by the presence of self-nuclei.

Figure 6b shows the parallel evolution of T_c of the dispersed PB phase as a function of the corresponding matrix T_s. For PB in these two blends, the increase in the crystallization kinetics starts from the SN Domain II of the corresponding matrix, due to the effect of the changing matrix T_c, causing the formation of thicker lamellar crystals in the matrix phase, which in turn favors PB droplet nucleation at the interface. The T_c of the dispersed phase continues to increase in the SN Domain III of the iPP phase for PB/iPP blend, but starts to decrease in the Domain III of the PVDF phase for PB/PVDF blend. Even though there is a decrease in the T_c in Domain III, it remains always higher than the T_c recorded in Domain I. Furthermore, the T_c of PB in the PB/iPP blend is always higher than that of PB in the PB/PVDF blend, independently of the considered SN Domain. Considering that the average size of PB droplets is 0.24 µm in iPP matrix and 1.37 µm in PVDF matrix, a lower T_c of PB in Domain I for PB/iPP blend would be expected, due to a more probable loss of nucleating impurities. However, what happens in practice is that the T_c of PB is always higher in the case of PB/iPP with respect to PB/PVDF blends. This finding can be attributed to the higher nucleating efficiency of iPP surface toward PB, likely due to the known epitaxial relationship,^[21] while for PVDF, an epitaxial mechanism is either lacking or the crystallographic mismatch between the two structures is more significant than that existing for the PB/iPP system. Moreover, the T_c of PB in the blends is lower than that of neat PB indicated by the horizontal line in Figure 6b (the complete self-nucleation behavior of the neat PB is reported in Figure S2, Supporting Information), suggesting that some nucleating impurities of neat PB are transferred to the matrices during melt blending.

3.2 Self-Nucleation of the Matrix and Non-Isothermal Crystallization of the Dispersed Phase in PCL/Px Blends

SEM micrographs of the cryofractured PCL/PB and PCL/PVDF blends are shown in Figure 7, while the blends of PCL with the other matrices are reported in Figure S3 (Supporting Information). Similar to the PB/Px case, the micrographs reveal a sea-island morphology, typical of immiscible blends. The histograms of the size distribution of PCL droplets in the different blends are shown in Figure S4 (Supporting Information), while the average number, volume diameter (d_n and d_v), and dispersity (D) are reported in Table S2 (Supporting Information). It is noteworthy that a wide variety of droplet sizes is produced, as changing the matrix in the blends, the average size of droplets increased from 0.16 to 16.22 µm for PVDF and PBS, respectively. The average size is instead more similar and in the range of 0.6–1.8 µm for PCL/iPP, PCL/PB, and PCL/HDPE blends.

Figure 8 shows the DSC cooling and heating curves of neat PCL and PCL/Px blends obtained at 10 °C min^–1. The peak crystallization and melting temperatures of neat PCL are 31.8 and 55.8 °C, respectively. For blended systems, the larger peaks are always attributed to the crystallization and melting of the matrix, accordingly to the blend composition. Comparing the crystallization peak of PCL in neat and blended systems, an increase in the PCL T_c value (reported in Figure 8a) is highlighted in all the blends except for PCL/PVDF and PCL/PB (≈1 °C below T_c of neat PCL). The increased nucleation rate of PCL droplets in the different blends can be attributed to interfacial nucleation with various efficiencies or to the migration of active nucleating heterogeneities from the matrix to the PCL phase.^{[10, 28-30, 34, 40]} To distinguish between the two effects, the self-nucleation protocol of the matrix phase will be employed in the following. The decreased T_c of PCL/PVDF can only be attributed to a loss of nucleating impurities by the dispersed phase toward the matrix, while the PVDF surface can still act as a nucleating substrate, although with low efficiency.

Figure S5a (Supporting Information) reports the DSC cooling scans after self-nucleation of the neat PCL at the indicated T_s, while Figure S5b (Supporting Information) shows the subsequent heating scans. In neat PCL and at T_s higher than 80 °C, no appreciable changes in the crystallization peak are observed, which is indicative of SN Domain I. Starting from 80 °C down to a low T_s of 60 °C, a clear increase in the crystallization temperature is recorded. In this range of temperatures, the thermal treatment applied during SN creates self-nuclei in PCL, which is typical of Domain II. Below 60 °C, more crystals undergo annealing and lamellar thickening during self-nucleation, and changes in the final heating scans occur. In particular, a small melting peak at temperatures higher than the main melting peak, as typical of SN Domain III, is not observed in this case. Instead, the full melting peak slightly shifts to higher temperatures with decreasing T_s.

Figure 9 presents DSC cooling and heating scans after self-nucleation of the HDPE matrix phase within the PCL/HDPE blend at the indicated T_s. At T_s > 138 °C, both crystallization and melting of the HDPE phase are invariant. In this temperature range of Domain I, the crystallization behavior is controlled only by high-temperature-resistant nucleating heterogeneities. Domain III starts at T_s values ≤ 128 °C. In this temperature range, the T_c is increasing with the decrease in T_s, however, the sample is only partially molten, and the unmolten HDPE crystal fragments experience annealing and become thicker, resulting in a second melting peak at higher temperatures in the final heating scan (Figure 9b). This indicates that, with the adopted self-nucleation temperature window of 2 °C, no Domain II can be highlighted, and a direct transition from Domain I to Domain III is observed, as typical for HDPE. Moreover, in the studied T_s temperature range, there is a negligible variation in the crystallization behavior of PCL droplets after SN of HDPE matrix (Figure 9c). However, it is noteworthy that the crystallization temperature of PCL droplets in the PCL/HDPE blend is ≈6 °C higher than that of the isotropic melt of neat PCL.

Another example is reported in Figure 10, which shows the DSC cooling and heating scans after self-nucleation of the PB phase within the PCL/PB blend. As we described previously, the three characteristic SN domains are distinguished depending on the T_s range: red for Domain I (190–180 °C), blue for Domain II (180–150 °C), and green for Domain III (150–130 °C) as proposed by Müller et al.^{[16, 32]} The crystallization behavior of dispersed PCL during SN of PB phase in PCL/PB blend is analogous to that of dispersed PB in PVDF or iPP matrix. More specifically, the crystallization peak of PCL in the PCL/PB blend moves toward higher temperatures, as the T_s of PB decreases (Figure 10c). For the lowest T_s values employed, a slight decrease of PCL T_c is observed. We note that the maximum increase in PCL crystallization temperature is ≈3 °C, thus revealing a minor surface nucleation effect in comparison to the HDPE substrate (≈7 °C). Analogous phenomena with self-nucleation of other matrices within PCL/iPP, PCL/PVDF, and PCL/PBS blends are shown in Figures S6–S8 (Supporting Information), respectively. All these systems, except for the PCL/HDPE blend, show a relationship between the dispersed phase's T_c and the matrix's self-nucleation temperature. The trend of increasing crystallization temperature of the dispersed phase with decreasing T_s of the matrix is thus believed to be a general observation for double semicrystalline immiscible blends, being observed for six systems in the present work and one blend in previous literature.^[10] This relationship strongly suggests the role of the interface in the nucleation of dispersed droplets. Notably, by comparing Figures 10 and Figure 10, and Figures S6–S8 (Supporting Information), distinct nucleation effects, i.e., different PCL T_c, are exhibited by the different matrices, pointing toward a certain nucleation efficiency scale between the various polymeric interfaces.

The self-nucleation effect of the matrix on PCL crystallization in the various blends is better grasped with the help of Figure 11. The self-nucleation of neat PCL is also reported, for the sake of comparison. An initial invariance of PCL crystallization temperature with T_s in Domain I of the matrix is observed for all the investigated systems. Then a typical increase of PCL T_c with decreasing T_s is clearly identified for several systems, although the extent of the variation is highly matrix-dependent. Such enhancement usually occurs within Domain II of the considered matrix, where the increase in T_c of the matrix corresponds to thicker lamellae as a nucleating template (see the indicated domains of HDPE and PB systems as an example). Interestingly, lower T_s values (in Domain III) might lead to lower PCL T_c in some systems. This observation is interpreted as the existence of a maximum lamellar thickness, due to the lower lamellar thickening rate with decreasing matrix T_s.

The data in Figure 11 also reveal a clear nucleating efficiency scale of the different matrices. In particular, considering for each blend the maximum PCL T_c, the nucleating efficiency decreases in the order: HDPE>iPP>PB>PBS>PVDF. Notably, the crystallization temperature of PCL/HDPE is higher than the one recorded for neat PCL self-nucleated at T_s,ideal, i.e., at the lowest temperature of Domain II. This implies that, according to Fillon et al. nucleating efficiency scale,^{[41, 42]} the efficiency value of HDPE nucleating PCL is above 100%, i.e., it is a case of supernucleation.^[43-45] Such high nucleating efficiency is most probably due to the known excellent crystallographic matching between PCL and HDPE, already demonstrated by Yan et al. in thin polymer films.^[22] The epitaxial crystallization mechanism also explains the second ranking of iPP surfaces, as a certain degree of matching between some given PCL and iPP crystal planes was also reported.^[24] In fact, our nucleation efficiency observation, i.e., HDPE>iPP, is in agreement with previous literature on thin films.^[24] For the rest of the substrates, no specific epitaxial relationship is known. However, the matrix self-nucleation methods clearly prove the occurrence of surface nucleation, whose efficiency can be attributed to a poorer lattice matching with PCL or simply non-specific PCL/substrate interactions. It is worth noting that, for the PCL/PVDF blend, although nucleation clearly occurs at the surface (see the increase of PCL T_c with decreasing T_s of PVDF), the crystallization temperature of the dispersed PCL phase is lower than that of the neat polymer. This can only indicate that some nucleating heterogeneities are transferred from the PCL phase to the PVDF matrix domain during melt mixing, similar to what occurred for the PB-based blends.

To offer a different view of the surface nucleation phenomenon, the T_c values for the dispersed phase are reported in Figure 12 as a function of the matrix T_c in both PB and PCL-based blends. For practically all the systems, the T_c of PB (Figure 12a) or of PCL (Figure 12b) increases linearly with the crystallization temperature of the various matrices, which in turn is increased by the effect of self-nucleation. Given that the lamellar thickness of the matrix crystals is defined by the undercooling, i.e., by the crystallization temperature, we proposed that the relationship between the two T_c represented in Figure 12 parallels the increase of droplets nucleation kinetics when in contact with thicker lamellae of the matrix polymer. We note that the increased lamellar thickness for self-nucleated iPP has been demonstrated already for the case of HDPE/iPP blends.^[10] The reason behind the enhanced crystallization kinetics of the dispersed phase when increasing the crystal dimension of the matrix is tentatively ascribed to a template nucleation model.^[35] Thicker lamellae offer a better nucleating substrate for forming dispersed phase nuclei, since the crystallizing polymer chain segments are less likely in contact with the amorphous part of the matrix phase morphology, thus experiencing a lower energy penalty and hence a lower nucleation energy barrier. A similar concept has been put forward in our work in the case of cross-nucleation between polymorphs of PB.^[38]

Also from the plots of Figure 12, the different nucleating efficiency of the various polymer matrices can be easily captured, by judging the maximum T_c of each series. The comparison with the crystallization of neat polymers is also straightforward: it can be seen that both PVDF and iPP have an anti-nucleation effect on PB dispersed phase, due to impurity transfer to the matrices. The same is true for the PVDF matrix in PCL/PVDF blend, while HDPE substrate in PCL/HDPE blends displays a super-nucleation effect, possibly due to the good epitaxial matching, as previously discussed.^{[22, 23]}