**Development of Injection-Molded Polylactide Pieces with High Toughness by the Addition of Lactic Acid Oligomer and Characterization of Their Shape Memory Behavior**

**Diego Lascano 1,2, Giovanni Moraga 1, Juan Ivorra-Martinez 1, Sandra Rojas-Lema 1,2, Sergio Torres-Giner 3,\*, Rafael Balart 1, Teodomiro Boronat <sup>1</sup> and Luis Quiles-Carrillo <sup>1</sup>**


Received: 11 November 2019; Accepted: 11 December 2019; Published: 14 December 2019

**Abstract:** This work reports the effect of the addition of an oligomer of lactic acid (OLA), in the 5–20 wt% range, on the processing and properties of polylactide (PLA) pieces prepared by injection molding. The obtained results suggested that the here-tested OLA mainly performs as an impact modifier for PLA, showing a percentage increase in the impact strength of approximately 171% for the injection-molded pieces containing 15 wt% OLA. A slight plasticization was observed by the decrease of the glass transition temperature (Tg) of PLA of up to 12.5 ◦C. The OLA addition also promoted a reduction of the cold crystallization temperature (Tcc) of more than 10 ◦C due to an increased motion of the biopolymer chains and the potential nucleating effect of the short oligomer chains. Moreover, the shape memory behavior of the PLA samples was characterized by flexural tests with different deformation angles, that is, 15◦, 30◦, 60◦, and 90◦. The obtained results confirmed the extraordinary effect of OLA on the shape memory recovery (Rr) of PLA, which increased linearly as the OLA loading increased. In particular, the OLA-containing PLA samples were able to successfully recover over 95% of their original shape for low deformation angles, while they still reached nearly 70% of recovery for the highest angles. Therefore, the present OLA can be successfully used as a novel additive to improve the toughness and shape memory behavior of compostable packaging articles based on PLA in the new frame of the Circular Economy.

**Keywords:** PLA; OLA; impact modifier; shape memory; packaging applications

#### **1. Introduction**

Polylactide (PLA) is a linear thermoplastic biodegradable polyester that can be obtained from starch-rich materials by fermentation to give lactide, which is polymerized at the industrial scale by ring-opening polymerization (ROP) [1]. PLA is a sound candidate to substitute some plastic commodities such as polypropylene (PP) or polystyrene (PS) in packaging applications [2–6], electronics [7], automotive [8], agriculture [9], textile, consumer goods, 3D printing applications [10–13], biomedical devices [14,15], pharmaceutical carriers [16,17], etc. The main advantage of PLA over other biopolymers is its relatively low cost and overall balanced properties and processability [18,19],

resulting in compostable articles [20]. In 2018, PLA production represented 10.3% of the worldwide production capacity of bioplastics, reaching nearly 220,000 tons/year and it is estimated a growth around 60% by 2023 [21].

Although PLA is a very versatile biopolymer, it results in extremely brittle materials with very low elongation at break and low toughness [22]. Many research works have been focused on overcoming or, at least, minimizing the intrinsic brittleness of PLA materials. One possible strategy is copolymerization with long aliphatic monomers, as suggested by Zhang et al. [23]. This is the case of the new type of polyester amide (PEA) copolymers developed by Zou et al. [24] combining poly(L-lactic acid) (PLLA) and poly(butylene succinate) (PBS) flexible segments. A similar approach was developed by Lan et al. [25] based on PLA-*co*-PBS copolymers. Nevertheless, the most widely used approach is blending with other biopolymers due to its cost effectiveness. For instance, Garcia-Campo et al. [26,27] reported interesting toughening effects on ternary blends composed of PLA, poly(3-hydroxybutyrate) (PHB), and different rubbery polymers such as PBS, poly(butylene succinate-*co*-adipate) (PBSA) or poly(ε-caprolactone) (PCL). Recently, Sathornluck et al. [28] developed binary blends of PLA with epoxidized natural rubber (ENR), which can positively contribute to improving toughness due to the effect of the rubber phase finely dispersed in a brittle PLA matrix. Su et al. [29] and Zhang et al. [30] have also reported different approaches to overcome the intrinsic brittleness of PLA by adding different contents of PBS using reactive compatibilizers. Fortelny et al. [31] recently reported, however, that it is possible to obtain "super-toughened" PLA formulations by blending PLA with poly(ε-caprolactone) (PCL) without any compatibilizer but using the appropriate PCL particle size. Addition of modified or unmodified particles is another way to overcome the intrinsic PLA brittleness. For instance, Wang et al. [32] improved the PLA toughness by means of silanized helical carbon nanotubes (CNTs). Similarly, Li et al. [33] reported the positive effect of cellulose nanofibers (CNFs) and a Surlyn® ionomer to improve the interfacial interaction between CNFs and PLA, leading to a remarkable increase in toughness. The work carried out by Gonzalez-Ausejo et al. [34] reported a clear improvement in toughness by using sepiolite nanoclays as gas barrier and compatibilizer in PLA/poly(butylene adipate-*co*-terephthalate) (PBAT) blends.

Even though copolymerization and blending are some of the most used procedures to overcome the intrinsic brittleness of PLA, plasticization is another interesting approach. Plasticizers have been widely used in PLA formulations to reduce fragility by decreasing the glass transition temperature (Tg). Plasticizers contribute to an increase in ductile properties, such as elongation at break, but mechanical resistant properties are also negatively decreased. Therefore, the use of plasticizers not always provides enhanced toughness since it depends on both ductile and resistant properties. For instance, Tsou et al. [35] improved toughness of PLA by using an adipate ester plasticizer. It has also been reported the positive effect of combining two plasticizers: one solid plasticizer, namely poly(ethylene glycol) (PEG), and a liquid plasticizer derived from soybean oil [36]. Indeed, many research works have been focused on using environmentally friendly plasticizers, which can contribute to improving toughness without compromising the overall biodegradability of PLA. In this regard, Kang et al. [37] described the usefulness of cardanol obtained from cashew nutshell liquid (CNSL) to obtain a 2.6-fold increase in toughness over neat PLA. Carbonell-Verdu et al. [38] developed a new series of dual plasticizer and compatibilizer additives derived from cottonseed oil subjected to epoxidation and/or maleinization as an environmentally friendly solution instead of using conventional epoxy-styrene acrylic oligomers. Other similar multi-functionalized vegetable oils have demonstrated good compatibilizing effects on PLA [39,40] and also on its blends with other biopolymers [41]. Ferri et al. [42] recently reported a remarkable increase in toughness of neat PLA by using maleinized linseed oil (MLO) as a bio-based plasticizer, which yielded a slight plasticization that overlapped with chain extension, branching, and some cross-linking. Tributyl citrate (TBC) has also been proposed as an effective plasticizer for PLA by Notta-Cuvier et al. [43], who reported the synergistic effect of TBC in PLA formulations containing halloysite nanotubes (HNTs).

Oligomers of lactic acid (OLAs) can also provide plasticization to PLA and PLA-based materials [44]. As reported by Burgos et al. [45], OLAs can contribute to a significant decrease in Tg, though their role as impact modifiers require further research. Therefore, this study focuses on the effect of the addition of a new type of OLA on PLA pieces prepared by injection molding. In particular, the effect of varying the OLA content on the mechanical, thermal, thermomechanical properties and the morphology of PLA was reported. Finally, the shape memory behavior of the pieces was analyzed and related to the OLA dispersion within the PLA matrix.

#### **2. Experimental**

#### *2.1. Materials*

PLA was supplied by NatureWorks LLC (Minnetonka, MN, USA) as IngeoTM Biopolymer 6201D. This PLA grade contains 2 mol% D-lactic acid. It was supplied in pellet form and it had a density of 1.24 g cm−<sup>3</sup> whereas its melt flow index (MFI) was 20 g/10 min, measured at 210 ◦C and 2.16 kg. OLA was kindly supplied by Condensia Química S.A. (Barcelona, Spain) as Glyplast OLA 2. It was provided in a liquid form with a viscosity of 90 mPa.s at 40 ◦C. According to the manufacturer, it has an ester content >99%, a density of 1.10 g cm−3, a maximum acid index of 2.5 mg KOH g−1, and a maximum water content of 0.1 wt%.

#### *2.2. Manufacturing of OLA-Containing PLA Pieces*

As PLA is very sensitive to moisture, the biopolyester pellets were dried at 60 ◦C for 24 h. The OLA content varied in the 0–20 wt% at weight steps of 5 wt%. The terminology used for the formulations was "PLA-OLA x%", where x represents the weight fraction of OLA in PLA. All the compositions were compounded in a twin-screw co-rotating extruder Coperion ZS-B 18 (Stuttgart, Germany) equipped with a main hopper in which the PLA pellets were fed and a side feeder for liquids to feed OLA. The liquid feeder was heated at 50 ◦C during compounding to decrease the OLA viscosity and allow efficient mixing. The rotating speed was set to 150 rpm while the temperature profile was modified according to the formulation as shown in Table 1. These temperatures were selected due to the change in viscosity produced by the OLA addition in order to optimize the processing conditions for each formulation. After extrusion, the different strands were cooled in air and then pelletized using an air-knife unit.


**Table 1.** Temperature profile of the seven barrels in the twin-screw extruder during the compounding of the polylactide (PLA)/oligomer of lactic acid (OLA) formulations.

The compounded pellets were stored in an air-circulating oven at 60 ◦C for 24 h to avoid moisture gain. Standard samples for characterization were thereafter obtained in an injection molding machine Meteor 270/75 from Mateu & Solé (Barcelona, Spain). The temperature profile of the four barrels was programmed as indicated in Table 2. Similar to the extrusion process, it was necessary to optimize the temperature profile for each formulation due to the drop in the melt viscosity after the OLA addition.


**Table 2.** Temperature profile of the four barrels and maximum pressure in the injection-molding machine during the manufacturing of polylactide (PLA)/oligomer of lactic acid (OLA) pieces.

#### *2.3. Mechanical Characterization*

Mechanical properties of the OLA-containing PLA pieces were obtained in tensile conditions as indicated by ISO 527-1:1996. A universal testing machine ELIB 30 from S.A.E. Ibertest (Madrid, Spain) was used. At least six different samples were tested at room temperature using a load cell of 50 kN. The cross-head speed rate was set to 10 mm min−1. As recommended by the standard, the tensile modulus (Et), tensile strength at break (σb), and elongation at break (%εb) were determined and averaged. The toughness was estimated using the Charpy method following the guidelines of ISO 179-1:2010 with a 6-J pendulum from Metrotec (San Sebastián, Spain). Unnotched rectangular samples with dimensions 80 <sup>×</sup> 10 mm<sup>2</sup> and a thickness of 4 mm were tested. The impact strength was obtained from testing five different samples and calculated as the absorbed-energy per unit area (kJ m<sup>−</sup>2). The Shore D hardness was determined in a 673-D durometer from J. Bot Instruments (Barcelona, Spain) according to ISO 868:2003. Ten different measurements were collected from randomly selected zones and various samples were tested to obtain the average values.

#### *2.4. Microscopy*

The morphology of the fracture surfaces of the OLA-containing PLA pieces was studied by field emission scanning electron microscopy (FESEM) after the impact tests. A ZEISS Ultra 55 FESEM microscope from Oxford Instruments (Abingdon, UK) operating at 2 kV was used to collect the FESEM images at 1000×. To avoid electrical charging during observation, samples were previously coated with an ultrathin gold-palladium alloy in an EMITECH SC7620 sputter-coater from Quorum Technologies Ltd. (East Sussex, UK) in an argon atmosphere.

#### *2.5. Thermal Characterization*

Thermal properties of the OLA-containing PLA pieces were obtained by differential scanning calorimetry (DSC) and thermogravimetric analysis (TGA). DSC characterization was performed on a Mettler-Toledo DSC calorimeter DSC821e (Schwerzenbach, Switzerland). To carry out the DSC runs, small specimens of each composition with an average weight of 5–7 mg were placed in standard aluminum crucibles with a total volume of 40 mL and sealed with a cap. Then, the samples were subjected to a three-step temperature program. A first heating cycle from 30 ◦C to 200 ◦C was followed by a cooling step down to −60 ◦C and, finally, a second heating cycle from −60 ◦C to 350 ◦C was applied. The heating and cooling rates were set to 10 ◦C min−<sup>1</sup> whereas the selected atmosphere was nitrogen at a flow-rate of 30 mL min<sup>−</sup>1. The resultant DSC curves allowed obtaining Tg, the cold crystallization peak temperature (Tcc), and the melting peak temperature (Tm). Besides, the cold crystallization (Δ*Hcc*) and melting enthalpies (Δ*Hm*) were obtained from the integration of the corresponding peaks. The maximum degree of crystallinity (%χcmax) was calculated as indicated in Equation (1):

$$\% \chi\_{\text{cumx}} = \frac{\Delta H\_{\text{m}}}{\Delta H\_{\text{m}}^{0}} \cdot \frac{100}{w} \tag{1}$$

where *w* (g) stands for the weight fraction of PLA and Δ*Hm<sup>o</sup>* (J g<sup>−</sup>1) represents the theoretical melting enthalpy of a fully crystalline PLA polymer, which is close to 93.7 J g−<sup>1</sup> [46–49].

The effect of the OLA addition on the thermal stability of PLA was studied by TGA in a TGA/SDTA 851 thermobalance from Mettler-Toledo (Schwerzenbach, Switzerland). The temperature sweep was scheduled from 30 ◦C to 700 ◦C at a heating rate of 10 ◦C min−<sup>1</sup> in air atmosphere. Samples with an average weight comprised between 5 and 7 mg were placed on standard alumina crucibles and sealed with the corresponding cap. The onset degradation temperature, which was assumed at a weight loss of 5 wt% (T5%), and the maximum degradation rate temperature (Tdeg) were obtained. All DSC and TGA tests were run in triplicate to obtain reliable results.

#### *2.6. Viscosity Characterization*

To study the influence of the OLA addition on the viscosity of PLA, cylindrical disks sizing 40 mm diameter and 5 mm thickness were manufactured by hot-press molding in a BUEHLER SimpliMet 1000 (Lake Bluff, IL, USA) at a temperature of 165 ◦C and a pressure of 7.5 ton. Parallel-plate oscillatory rheometry (OR) was conducted in an AR-G2 rheometer from TA Instruments (New Castle, DE, USA) to obtain the evolution of the complex viscosity (|η*\**|) as a function of the angular frequency. The selected isothermal temperature was 200 ◦C and the angular frequency varied in the 100–0.01 rad s−<sup>1</sup> range. The maximum shear strain (γ) was set to 1% and the tests were carried out in air atmosphere in triplicate.

#### *2.7. Thermomechanical Characterization*

The effect of the OLA addition on the thermomechanical behavior of PLA was carried out by dynamic mechanical thermal analysis (DMTA) in a DMA 1 from Mettler-Toledo (Schwerzenbach, Switzerland) working in single cantilever flexural conditions. Rectangular samples with dimensions of 10 <sup>×</sup> 7 mm<sup>2</sup> and a thickness of 4 mm were subjected to a dynamic heating program from <sup>−</sup>30 ◦C to 130 ◦C at a constant heating rate of 2 ◦C min−1. The maximum deflection in the free edge was set to 10 μm and the selected frequency was 1 Hz. The storage modulus (E*'*) and the dynamic damping factor (*tan* δ) were collected as a function of increasing temperature.

The dimensional stability of the OLA-containing PLA pieces was studied by thermomechanical analysis (TMA) in a Q400 thermomechanical analyzer from TA Instruments (New Castle, DE, USA). The applied force was set to 0.02 N and the temperature program was scheduled from −30 to 120 ◦C in air atmosphere (50 mL min<sup>−</sup>1) at a constant heating rate of 2 ◦C min<sup>−</sup>1. The coefficient of linear thermal expansion (CLTE) of the PLA pieces, both below and above Tg, was determined from the change in dimensions versus temperature.

#### *2.8. Characterization of the Shape Memory Behavior*

The flexural method was used to evaluate the extension of the shape memory behavior. To this end, samples were compression-molded into sheets in the hot-press molding with a thickness of 0.4–0.5 mm. Two different parameters were then calculated to analyze the shape memory behavior in flexural conditions, namely the shape memory recovery (*Rr*) and stability ratio (*Rf*) [50,51]. The procedure is described as follows. In the first stage, the sheet specimens were deformed at a particular angle θ*f*. For this, the samples were heated at 65 ◦C and forced to adapt between two aluminum sheets with different angles (15◦, 30◦, 60◦, and 90◦) to form a sandwich structure: aluminum/PLA sheet/aluminum. The sandwich was then clamped to allow a permanent deformation to the desired angle. Thereafter, the sandwich was cooled down to 14 ◦C to achieve its permanent shape after a slight temporary recovery of θ*t*. Finally, the specimen was heated above its Tg in an air circulating oven at 65 ◦C for 3 min. After this, the final angle, that is, θ*f*, was measured. At least three different sheets were tested for each composition to obtain reliable values. The same procedure has been used to characterize the shape recovery behavior in several polymer systems [52,53]. The values of *Rf* and *Rr* were calculated using the following equations:

$$\% \mathcal{R}\_f = \frac{(\pi - \theta\_t)}{\left(\pi - \theta\_f\right)}.\tag{2}$$

$$\% \mathcal{R}\_r = \frac{\left(\pi - \theta\_f\right) / \left(\pi - \theta\_r\right)}{\left(\pi - \theta\_f\right)}.\tag{3}$$

#### **3. Results and Discussion**

#### *3.1. E*ff*ect of OLA on the Mechanical Properties of PLA*

Table 3 gathers the main results obtained after the mechanical characterization of the PLA pieces with different OLA contents. As expected, the progressive addition of OLA led to lowering σ<sup>b</sup> values from 64.6 MPa, for the neat PLA, down to 37.4 MPa, for the PLA pieces containing 20 wt% OLA. This decreasing tendency was almost linear as it can be seen in the table for the other compositions. This tendency on mechanical strength is the typical a plasticizer produces on the base polymer. Some other OLA additives have demonstrated a similar effect on mechanical properties by increasing remarkably ductility, mainly in the film form, as reported by Burgos et al. [45]. In the latter work, it was reported an increase in ε<sup>b</sup> from 4% to 315% with an OLA content of 25 wt%, but it is worthy to note the OLA previously used was designed for plasticization of PLA films. Such dramatic increase in the ε<sup>b</sup> was not observed when using OLA in this work since the primary use of this type of OLA was to improve impact strength, as indicated by the supplier and it will be discussed further.

**Table 3.** Mechanical properties of the polylactide (PLA) pieces with different weight contents of oligomer of lactic acid (OLA) in terms of: tensile modulus (Et), strength at break (σb), elongation at break (εb), impact strength, and Shore D hardness.


Regarding mechanical ductility, the neat PLA piece showed a very low value of 7.9% and the addition of OLA did not promote an increase in εb, but a slight decrease down to values of 5%. It is not usual that a plasticizer promotes a decrease in ductility since the typical effect of a plasticizer is a decrease in the tensile resistant properties (σ<sup>b</sup> and Et) and an increase in ductile properties (εb). Nevertheless, it has been reported that some plasticizers promote a clear plasticization that is detectable by a decrease in Tg while no improvement in ductility occurs. This atypical behavior was reported by Ambrosio-Martín et al. [54] in PLA films blended with different synthesized OLAs. A ε<sup>b</sup> value of 5.25% was reported for neat PLA, while the addition of 25 wt% of a purified OLA yielded a ε<sup>b</sup> of 2.52%. It was also reported a slight increase in Et and an apparent decrease in σb, in a similar way as obtained in this work. It was concluded that, although there is clear evidence of the mechanical plasticization of OLA-containing PLA films, they were not more deformable, which is also in agreement with the work performed by Courgneau et al. [55]. Concerning Et, the neat PLA piece was characterized by a value of nearly 2.2 GPa and the values remained in the 2.2–2.4 GPa range after the addition of OLA. Thus, the main effect of this type of OLA on the tensile mechanical properties was a remarkable decrease in σb, which representative for some plasticization, but also a slight decrease in εb. Interestingly, as it can also be seen in Table 3, the addition of OLA successfully increased the impact strength of PLA. Neat PLA showed an impact strength of 25.7 kJ m<sup>−</sup>2, which indicates a brittle behavior, and the only addition of 5 wt% OLA provided a slight increase in the impact strength to 30.4 kJ m<sup>−</sup>2. Nevertheless, the most remarkable changes were obtained for OLA additions of 10 wt% and 15 wt%, showing impact strength values of 54.2 kJ m−<sup>2</sup> and 69.7 kJ m−2, respectively. Therefore, the PLA piece with 15 wt% OLA presented the maximum impact strength with a percentage increase of approximately 171%

with regard to the neat PLA. It is also worthy to mention that the PLA piece containing 20 wt% OLA showed a decrease in impact strength in comparison with the other OLA-containing PLA pieces, thus suggesting certain OLA saturation in the PLA matrix. In this regard, Fortunati et al. [56] reported the use of isosorbide diester (ISE) as plasticizer for PLA. It was observed plasticizer saturation at 20 wt% ISE and this was attributed to a limitation of Tg decrement. In addition, a noticeable decrease in ε<sup>b</sup> was observed once the plasticizer saturation was achieved. Furthermore, Ferri et al. [57] reported the plasticization of PLA by fatty acid esters, observing a remarkable decrease in impact strength above 5 parts per one hundred parts (phr) of PLA. Accordingly, a remarkable decrease in Tg was attained for contents of up to 5 phr whereas, above this, the Tg values did not change in a noticeable way, corroborating the relationship between the impact and thermal properties. Regarding hardness, one can observe that the Shore D values remained nearly constant after the OLA addition, showing values in the 78–82 range. Therefore, the most important feature this OLA can potentially provide to the mechanical properties of PLA is a remarkable improvement in impact strength while the elasticity can be slightly improved and ductility reduced. This particular mechanical behavior could be ascribed to an increase in the sample crystallinity and also to the presence of soft domains of OLA dispersed within the PLA matrix, simultaneously improving impact strength and reducing flexibility.

As previously indicated, one of the most widely used strategies to improve toughness in PLA-based formulations is blending with rubber-like polymers such as PCL [58], PBS [29,59], or PBAT [60,61]. In these immiscible blends, the energy absorption is related to presence of finely dispersed rubber-like small polymer droplets embedded in the brittle PLA matrix. In some cases, a synergistic effect can be found when different reactive or non-reactive compatibilizers are used. In this work, OLA has the same chemical structure than PLA, thus leading to miscibility without the need of compatibilizers. In this regard, Burgos et al. [45] have reported the similarity between the solubility parameters of both PLA and OLA, which plays an essential role in miscibility. According to this, Figure 1 gathers the FESEM images corresponding to the fracture surfaces from the impact tests of the PLA pieces with the different OLA loadings. Figure 1a shows the fracture surface of the neat PLA piece. As one can observe in this micrograph, the surface was smooth with multiple microcracks presence, which is an indication of a brittle behavior. As opposite, Figure 1b shows that the fracture surface morphology of the PLA piece with 5 wt% OLA changed noticeably. Phase separation could not be detected due to the high chemical affinity between PLA and OLA and the microcracks were not observed but, in contrast, macrocracks were produced. Therefore, the presence of OLA seems to inhibit microcrack formation and growth and, therefore, the cracks could grow to a greater extent thus leading to a rougher surface that is responsible for higher energy absorption during impact. Figure 1c,d show the FESEM images corresponding to the fracture surfaces of the PLA pieces with 10 wt% and 15 wt% OLA, respectively. In these images, the above-mentioned effect was more intense then showing rougher surfaces that are related to enhanced energy absorption. Finally, Figure 1e shows that the PLA piece containing 20 wt% OLA presented a similar fracture surface than the other pieces and, despite there was a clear loss of toughness, its morphology did not allow identifying phase separation.

**Figure 1.** Field emission scanning electron microscopy (FESEM) images of the fracture surfaces of the of the polylactide (PLA) pieces with different weight contents of oligomer of lactic acid (OLA): (**a**) 0 wt%; (**b**) 5 wt%; (**c**) 10 wt%; (**d**) 15 wt%; and (**e**) 20 wt%. Images were taken at 1000× and scale markers are of 10 μm.

#### *3.2. E*ff*ect of OLA on the Thermal and Rheological Properties of PLA*

Figure 2 shows a comparative graph with the characteristic DSC thermograms during first heating corresponding to the neat PLA piece and the PLA pieces with different OLA loadings. Table 4 summarizes the main thermal values obtained from the thermograms. One can observe that Tg of neat PLA was located close to 63 ◦C. Then, the PLA sample cold crystallized indicating that the biopolyester chains could not crystallize in the injection mold. The process of cold crystallization was characterized by a peak at 109.8 ◦C. Finally, the melting process was defined by a peak temperature of 170.9 ◦C in which the whole crystalline fraction melted. The effect of the OLA addition on the thermal properties was remarkable. Concerning Tg, a clear decreasing tendency can be observed. In particular, Tg was reduced after the OLA addition down to 50.8 ◦C, thus indicating plasticization. Ambrosio-Martín et al. [54] reported a decrease in Tg with addition of OLA to PLA films from 60 ◦C to 27.7 ◦C, concluding that this fact is directly related to the mechanical properties of OLA and it depends on the synthesis procedure [62]. In the present work, the Tg value of PLA decreased after the addition of OLA but this reduction was much lower than other reported to other plasticizers. For instance, Ljungberg et al. [63] reported a Tg decrease of 30 ◦C with 15 wt% addition of different plasticizers

such as triacetin, tributyl citrate (TBC), triethyl citrate (TEC), acetyl tributyl citrate (ATBC), and acetyl triethyl citrate (ATEC).

Δ Δ

χ

**Figure 2.** Differential scanning calorimetry (DSC) thermograms corresponding to the polylactide (PLA) pieces with different weight contents of oligomer of lactic acid (OLA).

**Table 4.** Thermal properties of the polylactide (PLA) pieces with different weight contents of oligomer of lactic acid (OLA) in terms of: glass transition temperature (Tg), cold crystallization temperature (TCC), cold crystallization enthalpy (Δ*HCC*), melting temperature (Tm), melting enthalpy (Δ*Hm*), and degree of crystallinity (χ*cmax*).


Furthermore, the added OLA induced an internal lubricating effect that shifted the cold crystallization of PLA to lower temperatures due to an increase of chain mobility. Then Tcc lowered to values in the range of 96–101 ◦C with the different OLA loadings. Other authors suggested a specific nucleating effect provided by the short length OLA molecules, which are more readily to pack the PLA macromolecular structure thus favoring the cold crystallization process [64]. In addition, a small and broad exothermic peak was seen in the PLA sample processed with OLA, particularly noticeable at the lowest OLA contents. This exothermic peak is related to a pre-melt crystallization just before melting. In this regard, one can consider that the presence of OLA promoted the formation of different crystallites [63]. As reported by Maróti et al. [65] and also Maiza et al. [66], this peak has been observed in neat PLA depending on the heating rate and the applied thermal cycle. One can also observe a slight decrease in the Tm value with the increasing OLA content. Similar findings have been reported by Burgos et al. [62] in PLA films with different OLAs.

In addition to the characteristic values of Tg, Tcc, and, Tm, the enthalpies corresponding to the cold crystallization and melting processes, that is, Δ*Hcc* and Δ*Hm*, respectively, were collected. The maximum degree of crystallinity, that is, χ*cmax*, which does not consider the amount of crystals formed during cold crystallization, was 35.6% for the neat PLA. Then, χ*cmax* increased up to values of around 50% for the compositions containing 5–15 wt% OLA, while slightly lower values of crystallinity were obtained for the composition containing 20 wt% OLA, that is, 45.8%. This increase in crystallinity can be related to the plasticizing effect of OLA, as earlier reported by Burgos et al. [62] in PLA formulations with 15 wt% of different OLAs. This latter study also reported a decrease in the Tm value of approximately 5 ◦C.

Regarding thermogravimetric characterization, Figure 3 shows the mass versus temperature (Figure 3a) and the first derivative (DTG) versus temperature (Figure 3b) curves for all the PLA pieces. The main results of the thermal decomposition of PLA-OLA blends are summarized in Table 5. One can observe that the neat PLA was much more thermally stable than the toughened PLA formulations with the different OLA loadings. As the OLA content increased, the characteristic TGA curves in Figure 3a shifted to lower temperatures, thus indicating a decrease in thermal stability. DTG curves were very useful to determine the maximum degradation rate temperature (Tdeg), which was seen as peaks in Figure 3b. It can be seen in the graph that there was a clear decreasing tendency of Tdeg with increasing OLA content. Furthermore, the residual mass for all PLA formulations with OLA was almost the same, being below 1 wt%.

**Figure 3.** (**a**) Thermogravimetric analysis (TGA) and (**b**) first derivate thermogravimetric (DTG) curves corresponding to the polylactide (PLA) pieces with different weight contents of oligomer of lactic acid (OLA).

**Table 5.** Main thermal parameters of the polylactide (PLA) pieces with different weight contents of oligomer of lactic acid (OLA) in terms of: onset temperature of degradation (T5%), degradation temperature (Tdeg), and residual mass at 700 ◦C.


Therefore, OLA induced a remarkable reduction in the thermal stability of PLA, as similarly described by Ambrosio-Martín et al. [54]. In this regard, Burgos et al. [62] also reported individual TGA characterization of different OLAs with their corresponding thermal degradation parameters. A variation in the onset degradation temperature was observed from 179 ◦C to 214 ◦C. The lower thermal stability was related to the lower Tg values. Moreover, it was reported a maximum degradation

rate temperature ranging from 259 ◦C to 291 ◦C. These characteristic degradation temperatures are remarkably lower than those of the neat PLA herein studied. As the OLA content in PLA pieces increased, both the T5% and the Tmax values showed a clear decreasing tendency. This decrease was more pronounced in the case of T5%, which varied from 336.1 ◦C, for the neat PLA piece, to 254.8 ◦C, for the PLA piece containing 20 wt% OLA.

As indicated previously, OLA provided a lubricating effect, which could potentially lead to a subsequent decrease in the viscosity. Figure 4 shows the effect of the different OLA loadings on the |η*\**| values of the PLA sheets as a function of the angular frequency. As it can be seen, the effect of the angular frequency on the complex viscosity was not highly pronounced but it was possible to detect a decreasing tendency of |η*\**| with increasing the OLA content. This confirms that OLA lowers the viscosity of the PLA melt during processing and it can therefore potentially act as a processing aid. In fact, as indicated previously, it was necessary to adjust the thermal profile for optimum processing when the OLA content was modified.

**Figure 4.** A comparative plot of the complex viscosity (|η*\**|) of the polylactide (PLA) sheets with different weight contents of oligomer of lactic acid (OLA) at a constant temperature of 200 ◦C as a function of increasing angular frequency.

#### *3.3. E*ff*ect of OLA on the Thermomechanical Properties of PLA*

The results described above indicated a definite improvement in the PLA toughness by using OLA as an impact modifier. Mechanical characterization showed a decrease in mechanical strength while ductility was also slightly reduced. DMTA allows characterization of mechanical properties in dynamic conditions (sinusoidal applied stress) as a function of a heating cycle. Figure 5 shows the DMTA curves for the neat PLA piece and the PLA pieces containing different loadings of OLA. The results of the thermomechanical properties obtained by DMTA are summarized in Table 6. The variation of the E' values of the neat PLA (Figure 5a) showed a dramatic drop between 50 ◦C and 70 ◦C, which is representative of the α-relaxation process of the PLA chains as the glass transition region was surpassed. In particular, a three-fold decrease in E' was observed. As the OLA loading increased, the E' curves shifted to lower temperatures thus indicating a decrease in Tg, as previously observed by DSC analysis. In particular, the E' values decreased from 1500 MPa, for the neat PLA piece, to 1197 MPa, for the

PLA piece blended with 20 wt% OLA at 30 ◦C. As reported by other authors, both nucleating agents and plasticizers play a crucial role in the DMTA behavior of PLA [67,68]. Although the characteristic E' curves showed a decrease in Tg, more accurate values can be obtained by determining the peak maximum of *tan* δ, as observed in Figure 5b. Neat PLA showed a Tg of 68.2 ◦C and the Tg values decreased progressively as the OLA loading increased, reaching a minimum value of 49.4 ◦C for the PLA piece containing 20 wt% OLA. These results are in total agreement with the above-described results obtained during the DSC characterization.

**Figure 5.** Evolution as a function of temperature of the (**a**) storage modulus (E') and (**b**) dynamic damping factor (*tan* δ) of the polylactide (PLA) pieces with different weight contents of oligomer of lactic acid (OLA).

**Table 6.** Main thermomechanical parameters of the polylactide (PLA) pieces with different weight contents of oligomer of lactic acid (OLA) in terms of: storage modulus (E') measured at 30 ◦C and 70 ◦C, glass transition temperature (Tg), and coefficient of linear thermal expansion (CLTE) below and above Tg.


\* Tg was obtained as the peak maximum of the dynamic damping factor (*tan* δ).

Another attractive thermomechanical property is the effect of temperature on the dimensional stability of the PLA-based materials with different OLA loadings. TMA is a very useful technique to determine the CLTE values, which is a crucial property related to dimensional stability in terms of temperature exposition. Table 6 gathers these coefficients for the neat PLA piece and the PLA pieces containing different amount of OLA. As one can observe in the table, two different CLTE values were determined, one corresponding to the slope below Tg and another one corresponding to that above Tg. Plasticization can be observed by seeing the CLTE both below and above Tg. A slight increasing tendency was obtained, thus, indicating more ductility. The CLTE below Tg changed from 79.9 μm m−<sup>1</sup> K−<sup>1</sup> to 91.6 μm m−<sup>1</sup> K<sup>−</sup>1. The maximum change was therefore 11.7 μm m−<sup>1</sup> K<sup>−</sup>1, which is a very narrow range, typical of values below Tg then indicating excellent dimensional stability. Above Tg, the maximum change was 29.1 μm m−<sup>1</sup> K<sup>−</sup>1, which is in accordance with the typical plastic thermomechanical behavior above Tg.

#### *3.4. E*ff*ect of OLA on the Shape Memory Behavior of PLA*

Initially, the shape memory behavior was studied qualitatively by introducing the sheet specimens into a glass tube at room temperature, as observed in Figure 6a, and remained inside for 5 min to retain the shape. Then, the crimped PLA sheets were immersed in a water bath at 70 ◦C, above the biopolyester's Tg, and allowed to recover their shape. As can it be seen in Figure 5b, flat sheet shapes were obtained for the PLA materials containing >10 wt% OLA loadings in a short period of 4–10 s, therefore giving support to the significant effect of OLA on the shape memory behavior of PLA. Similar results, under the same conditions, were reported in the development of poly(l-lactide-*co*-ε-caprolactone) (PLACL), which showed recovery times of approximately 20 s [69]. Another study was focused on PLA/thermoplastic polyurethane (TPU) blends in which the recovery times at 70 ◦C were very similar to those obtained in this work, that is, 7–12 s [50].

**Figure 6.** Photographs of the qualitative study of the shape memory recovery capacity of the polylactide (PLA) sheets with different weight contents of oligomer of lactic acid (OLA): (**a**) initial deformation of the sheets by introducing them into a glass tube and (**b**) recovered shape of the sheets after heating at 70 ◦C.

In addition to the qualitative characterization shown above, a quantitative study was carried out to evaluate the influence of the OLA impact modifier on the shape memory behavior of PLA. Figure 7 shows the evolution of *Rr* for different angles as the OLA content was increased. As it can be seen in the plot, the neat PLA sheet showed limited shape recovery properties and, obviously, it highly depended on the deformation angle. This shape memory recovery was close to 77% for a deformation angle of 90◦ and it was remarkably lower with more aggressive deformations. For instance, PLA could only recover 55.5% when the initial deformation angle was 15◦. As the OLA loading increased, the ability of PLA to recover its initial flat shape (angle of 180◦) increased considerably. It is worthy to note the effect of the addition of 20 wt% OLA, which yielded an almost constant increase in *Rr* of approximately 20% for all the tested angles. Therefore, for this OLA loading, the shape recovery ability of PLA remarkably improved thus leading to an exciting shape memory behavior. As reported by Leonés et al. [70], a good shape memory behavior can be attained in electrospun PLA-based fibers containing different OLA loadings in the 10–30 wt% range.

**Figure 7.** Evolution of the percentage of shape memory recovery (%*Rr*) of polylactide (PLA) sheets with different weight contents of oligomer of lactic acid (OLA) at different initial deformation angles: 15◦, 30◦, 60◦, and 90◦.

The Rf value is representative for the dimensional stability below Tg, after an initial deformation. As it can be seen in Table 7, Rf was very high for all the systems, thus indicating excellent dimensional stability after the first deformation. In general, as the deformation angle was lower, for instance 90º, the stability ratio was higher, showing recovery values over 98%. Alternatively, if the initial deformation angle was very aggressive, for instance 15◦, a slight decrease in the stability ratio can be observed down to values of nearly 85%. Anyway, these stability ratios can be considered high for all compositions and angles. In this regard, Jing et al. [50] reported lower stability ratios in PLA/TPU blends, which showed less dimensional stability than the materials developed herein. Therefore, one can conclude indicating that this type of OLA is also an exceptional additive for shape memory recovery as shown by the high %Rr and %Rf values obtained.


**Table 7.** Variation of the percentage of stability ratio (%Rf) after different initial deformation angles (θf) for the polylactide (PLA) sheets with different weight contents of oligomer of lactic acid (OLA).

#### **4. Conclusions**

The positive effect of OLA to improve PLA toughness was evaluated in this study. The addition of 15 wt% OLA provided an increase in the impact strength from 25.7 kJ m−<sup>2</sup> to almost 70 kJ m−2, thus showing an extraordinary effect on toughness. Furthermore, the OLA impact modifier also provided a mechanical plasticization as observed by a decrease in the tensile strength though a slight reduction in ductility was also noticed. This plasticizing effect was observable by DSC in which the characteristic Tg of the neat PLA was reduced from 63.3 ◦C to 50.8 ◦C for the PLA piece with 20 wt% OLA. Moreover, since OLA was less thermally stable than PLA, a decrease in the onset degradation temperature, reported as T5%, was observed as the OLA loading increased. Nevertheless, the T5% value corresponding to the PLA piece with the highest OLA loading was still high, that is, 254.8 ◦C, which successfully allows processing these blends without thermal degradation. Another exciting feature that this type of OLA can provide to PLA is the improvement of its shape memory behavior. In particular, the shape memory recovery parameter, that is, Rr, was very high compared to other PLA-based blends or plasticized PLA systems, thus showing the extraordinary effect of this OLA on the shape memory recovery. As a general conclusion, the here-studied OLA additive represents an interesting technical and environmentally friendly solution to improve the intrinsic brittleness of PLA and it also contributes to somewhat plasticization that allows enhanced shape memory recovery properties. The resultant toughened PLA materials can be of high interest for the development of compostable packaging articles, such as food trays and films, or disposable articles, such as cutlery and straws.

**Author Contributions:** Conceptualization, R.B. and D.L.; methodology, G.M., S.R.-L.; validation, L.Q.-C., T.B. and S.T.-G.; formal analysis, J.I.-M. and R.B.; investigation, D.L.; data curation, G.M., D.L., J.I.-M., and S.R.-L.; writing—original draft preparation, L.Q.-C. and T.B.; writing—review and editing, R.B., S.T.-G.; supervision, T.B., L.Q.-C., and D.L.; project administration, S.T.-G. and R.B.

**Funding:** This research work was funded by the Spanish Ministry of Science, Innovation, and Universities (MICIU) project numbers RTI2018-097249-B-C21 and MAT2017-84909-C2-2-R.

**Acknowledgments:** L.Q.-C. wants to thank Generalitat Valenciana (GVA) for his FPI grant (ACIF/2016/182) and the Spanish Ministry of Education, Culture, and Sports (MECD) for his FPU grant (FPU15/03812). D.L. thanks Universitat Politècnica de València (UPV) for the grant received through the PAID-01-18 program. S.T.-G. is recipient of a Juan de la Cierva contract (IJCI-2016-29675) from MICIU. S.R.-L. is recipient of a Santiago Grisolía contract (GRISOLIAP/2019/132) from GVA. J.I.-M. wants to thank UPV for an FPI grant PAID-01-19 (SP2019001). Microscopy services of UPV are acknowledged for their help in collecting and analyzing the microscopy images. Authors also thank Condensia Química S.A. for kindly supplying Glyoplast OLA 2.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### *Article* **E**ff**ect of Almond Shell Waste on Physicochemical Properties of Polyester-Based Biocomposites**

#### **Marina Ramos 1, Franco Dominici 2, Francesca Luzi 2, Alfonso Jiménez 1, Maria Carmen Garrigós 1,\*, Luigi Torre <sup>2</sup> and Debora Puglia 2,\***


Received: 15 March 2020; Accepted: 1 April 2020; Published: 6 April 2020

**Abstract:** Polyester-based biocomposites containing INZEA F2® biopolymer and almond shell powder (ASP) at 10 and 25 wt % contents with and without two different compatibilizers, maleinized linseed oil and Joncryl ADR 4400®, were prepared by melt blending in an extruder, followed by injection molding. The effect of fine (125–250 m) and coarse (500–1000 m) milling sizes of ASP was also evaluated. An improvement in elastic modulus was observed with the addition of< both fine and coarse ASP at 25 wt %. The addition of maleinized linseed oil and Joncryl ADR 4400 produced some compatibilizing effect at low filler contents while biocomposites with a higher amount of ASP still presented some gaps at the interface by field emission scanning electron microscopy. Some decrease in thermal stability was shown which was related to the relatively low thermal stability and disintegration of the lignocellulosic filler. The added modifiers provided some enhanced thermal resistance to the final biocomposites. Thermal analysis by differential scanning calorimetry and thermogravimetric analysis suggested the presence of two different polyesters in the polymer matrix, with one of them showing full disintegration after 28 and 90 days for biocomposites containing 25 and 10 wt %, respectively, under composting conditions. The developed biocomposites have been shown to be potential polyester-based matrices for use as compostable materials at high filler contents.

**Keywords:** almond shell waste; reinforcing; polyester-based biocomposites; physicochemical properties; disintegration

#### **1. Introduction**

Almond is characterized by its high nutritional value, although information reported so far mainly concerns its edible kernel or meat. Other parts also present in the almond fruit are the middle shell, outer green shell cover or almond hull and a thin leathery layer known as brown skin of meat or seed coat [1]. Almonds are used as a fruit in snack foods and as ingredients in a variety of processed foods, especially in bakery and confectionery products. However, almond production generates large amounts of almond by-products since the nutritional and commercial relevance of almonds is restricted to the kernel. In particular, almond shell is the name given to the ligneous material forming the thick endocarp or husk of the almond (*Prunus amygdalus L.*) tree fruit. It is principally composed of cellulose (ranging from 29.8 to 50.7 wt %), hemicellulose (from 19.3 to 29.0 wt %) and lignin (from 20.4 to 50.7 wt %) [1]. This by-product is normally incinerated or dumped without control, which results in the production of large amounts of waste and pollution [2]. Several researchers have focused on different alternatives for using almond shell wastes based on their potential uses as biomass to produce renewable energy [3]; as a source of organic biopesticides [4], heavy metal

adsorbents [5], dye adsorbents [6], growing media [7], the preparation of activated carbons [8] and xylo-oligosaccharides [9], antioxidants [10] or as additives in eco-friendly composites [11–13].

The development of eco-friendly composites arises from the need for reducing environmental problems generated by industrial processes. In this scenario, agricultural waste utilization has become a potential option for the development of eco-friendly composites. This powerful area of interest presents several benefits such as biodegradability in combination with bio-based or natural polymers, light weight, low cost and easy processing [14,15]. Among the wide variety of lignocellulosic wastes, almond shell powder has been already considered as filler for commodity plastics, such as polypropylene [16–20], polyethylene [21], poly(methyl methacrylate) [22] and toughened epoxies [23,24].

Looking at a more environmentally friendly use, the role of ASP has been recently studied in enhancing the mechanical performance of some melt compounded biopolymers [25–27]. Nonetheless, due to the lack of miscibility between hydrophobic polymer matrices and highly hydrophilic almond shell fillers, the obtained green composites usually presented poor ductility and low thermal stability. In order to increase the interaction between them, several solutions have been proposed, such as silanization, acetylation and maleic anhydride modification [28]. Plasticizers could also act as internal lubricants, thus allowing chain mobility, which enhances processability and improves thermal stability and ductility. Recently, vegetable oils have been proposed as environmentally friendly compatibilizers as an alternative to conventional petroleum-based ones [29]. Specifically, maleinized linseed oil (MLO) has been used as a compatibilizer in biopolymer/ASP composites [11,30,31]. In these works, authors discussed plasticization and compatibilization effects provided by MLO due to the interaction between succinic anhydride polar groups contained in MLO and hydroxyl groups in ASP (hydroxyl groups in cellulose). The compatibilizing effect was obtained by melt grafting for the formation of new carboxylic ester bonds through the reaction of maleic anhydride functionalities present in MLO with the hydroxyl groups of both the polyester terminal chains and cellulose on the ASP surface. On the other hand, the possibility of improving the stress transfer between the filler and the polymer can be realized by reactive processing with chain extenders [32].

The main aim of the present work is the development and characterization of new biocomposites prepared using a commercial INZEA® biopolymer (mainly composed of a polyester-based matrix) containing almond shell powder at 10 and 25 wt % contents. The effect of adding two different milling sizes (125–250 μm and 500–1000 μm) in the biocomposites preparation was also evaluated. In addition, the potential of maleinized linseed oil as a compatibilizer was studied. The effect of this vegetable-oil-derived compatibilizer was also compared with a conventional epoxy styrene-acrylic oligomer (Joncryl ADR 4400) in terms of mechanical properties, thermal stability and blend morphology. Biocomposites containing 10 and 25 wt % of ASP at two grinding levels were submitted to a disintegration test in order to verify the effectiveness of the developed polyester/ASP composites to be used as compostable materials.

#### **2. Materials and Methods**

#### *2.1. Materials*

INZEA® biopolyester commercial grade, with a density of 1.23 g cm−<sup>3</sup> measured at 23 ◦C, a moisture content <0.5% and a melt flow rate of 19 g/10 min (2.16 kg, 190 ◦C), was kindly supplied by Nurel (Zaragoza, Spain). Almond shell (AS) waste used as filler was supplied by Fecoam (Murcia, Spain) as an agricultural by-product and pulverized with a high-speed rotor mill (Ultra Centrifugal Mill ZM 200, RETSCH, Haan, Germany). The obtained particles were sieved by selecting the sizes of the ground shells in the ranges of 125–250 μm as fine grain (F) and 500–1000 μm as coarse grain (C) to evaluate the effect of particle size in the composites. Two different compatibilizers were selected to improve the compatibility between the polyester-based matrix and the natural filler: a synthetic polymer chain extender with recognized efficacy as compatibilizer supplied as Joncryl ADR 4400® (J44) (BASF S.A, Barcelona, Spain) and a biodegradable additive obtained by the maleinizing treatment

of linseed oil, supplied as Veomer Lin by Vandeputte (Mouscron, Belgium, viscosity of 10 dPa s at 20 ◦C and an acid value of 105–130 mg KOH g<sup>−</sup>1).

#### *2.2. Biocomposites Preparation*

Biocomposite materials were obtained using the melt blending method by mixing the biopolymer matrix with the almond particles, obtained by grinding and sieving the almond shells as previously described and the additives according to the proportions shown in Table 1. A co-rotating twin-screw extruder, Xplore 5 & 15 Micro Compounder by DSM, was used by mixing at a rotating speed of 90 rpm for 3 min and setting a temperature profile of 190–195–200 ◦C in the three heating zones from feeding section to die. A Micro Injection Molding Machine 10 cc by DSM, coupled to the extruder and equipped with adequate molds, was used to produce samples for flexural tests according to the standards. An appropriate pressure/time profile was used for the injection of each type of sample, while the temperatures of the injection barrel and the molds were set, respectively, at 210 and 30 ◦C.

**Table 1.** Formulations obtained in this work and their codification.


\* Fine grain (F): 125–250 μm; coarse grain (C): 500–1000 μm. J44: Joncryl ADR 4400; MLO: maleinized linseed oil.

#### *2.3. Almond and Biocomposites Characterization*

#### 2.3.1. Field Emission Scanning Electron Microscopy

Morphological characterization of ASP was carried out using a field emission scanning electron microscope (FESEM), Supra 25 by Zeiss (Oberkochen, Germany). The surfaces and the fractures of biocomposites were analyzed with a FESEM Merlin VP Compact by ZEISS. In both cases, micrographs were taken using an accelerating voltage of 5 kV at different magnifications. Samples were previously gold-sputtered with an Automatic Sputter Coater, B7341 by Agar Scientific (Stansted, Essex, UK), operating with a vacuum atmosphere (0.1–0.005 mbar) and low current (0–50 mA) to provide electric conductivity.

#### 2.3.2. Thermal Characterization

Thermogravimetric analysis of ASP was performed with thermogravimetric analysis (TGA; (Seiko Exstar 6300, Tokyo, Japan). Approximately 5 mg of samples were heated from 30 to 600 ◦C at 10 ◦C min−<sup>1</sup> under nitrogen atmosphere (flow rate 200 mL min<sup>−</sup>1).

Differential scanning calorimetry (DSC) tests were conducted for the determination of thermal events by using a DSC (Q1000, TA Instruments, New castle, DE, USA) under a nitrogen atmosphere (50 mL min<sup>−</sup>1). A 3 mg amount of samples were introduced in aluminum pans (40 μL) and they were submitted to the following thermal program: <sup>−</sup>30 ◦C to 250 ◦C at 10 ◦C min<sup>−</sup>1, with two heating and one cooling scans.

The thermal degradation behavior of biocomposites in composting conditions was evaluated by thermogravimetric analysis (TGA/SDTA851e/SF/1100, Mettler Toledo, (Schwarzenbach, Switzerland). Around 5 mg of samples were used to perform dynamic tests in a nitrogen atmosphere (200 mL min–1) from 30 ◦C to 700 ◦C at 10 ◦C min<sup>−</sup>1.

#### 2.3.3. Mechanical Properties

Flexural tests were carried out by using a universal test machine LR30K (Lloyd Instruments Ltd., Bognor Regis, UK) at room temperature. A minimum of five different samples was tested using a 0.5 kN load cell, setting the crosshead speed to 2 mm min−<sup>1</sup> for three points bending test, as suggested by ISO 178 Standard.

#### 2.3.4. Disintegrability in Composting Conditions

Disintegration tests in composting conditions were performed, in triplicate, by following the ISO 20200 Standard method using a commercial compost with a certain amount of sawdust, rabbit food, starch, oil and urea [33]. Tested samples were obtained from the previously prepared dog-bone-shaped bars, which were cut in pieces (5 <sup>×</sup> 10 <sup>×</sup> 2 mm3), buried at a 5 cm depth in perforated boxes and incubated at 58 ◦C. The aerobic conditions were guaranteed by mixing the compost softly and by the periodical addition of water according to the standard requirements.

Different disintegration times were selected to recover samples from burial and further tested: 0, 4, 7, 15, 21, 28, 40, 69 and 90 days. Samples were immediately washed with distilled water to remove traces of compost extracted from the container and further dried at 40 ◦C for 24 h before gravimetric analysis. The disintegrability value for each material at different times was obtained by normalizing the sample weight with the value obtained at the initial time.

The evolution of disintegration was monitored by taking photographs of recovered samples for visual evaluation of physical alterations with disintegration time. In addition, thermal (DSC, TGA) properties upon disintegrability tests were also studied.

#### *2.4. Statistical Analysis*

Statistical analysis of experimental data was performed by one-way analysis of variance (ANOVA) using SPSS 15.0 (IBM, Chicago, IL, USA) and expressed as means ± standard deviation. Differences between average values were assessed based on the Tukey test at a confidence level of 95% (*p* < 0.05).

#### **3. Results**

#### *3.1. Characterization of Almond Shell Powder*

#### 3.1.1. Morphological Analysis

Low-magnification FESEM micrographs of almond shell waste at the two different studied sizes (125–250 μm, fine, and 500–1000 μm, coarse), reported in Figure 1, showed a general view typical of fillers obtained after grinding and sieving processes. Almond shell was shown to have a sheet-like structure. In addition, a series of 1 μm pores were observed on the almond surface (see inserts) [34]. Most of the particles were characterized by a spherical shape, though some aggregates, as well as flat and long rod-like particles, were also observed. A detail of the particle surface can be seen in the high-magnification FESEM images. The micrographs revealed that the particles were irregular in shape and presented a rough surface, more likely resulting from the crushing process due to the high hardness of this type of filler. Some granular features can also be observed, which resemble the original grainy and wavy structure of almond shell [35].

**Figure 1.** FESEM images of fine (**a**) and coarse (**b**) almond shell powders.

#### 3.1.2. Thermal Properties

TG and derivative DTG profiles obtained for ASP (coarse size) under the nitrogen atmosphere at a heating rate of 10 ◦C min<sup>−</sup>1, reported in Figure 2, showed the typical thermal degradation profile for biomasses with three well-demarked steps for moisture release, devolatilization and char formation. Weight loss in the lower temperature region can be attributed to the loss of moisture, while major weight loss was observed at temperatures ranging from 225 to 365 ◦C, over which hemicellulose and cellulose decomposition occurs, leading to the formation of pyrolysis products (volatiles, gases and primary biochar) [36]. This phenomenon was followed by a slow weight loss until 600 ◦C, which was attributable to the continuous devolatilization of biochar caused by a further breakdown of C–C and C–H bonds.

**Figure 2.** TGA analysis of almond coarse shell powder.

#### *3.2. Characterization of ASP Biocomposites*

#### 3.2.1. Flexural Tests

Flexural tests provided information on the effect of the amount and size of ASP incorporated in the composite, as well as on the effect of adding the studied compatibilizing additives. In general, results show that the addition of the filler did not improve the maximum strength or strain at break with respect to the reference polymer matrix (Table 2). The elastic modulus of the biocomposites was improved only in formulations containing 25 wt % of both fine and coarse filler. The comparison of flexural tests for formulations with 10 wt % of ASP (Table 2 and Figure 3) showed that the size of the coarse grain gives greater rigidity than the fine grain by increasing strength and modulus but slightly reducing elongation. In fact, the presence of 10 wt % of fine filler in the INZEA\_10ASF

biocomposite produced a maximum strength value of 44 MPa with a flexural modulus of 1473 MPa, lower than the biocomposite with the coarse filler INZEA\_10ASC, which showed σmax = 47 MPa and E = 1699 MPa values. Nabinejad et al. [37] reported that the surface roughness and high surface area of coarse powder particles could have a positive effect on the mechanical performance of composites, being able to restrict the polymer chain mobility. Porosity and roughness of the hydrophilic surface for coarse almond shell powder would be expected to increase its wettability by the polymer matrix. So, as a result, the INZEA\_10ASC composite showed high stiffness values compared to composites containing fine filler (INZEA\_10ASF) with low surface roughness and porosity. Additionally, results from Zaini et al [38] confirmed that composites filled with a larger-sized filler showed higher modulus, tensile and impact strengths, particularly at high filler loadings.


**Table 2.** Flexural parameters for almond shell powder (ASP)-based biocomposites (mean ± SD, *n* = 5).

σmax: flexural strength; ε at σmax: strain at maximum stress; E: Young's Modulus.

**Figure 3.** Stress–strain curves of polyester-based biocomposites containing 10 and 25 wt % of ASP at two grinding levels (F, C), with or without compatibilizers.

The addition of 5 wt % of MLO in the biocomposite formulation produced a compatibilizing effect lower than expected. In fact, the improvement in deformability moved from 4.4% to 4.9% of INZEA\_10ASF\_5MLO and from 4.0% to 5.4% of INZEA\_10ASC\_5MLO in the formulations with MLO and 10 wt % ASP fine and coarse, respectively, with a consistent reduction in both flexural strength and flexural modulus. If the modest plasticizing effect is excluded, the MLO did not produce improvements of the interface strength between the natural filler and the biopolymer matrix. Effectiveness of maleinized linseed oil has been demonstrated in composites based on lignin fillers and, specifically, from ground almond shells and some biodegradable polymers, such as poly(lactic acid) (PLA) and lignin. In this case, the poor result obtained with the INZEA matrix should be attributed to the

particular composition of the commercial biopolymer used (a bio-based blend mainly composed of a polyester matrix) [30,39,40]. Joncryl ADR 4400 added to the formulation with fine particles INZEA\_10ASF\_1J produced an increase in flexural strength from 44 to 48 MPa. This improvement is due to the compatibilizing effect of the polymer chain extender, which improves the adhesion between the matrix and the particles allowing a greater deflection. This effect is highlighted by the perfect overlap between the σ−ε curve of the unmodified INZEA\_10ASF and the INZEA\_10ASF\_1J curve which, thanks to the J44, extends up to 5.5% increasing the flexural strength [41]. The effect of Joncryl was negligible on biocomposites with coarse grain, since the compatibilization effect at the interface between the matrix and the filler was much lower (about 6%) than biocomposites with fine particles. In fact, when simplifying the particles as spherical and calculating the ratios between coarse and fine surfaces, with the same wt % content, an area ratio of 0.0625 was obtained.

When increasing the quantity of filler to 25 wt %, an increase in the rigidity of the biocomposites was obtained. The result is a general increase in flexural strength and moduli, which corresponds to a reduction in deflection. Even in formulations with a higher content of the natural filler, MLO did not produce any other effects than those already shown in the set with 10 wt % of ASP. The flexural strength of the biocomposite containing 25 wt % of fine particle (INZEA\_25ASF) rises to 50 MPa (44 MPa for INZEA\_10ASF), while the INZEA\_25ASC composite maintains the same value of 47 MPa as 10 wt % of filler, highlighting the achievement of the plateau of the coarse grain reinforcement. In the case of the higher amount of ASP, the compatibilizing effect of Joncryl appears even more evident, since a better interface bonding between the matrix and the filler was achieved. INZEA\_25ASF\_1J showed a further increase in flexural strength reaching 56 MPa, with an improvement in modulus to 2555 MPa and without excessively reducing the flexural deflection. In addition, INZEA\_25ASC\_1J with Joncryl showed an improvement in strength going up from 47 to 53 MPa, but in this case, the compatibilizing effect of J44 was less evident, because it occurred on a smaller interface surface, compared to fine-grain-sized biocomposites, due to the larger particle size [13].

#### 3.2.2. Morphological and Thermal Analysis

In Figure 4, FESEM images of fractured surfaces for INZEA/ASP composites (uncompatibilized and compatibilized biocomposites) are reported. As it can be observed, in the case of the unmodified matrix, the polymer-particle adhesion was very poor, both at low and high ASP contents, so important gaps can be found between the particles and the surrounding polyester matrix [11]. This morphological observation correlates with the above-described mechanical performance of the unmodified INZEA composites, in which the presence of ASP did not contribute to an improvement in mechanical performance. The addition of 5 wt % MLO provides some interaction as the gap seems to be reduced, indicating a limited but good compatibilizing effect of MLO modifier, but at the same time, the presence of microsized voids in the matrix compatibilized with 5 wt % of MLO was visible [30]. In the presence of a styrene-acrylic-based compatibilizer and oligomeric agent, an improvement in the interface with ASF and ASC was noted at higher contents (25 wt %), even if biocomposites containing the higher amount presented some gaps at the interface. So, it can be concluded the chain extender was effective at both filler contents [42].

Figure 5a,b shows the TG/DTG thermograms of the INZEA matrix with the addition of 25 wt % ASP at the two different grinding sizes. Only the trends in thermal behavior for the biocomposites containing the higher ASP content have been reported, being the ones at 10 wt % essentially inline (data not shown). The presence of a double degradation peak for the INZEA neat matrix gives us an indication of a material that degrades in two steps around 350 and 400 ◦C, that could match with the possible degradation temperatures of PLA and poly(butylene succinate) (PBS) polyesters. It is also important to note that at 900 ◦C, even the unmodified INZEA matrix maintains a residual mass of ca. 5 wt %, which is increased by the presence of the fillers. This is in accordance with the possible presence of an inorganic filler in the formulation of the commercial product.

**Figure 4.** FESEM images of INZEA-based biocomposites with 10 and 25 wt % of ASP at the two grinding levels (fine, coarse), with or without compatibilizers.

**Figure 5.** TG/DTG curves of INZEA-based biocomposites with 25 wt % of ASP at the two grinding levels (**a**,**b**) and INZEA biocomposites with 25 wt % of ASP in the presence of MLO or Joncryl compatibilizers (**c**,**d**).

The TG curves corresponding to INZEA/ASP composites showed an evident decrease at the onset degradation temperature from 324 ◦C for neat INZEA to 275 ◦C and 280 ◦C, respectively, for INZEA\_25ASC and INZEA\_25ASF (Table 3). This behavior is essentially due to the relatively low thermal stability of the lignocellulosic filler, which initiated its degradation at 198 ◦C (Figure 2), and negatively contributed to the reduction of the global thermal stability of the INZEA-based biocomposites, in agreement with previous studies on the same bio-based reinforcement [17,20]. The introduction of the ASP filler mainly affected the thermal stability of the polyester component with the lower Tpeak temperature (maximum mass loss rate): in detail, Tpeak1 moved from 351 ◦C to 311 ◦C and 334 ◦C for INZEA\_25ASF and INZEA\_25ASC, respectively, while the temperature for the second peak (Tpeak2) remained practically unchanged for all the different composites (Table 3). This behavior was related to the higher amount of ASP incorporated and its degradation over this temperature range. According to other authors, cellulose, hemicelluloses and lignin show a broad temperature range starting at about 250 ◦C and ending at 450 ◦C in a progressive weight loss process, and this degradation can reduce the thermal stability of biopolyesters [43–45]. Similar findings were reported by Liminana et al. [31] who observed a decrease of 11.2 ◦C in thermal stability of PBS with the addition of 30 wt % of almond shells.

**Table 3.** Tonset and Tpeak values of INZEA-based biocomposites with 25 wt % of ASP at the two grinding levels with and without MLO or Joncryl compatibilizers (mean ± SD, *n* = 3).


Tonset: initial degradation temperature; Tpeak1 and Tpeak2: first and second maximum degradation temperatures, respectively.

The addition of MLO (Figure 5c,d) exerted a limited positive effect on the overall thermal stability of the biocomposites (Table 3). Specifically, the onset degradation temperature increased to 285 ◦C for INZEA\_25ASF\_5MLO and substantially remained unchanged for INZEA\_25ASC\_5MLO (if compared to uncompatibilized matrices). On the other hand, the temperature of Tpeak1 was significantly improved, in comparison with the unmodified ASP biocomposites, at the two grinding sizes. In particular, Tpeak1 was delayed up to 328 ◦C for the INZEA\_25ASF\_5MLO material and 341 ◦C for the INZEA\_25ASC\_5MLO material. This increase in thermal stability could be directly related to the chemical interaction achieved by MLO, due to the establishment of covalent bonds between the lignocellulosic fillers and the polyester matrix. In addition, MLO could also provide a physical barrier that obstructs the removal of volatile products produced during decomposition. A similar effect on thermal stability was recently reported for PLA and epoxidized palm oil blends [46].

It has been also reported that the reactive extrusion of aliphatic polyesters, such as PLA, with styrene-epoxy acrylic oligomers can provide an increase in thermal stability due to the branching effect obtained during the extrusion process. As it has been found in our case (Figure 5c,d), the presence of Joncryl provided enhanced thermal resistance for the low Tpeak polyester phase, moving this value from 311 ◦C and 334 ◦C (INZEA\_25ASF and INZEA\_25ASC), respectively, to 329 ◦C and 339 ◦C for INZEA\_25ASF\_1J and INZEA\_25ASC\_1J (Table 3). This effect was already observed by Lascano et al in poly(lactic acid)/poly(butylene succinate-co-adipate) blends containing Joncryl [47].

#### *3.3. Disintegration Tests*

According to the characterization results previously obtained, the disintegrability of polyesterbased biocomposites with 10 and 25 wt % of ASP at two grinding levels (F, C) was studied to evaluate their degradation in natural environments. Formulations including the two studied compatibilizing agents were not included in this study, since the main goal of this characterization was the analysis of ASP size and content on disintegration behavior of the reference matrix.

The visual evaluation of all samples at different degradation times was carried out and results are shown in Figure 6. Some changes in sample surfaces submitted to composting conditions were clearly appreciable, showing all samples considerable modifications in color and morphology after 15 treatment days. These modifications can be associated to the beginning of the polymer matrix degradation, which can be related to the direct contact of biocomposites with the compost, by a gradually microorganism erosion from the surface to the bulk and to the moisture absorption, as it has been reported in the literature [45]. After 69 days of study, samples with 25 wt % of ASP have completely lost their morphology, and small fragments can be observed (in particular for INZEA\_25ASC). In addition, differences in color could be observed between samples with different amounts of ASP. Formulations with 10 wt % of ASP showed disintegration behavior similar to that of INZEA neat during all the study.

**Figure 6.** Visual appearance of INZEA-based biocomposites with 10 and 25 wt % of ASP at two grinding levels (F, C) at different testing days at 58 ◦C.

Figure 7 shows the evolution of disintegrability values (%) as a function of testing time for all biomaterials. According to de Olivera et al. [48], the first stage of the biodegradation mechanism is the release of enzymes that can cause the hydrolysis of the polymer matrix and break of polymer chains, creating functional groups capable of improving hydrophilicity and the microorganism's adhesion on the surface of the polymer matrix. The results obtained at longer times suggested that physical degradation progressed slowly with burial time, indicating that microorganisms required more time to produce suitable enzymes capable to break down polymer chains, resulting in an incomplete loss of the initial morphology and general rupture after 90 days for formulations with 10 wt % of ASP and the INZEA control. Formulations with 25 wt % of ASP significantly increased their disintegrability ratio compared to formulations with 10 wt % of ASP and INZEA control after 28 days of study, probably due to the higher amount of ASP, which enhances the high biodegradability of lignocellulosic residues, and to the poor fiber/matrix adhesion allowing and facilitating microorganisms attack and biodegradation rate by promoting biofouling and the adhesion of microorganisms to the surface [49]. Moreover, this increase in the disintegrability rate of the polymer matrix could be due to the presence of hydroxyl groups in ASP [49], which could play a catalytic role on the hydrolysis of the polymer, inducing an acceleration of polymer weight loss due to the higher filler addition [50]. Similar results were found by Wu [51] and de Oliveira et al. [48] when studying the biodegradation of composites obtained with poly(butylene adipate-co-terephthalate) and different natural fillers. The authors observed that the biodegradation rate of the composites increased with filler content. In this sense, the presence of high amounts of ASP can be related to a greater discontinuity in the polymer matrix which could facilitate

water penetration into the biocomposites producing a huge modification of the surface and generating a natural environment conducive to the growth of microorganisms [48]. After 90 days of study, almost 50% of the materials were disintegrated under composting conditions. However, lower weight loss ratios were obtained with lower amounts of ASP. This behavior suggests that the disintegration rate is more influenced and dependent on the polymer matrix being slower at lower ASP contents.

**Figure 7.** Disintegrability (%) of INZEA-based biocomposites with 10 and 25 wt % of ASP at two grinding levels (F, C) as a function of degradation time under composting conditions at 58 ◦C (mean ± SD, *n* = 3).

Figure 8 shows the DSC thermograms obtained during the second heating scan for all formulations as a function of composting time (0, 28 and 90 days). Two different peaks were observed at day 0, around 110 and 170 ◦C, indicating the presence of two main polyesters in the polymer matrix, in agreement with the behavior previously observed by TGA (Section 3.2.2). A similar DSC profile for the first and second peaks was described by Liminana et al. and Quiles-Carrillo et al. for the characterization of PBS-based and PLA-based composites reinforced with similar amounts of almond shells, respectively [12,52]. Both melting peak temperatures remained practically invariable after the addition of ASP at the two studied contents, 10 and 25 wt % (Table 4), showing a slight modification which could be related to the formation of more perfect crystals into the polymer matrix. Similar behavior was observed by Quiles-Carrillo et al. for PLA-based composites with 25 wt % of almond shell [12]. A slight decrease in the melting enthalpy of the first peak was observed with the addition of ASP which was more pronounced at 25 wt % This effect could be related to the nucleating effect exerted by the lignocellulosic filler on semicrystalline polymers acting the cellulose crystals of the almond shells as nucleating points [11].

Regarding the disintegration study, the DSC melting temperature of the second endothermic peak moved from 168.4 ± 3.7 ◦C for INZEA at day 0 to 142.4 ± 0.8 ◦C at day 90 (Table 4). Formulations with ASP did not show significant differences at day 0 compared to the values obtained for INZEA control. This decrease in around 26 ◦C was related to a rapid molecular mass reduction, implying that

small and imperfect crystals disappeared with degradation time [53,54]. INZEA and formulations with 10 wt % of ASP did not show significant differences in DSC values at the same disintegration time, maintaining similar values throughout the whole study. In this sense, the addition of 10 wt % of ASP could not be enough to achieve an acceptable weight loss ratio into the disintegration process under composting conditions.

**Figure 8.** DSC thermograms of INZEA-based biocomposites with 10 and 25 wt % of ASP at two grinding levels (F, C) after different degradation times ((**a**): 0 days, (**b**): 28 days, (**c**): 90 days) at 58 ◦C during the first heating scan (10 ◦C min<sup>−</sup>1).

Under composting conditions, a different behavior was observed for formulations with 25 wt % of ASP versus time, as the second DSC peak initially appearing around 170 ◦C started to disappear after 28 days of study (Figure 8). After 90 days, the appearance of the thermogram suggested that this polyester-based polymer was totally disintegrated by disappearing the corresponding glass transition temperature (around 50 ◦C) and melting peak around 170 ◦C (Table 4). This result was also related to the final appearance and percentage of disintegrability achieved in samples after 90 days. On the other hand, the observed peak around 110 ◦C remained unchanged after 90 days of study, with slight modifications on its profile, indicating that this polymer was not degraded yet.

This behavior was also confirmed by TGA analysis (Table 4). The obtained results demonstrated that after 28 days of study the maximum degradation temperature of the first polymer peak, Tpeak1, decreased around 70 ◦C respect to day 0, disappearing this degradation peak from the TGA curve after 40 days under composting conditions. These results confirm those obtained by DSC where some modification of the thermal profile was observed due to the disintegration of one main polyester component of the polymer matrix.



*Polymers* **2020** , *12*, 835

◦C

respectively; ΔHm1 and ΔHm2: First and second enthalpies of fusion, respectively; Tg: glass transition temperature.

#### **4. Conclusions**

In this work, biocomposite materials were obtained based on a polyester matrix (INZEAF2) and almond shell as a reinforcing agent at 10 and 25 wt % and two different milling sizes (125–250 μm and 500–1000 μm). MLO and Joncryl ADR 4400 were studied as compatibilizers. The reinforced effect of the addition of ASP was shown by increasing the elastic modulus of the biocomposites with both fine and coarse ASP at 25 wt % The addition of MLO and Joncryl produced some compatibilizing effect at 10 wt % while some gaps at the interface, visible by FESEM at 25 wt %, did not substantially affect the overall flexural properties. A double degradation pattern was obtained by TGA for the INZEA neat matrix indicating the degradation of the material in two steps and the presence of two different polyesters in the matrix. Some decrease in thermal stability of biocomposites was shown which was more pronounced at high ASP contents and it was related to the relatively low thermal stability and disintegration of the lignocellulosic filler over the studied temperature range. Some limited positive enhancement in thermal stability was obtained by adding the studied modifiers. The addition of 10 wt % of ASP was not enough to achieve an acceptable weight loss ratio of disintegration under composting conditions, whereas around 50% of disintegration (nearly double of the polymer control) was obtained by adding 25 wt % of ASP after 90 days. At these conditions, one of the polyester-based polymers was totally disintegrated by disappearing the corresponding glass transition temperature (around 50 ◦C), melting peak (around 170 ◦C) and degradation temperature (around 350 ◦C). On the other hand, the second polymer with Tm around 110 ◦C was not degraded yet.

The developed composites represent an interesting approach to reduce the overall cost of bio-based polyesters increasing also the added-value potential of almond agricultural wastes, obtaining environmentally friendly materials with specific reinforced and aesthetic functionalities and contributing to the circular economy approach. Further work will be needed in order to evaluate a possible improvement in filler–matrix adhesion and disintegration rate of the biocomposites by using different compatibilizers and adding higher amounts of lignocellulosic filler without compromising the final mechanical and thermal properties.

**Author Contributions:** Conceptualization, A.J., M.C.G., L.T. and D.P.; methodology M.R., F.D., and F.L.; validation, A.J., M.C.G., L.T. and D.P.; formal analysis, M.R., F.D., and F.L.; investigation, M.R., F.D., and F.L.; resources, A.J., M.C.G., L.T. and D.P.; data curation, M.R., F.D., F.L., A.J., M.C.G. and D.P.; writing—original draft preparation, M.R., F.D., and F.L.; writing—review and editing, A.J., M.C.G. and D.P.; supervision, A.J., M.C.G., L.T. and D.P.; funding acquisition, A.J., M.C.G., L.T. and D.P. All authors have read and agreed to the published version of the manuscript.

**Funding:** The authors express their gratitude to the Bio Based Industries Consortium and European Commission for the financial support to the project BARBARA: Biopolymers with advanced functionalities for building and automotive parts processed through additive manufacturing. This project has received funding from the Bio Based Industries Joint Undertaking under the European Union's Horizon 2020 research and innovation programme under grant agreement No 745578.

**Conflicts of Interest:** The authors declare no conflicts of interest.

#### **References**


© 2020 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### *Article* **Mechanical Recycling of Partially Bio-Based and Recycled Polyethylene Terephthalate Blends by Reactive Extrusion with Poly(styrene-***co***-glycidyl methacrylate)**

**Sergi Montava-Jorda 1, Diego Lascano 2,3, Luis Quiles-Carrillo 2, Nestor Montanes 2, Teodomiro Boronat 2, Antonio Vicente Martinez-Sanz 4, Santiago Ferrandiz-Bou <sup>2</sup> and Sergio Torres-Giner 5,\***


Received: 5 December 2019; Accepted: 3 January 2020; Published: 9 January 2020

**Abstract:** In the present study, partially bio-based polyethylene terephthalate (bio-PET) was melt-mixed at 15–45 wt% with recycled polyethylene terephthalate (r-PET) obtained from remnants of the injection blowing process of contaminant-free food-use bottles. The resultant compounded materials were thereafter shaped into pieces by injection molding for characterization. Poly(styrene-*co*-glycidyl methacrylate) (PS-*co*-GMA) was added at 1–5 parts per hundred resin (phr) of polyester blend during the extrusion process to counteract the ductility and toughness reduction that occurred in the bio-PET pieces after the incorporation of r-PET. This random copolymer effectively acted as a chain extender in the polyester blend, resulting in injection-molded pieces with slightly higher mechanical resistance properties and nearly the same ductility and toughness than those of neat bio-PET. In particular, for the polyester blend containing 45 wt% of r-PET, elongation at break (εb) increased from 10.8% to 378.8% after the addition of 5 phr of PS-*co*-GMA, while impact strength also improved from 1.84 kJ·m−<sup>2</sup> to 2.52 kJ·m−2. The mechanical enhancement attained was related to the formation of branched and larger macromolecules by a mechanism of chain extension based on the reaction of the multiple glycidyl methacrylate (GMA) groups present in PS-*co*-GMA with the hydroxyl (–OH) and carboxyl (–COOH) terminal groups of both bio-PET and r-PET. Furthermore, all the polyester blend pieces showed thermal and dimensional stabilities similar to those of neat bio-PET, remaining stable up to more than 400 ◦C. Therefore, the use low contents of the tested multi-functional copolymer can successfully restore the properties of bio-based but non-biodegradable polyesters during melt reprocessing with their recycled petrochemical counterparts and an effective mechanical recycling is achieved.

**Keywords:** bio-PET; r-PET; chain extenders; reactive extrusion; secondary recycling; food packaging

#### **1. Introduction**

The transition of the plastic industry from its traditional Linear Economy to a Circular Economy, a more valuable and sustainable model, is being spearheaded by the European Union (EU), where legislative measures are being introduced to eliminate excessive waste [1]. To this end, it is first necessary to promote sustainable polymer technologies that decouple plastics from fossil feedstocks [2]. Furthermore, it is also important to increase the quality and uptake of plastic recycling [3]. In this context, the food packaging sector it is among the most heavily scrutinized, given the large production and the short life cycle of their products [4]. Polyethylene terephthalate (PET) is a thermoplastic polyester that is widely used in the manufacture of bottles for water or beverages and also food trays due to its good mechanical properties, chemical resistance, clarity, thermal stability, barrier properties, and low production cost [5]. PET is fully recyclable and it is, indeed, one of the most recycled plastics in the world. Since 2012, PET monomaterial packaging has showed recycling rates of approximately 52% in the EU and 31% in the United States (US) [6]. Moreover, the recycling of PET articles is expected to increase, as exemplified by companies like Coca-Cola, which have already announced a full switch to recycled plastics for their beverage bottles in some EU countries, starting with at least 50% rates by the end of 2023 [7].

In recent years, the replacement of conventional or petrochemical polymer materials with those obtained from natural and renewable sources is of great interest for research due to the depletion of fossil resources and the cost of extracting them. The main interest and advantage of using bio-based polymers reflects the concept of the so-called "biorefinery system design" due to its ability to improve the environmental impact of a product by reducing greenhouse gas emissions, economizing fossil resources, exploring the possibility of using a local resource, and the use of by-products or even wastes [8]. Bio-based polymers, which can be either biodegradable or non-biodegradable, are certified according to international standards such as EN 16640:2015, ISO 16620-4:2016, ASTM 6866-18, and EN 16785-1:2015. In particular, EN 16640:2015, ISO 16620-4:2016, and ASTM 6866-18 measure the bio-based carbon content in a material through 14C measurements, while EN 16785-1:2015 measures the bio-based content of a material using radiocarbon and elemental analyses [9]. Indeed, the pattern of production of biopolymers is shifting from biodegradable to bio-based. Bio-based but non-biodegradable polymers represented 57.1% of the total biopolymer production in 2018, whereas biodegradable polymers accounted for 42.9% [10]. This is in contradiction to the public perception that most biopolymers are biodegradable. Among them, partially bio-based polyethylene terephthalate (bio-PET) is currently the most produced bioplastic, reaching ~27% of the total production in 2018, that is, 0.54 million tons per year. [11]. This is based on the fact that this "green polyester" offers an almost identical chemical structure and properties to its petrochemical counterpart, that is, PET.

Several studies on PET recycling methods have indicated that mechanical recycling appears to be the most desirable method for the management of PET waste, as compared to chemical recycling and incineration [12–14]. The chemical or tertiary recycling method involves depolymerization of the PET polymer by chemical agents (chemolysis) or temperature (pyrolysis) to obtain its constituent monomers, that is, monoethylene glycol (MEG) and terephthalic acid (TA), or their derivatives, which can be used for new polymerization processes [15,16]. In addition, the char from pyrolysis of washed PET wastes can successfully replace up to 50 wt% the epoxy resin used in the production of thermosets [17]. In mechanical or secondary recycling, PET waste is subjected to mechanical processes including shredding, grinding, melting, and, when necessary, as in the case of contaminated articles, is combined with washing and/or drying [18]. In this process, the polymers stay intact, which permits multiple reuses in the same or similar products. However, mechanical recycling is currently the preferred option for monomaterial packaging since it has the advantages of simplicity and low cost, requires little investment, uses established equipment, and has little adverse environmental impact. Despite these advantages, mechanical recycling of PET is difficult due to complexity and waste contamination [12,19,20]. Another issue observed during the mechanical recycling of PET by "fusion reprocessing" is that the polymer is habitually subjected to chemical, mechanical, thermal, and oxidative degradation, which decreases the molar mass and, finally, causes the deterioration of the performance and transparency of PET articles [21,22].

Different studies have conventionally supported the notion that the use of recycled polymer streams in a mixture with virgin polymer of the same nature is a good solution for improving the properties of recycled polymer materials and achieving mechanical recycling. Elamri et al. [23] investigated fibers from blends of recycled polyethylene terephthalate (r-PET) and virgin PET. An improvement in the melt processing of r-PET and its fibers, with similar mechanical characteristics to those obtained from virgin PET, was reported. In another study, Scarfato and La Mantia [24] studied recycled and virgin mixtures of polyamide 6 (PA6), showing mechanical and rheological properties similar to those of the virgin polyamide. Moreover, Elamri et al. [25] also investigated blends of recycled and virgin high-density polyethylene (HDPE), reporting a predictable linear behavior in the mechanical properties as a function of the recycled content. A novel, plausible solution to increase the properties of r-PET during melt reprocessing is the use of additives such as chain extenders [26–29]. This option represents a more economical and attractive strategy than chemical recycling processes for monomaterial packaging since these additives can be used during reprocessing cycles by extrusion or injection molding [26,27]. Chain extenders are additives containing at least two functional groups that are capable of reacting with the end groups of different macromolecular fragments, thereby creating new covalent bonds. During melt reprocessing, these additives are able to "reconnect" the previously broken polymer chains, leading to a polymer with higher molecular weight (MW) and restored properties [27,28].

The use of chain extenders during PET reprocessing has been widely investigated. Some studies have reported the use of pyromellitic dianhydride (PMDA) [30,31], organic phosphites [32,33], bis-oxazolines [34–36], bis-anhydrides [30,37], diisocyanates [29,34,38], diepoxides [34], epoxy-based oligomers [34,39], or oligomeric polyisocyanates [40]. However, new commercial chain extenders based on random copolymers of poly(styrene-*co*-glycidyl methacrylate) (PS-*co*-GMA) have recently arisen. In particular, the glycidyl methacrylate (GMA) multi-functionality has excellent affinity with condensation polymers, not only PET, but also polybutylene terephthalate (PBT) or polylactide (PLA). It performs by connecting the polymer chains through a chemical reaction of the GMA groups with the hydroxyl (–OH), carboxyl (–COOH), and amine (–NH2) terminal groups of the polycondensation polymers. This process can result in a branched structure that also offers higher melt strength for optimal processing conditions. Although some other block copolymers based on GMA groups have been previously used as a chain extenders in other types of polymer blends, such as poly(trimethylene terephthalate) (PTT)/polystyrene (PS) [41] or PET/polypropylene (PP) [42], their use in PET systems remains almost unexplored. Only Benvenuta et al. [43] recently synthesized reactive tri-block copolymers of styrene glycidyl methacrylate (SGMA) and butyl acrylate (BA) (SGMA-*co*-BA-*co*-SGMA) by reversible addition-fragmentation chain transfer (RAFT) and used them as chain extenders of r-PET. These additives turned out to be very effective at increasing the molar mass and intrinsic viscosity, diminishing the melt flow, and improving the melt elasticity and processability of r-PET.

Despite the large amount of work that has been undertaken on the recycling of PET and r-PET blends, those based on bio-PET are barely beginning to be studied. Furthermore, while the price of virgin PET remains stable [44], the use of r-PET is significantly increasing, and thus, novel and more sustainable technologies for PET recycling could encourage more competitive prices in the industry. In the present work, the properties of bio-PET and pre-consumer (uncontaminated) r-PET blends were analyzed to ascertain the potential for mechanical recycling of the next generation of PET materials. To this end, different amounts of r-PET were melt-mixed with bio-PET in a twin-screw industrial extruder and then shaped into pieces by injection molding. During melt compounding, PS-*co*-GMA was added and the effect of the random copolymer on bio-PET/r-PET blends was assessed through a detailed characterization, including measurements of the mechanical, thermal, and thermomechanical properties as well as the rheological behavior.

#### **2. Materials and Methods**

#### *2.1. Materials*

Bio-PET, commercial grade BioPET 001, was obtained from NaturePlast (Ifs, France). According to the manufacturer, this grade is produced from up to 30 wt% renewable materials. It has a melting temperature (Tm) between 240–260 ◦C, a true density of 1.3–1.4 g·cm−3, and an intrinsic viscosity between 75–79 mL·g−<sup>1</sup> while it is supplied with a water content of less than 0.4 wt%. Further details can be found elsewhere [45]. r-PET was obtained from remnants of the injection blowing process of contaminant-free food-use bottles supplied by the local company Plásticos Guadalaviar S.A. (Beniparell, Spain). The pre-consumer PET waste was shredded and supplied in the form of flakes. PS-*co*-GMA was provided by Polyscope Polymers B.V. (Geleen, The Netherlands) as XibondTM 920 in the form of pellets. It has a mass average MW of 50,000 g·mol−<sup>1</sup> and a glass transition temperature (Tg) of 95 ◦C. The manufacturer recommends a dosage level of 0.1–5 wt% while not exceeding 330 ◦C to avoid thermal degradation.

#### *2.2. Preparation of the Bio-PET*/*r-PET Blends*

Both bio-PET and r-PET were dried at 60 ◦C for 72 h, while PS-*co*-GMA was dried at 90 ◦C for 3 h. All materials were dried separately to avoid premature cross-linking and then premixed in a zipper bag. Table 1 summarizes the coding and compositions of the prepared samples.

**Table 1.** Code and composition of each sample according to the weight content of partially bio-based (bio-PET) and recycled polyethylene terephthalate (r-PET) in which poly(styrene-*co*-glycidyl methacrylate) (PS-*co*-GMA) was added as parts per hundred resin (phr) of blend.


The compounding process was carried out in a twin-screw extruder from Construcciones Mecánicas Dupra, S.L. (Alicante, Spain) equipped with a screw diameter of 25 mm and a length-to-diameter (L/D) ratio of 24. The rotating speed was set to 25 rpm and the temperature profile from the feeding to the die was: 240 ◦C (hopper)–245 ◦C–250 ◦C–255 ◦C (nozzle). The strands were cooled in air and granulated in pellets in an air-knife unit. The resultant pellets were dried at 60 ◦C for 72 h to remove moisture since PET has a high sensitivity to hydrolysis.

Standard pieces for characterization were obtained from the compounded pellets by injection molding in a Meteor 270/75 from Mateu&Solé (Barcelona, Spain). The temperature profile in the injection machine was programmed to 250 ◦C (hopper)–255 ◦C–255 ◦C–260 ◦C (injection nozzle). The resultant pieces were finally annealed at 60 ◦C for 72 h to further develop crystallinity, improve their dimensional stability, and remove any residual moisture.

#### *2.3. Mechanical Characterization*

Tensile properties were obtained in a universal test machine ELIB-50 from S.A.E. Ibertest (Madrid, Spain) following ISO 527-1:2012. The selected cross-head speed was 5 mm·min−<sup>1</sup> and the load cell was 5 kN. As recommended by the standard, the tensile modulus (Et), maximum tensile strength at break (σmax), and elongation at break (εb) values were determined. Shore D hardness was obtained in a 676-D durometer from Instruments J. Bot S.A. (Barcelona, Spain) as indicated in ISO 868:2003. The impact strength was obtained in a Charpy's pendulum of 1 J from Metrotec (San Sebastián, Spain) on V-notched samples with a radius notch of 0.25 mm according to ISO 179-1:2010. All the mechanical tests were performed under controlled conditions of 25 ◦C and 40% relative humidity (RH). At least 6 samples of each material were evaluated.

#### *2.4. Color Measurements*

Changes in color were measured in a colorimetric spectrophotometer ColorFlex from Hunterlab (Reston, VA, USA). The selected color space was the chromatic model L\* a\* b\* or CIELab (spherical color space). L\* stands for the luminance, where L\* = 0 represents dark and L\* = 100 indicates clarity or lightness. The a\*b\* pair represents the chromaticity coordinate, where a\* > 0 is red, a\* < 0 is green, b\* > 0 is yellow, and b\* < 0 is blue. The L\*a\*b\* coordinate values were obtained on five different samples and the color difference (ΔE\*) was calculated using the following Equation (1):

$$
\Delta \mathbf{E}^\* = \left[ \left( \Delta \mathbf{L}^\* \right)^2 + \left( \Delta \mathbf{a}^\* \right)^2 + \left( \Delta \mathbf{b}^\* \right)^2 \right]^{0.5} \tag{1}
$$

where ΔL\*, Δa\*, and Δb\* correspond to the differences between the color parameters of the tested pieces and the values of the neat bio-PET (B100) piece. Color change was evaluated as follows: Unnoticeable (ΔE\* < 1), only an experienced observer can notice the difference (ΔE\* ≥ 1 and < 2), an unexperienced observer notices the difference (ΔE\* ≥ 2 and < 3.5), clear noticeable difference (ΔE\* ≥ 3.5 and < 5), and the observer notices different colors (ΔE\* ≥ 5) [46].

#### *2.5. Microscopy*

The fracture surfaces of the injection-molded samples after the impact test were covered with a thin metal layer to provide the conducting properties for analysis by field emission scanning electron microscopy (FESEM). The sputtering process was carried out in a cathodic sputter-coater Emitech SC7620 from Quorum Technologies LTD (East Sussex, UK). Subsequently, the morphologies were observed in a CARL ZEISS Ultra-55 FESEM microscope from Oxford Instruments (Abingdon, UK). The acceleration voltage was set to 2.0 kV.

#### *2.6. Thermal Characterization*

Thermal characterization was carried out on a differential scanning calorimeter DSC-821 from Mettler-Toledo Inc (Schwarzenbach, Switzerland). Samples with a mean weight of 6.1 ± 1.2 mg were placed in aluminum pans and subjected to a dynamic thermal program in three stages: initial heating from 30 ◦C to 280 ◦C, cooling to 0 ◦C, and a second heating to 350 ◦C. The heating/cooling rates were set at 10 ◦C ·min−<sup>1</sup> for all three stages, and the atmosphere was air at a flow-rate of 66 mL·min−1. The DSC runs were obtained in triplicate to obtain reliable results. The values of Tg, cold crystallization temperature (Tcc), Tm, as well as the enthalpies corresponding to the melting process (ΔHm) and the cold crystallization process (ΔHcc) were obtained from the second heating step, whereas the temperature of crystallization (Tc) and enthalpy of the crystallization process (ΔHc) were obtained from the cooling step. The degree of crystallinity (χc) was calculated using Equation (2):

$$\% \chi\_{\text{c}} = \frac{\left(\Delta \text{H}\_{\text{m}} - \Delta \text{H}\_{\text{cc}}\right)}{\Delta \text{H}\_{100\%} \cdot \text{W}\_{\text{P}}} \cdot 100 \tag{2}$$

where Wp stands for the weight fraction of PET (%) in the sample and ΔH100% is the melting enthalpy of a theoretic 100% crystalline PET, that is, 140 J·g−<sup>1</sup> [47,48].

Thermal degradation was evaluated by thermogravimetry (TGA) in a TGA/SDTA851 thermobalance from Mettler-Toledo Inc (Schwarzenbach, Switzerland). Samples with an average weight of 5–7 mg were subjected to a temperature sweep from 30 ◦C to 700 ◦C at a constant heating rate of 20 ◦C·min−<sup>1</sup> in air atmosphere with a flow-rate of 50 mL·min−1. The onset degradation temperature, which was assumed at a weight loss of 5 wt% (T5%), and the maximum degradation rate temperature (Tdeg) were obtained from the TGA curves. All thermal tests were run in triplicate.

#### *2.7. Thermomechanical Characterization*

Dynamical mechanical thermal analysis (DMTA) was conducted in an oscillatory rheometer AR-G2 from TA Instruments (New Castle, DE, USA), equipped with a special clamp system for solid samples working in a combination of shear-torsion stresses. The maximum shear deformation (%γ) was set to 0.1% at a frequency of 1 Hz. The thermal sweep was from 20 ◦C to 160 ◦C at a heating rate of 5 ◦C min<sup>−</sup>1. Thermomechanical analysis (TMA) was carried out in a Q-400 thermoanalyzer, also from TA Instruments. A dynamic thermal sweep from 0 ◦C to 140 ◦C at a constant heating rate of 2 ◦C·min−<sup>1</sup> with an applied load of 20 mN was performed. All thermomechanical tests were also done in triplicate.

#### **3. Results and Discussion**

#### *3.1. Mechanical Properties of the Bio-PET*/*r-PET Blends*

Table 2 shows the mechanical properties in terms of Et, σmax, εb, Shore D hardness, and impact strength for bio-PET, r-PET, and their blends. One can observe that neat bio-PET resulted in elastic and flexible pieces, having Et and σmax values of approximately 600 MPa and 57 MPa, respectively, whereas <sup>ε</sup><sup>b</sup> nearly reached a value 495%. Shore D hardness and impact strength values were 92.3 kJ·m−<sup>2</sup> and 2.87 kJ·m−2, respectively. These mechanical properties were similar to those reported in our earlier research work [45], though the present pieces were slightly less rigid and more flexible. The performance differences could be related to variations in the moisture content since the PET polyester can easily absorb water due to the high hygroscopicity that plasticizes it [49]. In contrast, neat r-PET produced pieces with a similar mechanical strength but a significantly lower mechanical ductility and toughness. In particular, <sup>ε</sup><sup>b</sup> was 23.1%, while the impact strength was 0.82 kJ·m−2, that is, a respective 21- and 3.5-fold reduction in comparison with virgin bio-PET. This ductility impairment has been associated with the decrease in the PET's MW, mainly caused by a phenomenon of chain scission [50,51]. Therefore, as expected, increasing the amount of r-PET in the blend induced a progressive reduction in the ductile performance of the pieces, reaching ε<sup>b</sup> values in the 10–11% range for the highest contents. The incorporation of r-PET in contents of 30 wt% and 45 wt%, however, increased the E value by approximately 40%. This mechanical behavior has been previously ascribed by Mancini and Zanin to an increase in crystallinity [50] but it can also be associated with the attained reduction in ductility.


**Table 2.** Summary of the mechanical properties of the partially bio-based and recycled polyethylene terephthalate (bio-PET/r-PET) blends in terms of tensile modulus (Et), maximum tensile strength (σmax), elongation at break (εb), Shore D hardness, and impact strength.

The incorporation of 1 phr of PS-*co*-GMA into the blend system had a slight influence on the mechanical performance of the bio-PET/r-PET pieces. However, the use of contents of 3 phr and 5 phr successfully yielded a significant increase in the ductility and toughness in the r-PET/bio-PET pieces. In particular, for the blend piece containing 45 wt% r-PET, ε<sup>b</sup> increased from 10.8% to 312.9% and

378.8% for the same pieces that were melt-processed with 3 phr and 5 phr of PS-*co*-GMA, respectively. The same trend was observed for the impact strength values, increasing from 1.84 kJ·m−<sup>2</sup> to 2.43 kJ·m−<sup>2</sup> and 2.52 kJ·m<sup>−</sup>2, respectively. This improved ductility can be ascribed to the chain extension mechanism of PET achieved during extrusion by the presence of the PS-*co*-GMA. This reactive copolymer contains multiple GMA groups that can react with the –OH and –COOH terminal groups of both bio-PET and r-PET. As a result of the so-called reactive extrusion process, a macromolecule of higher MW based on a linear chain-extended, branched, or even cross-linked structure is formed. The proposed chain extension mechanism is shown in Figure 1. Thus, the resultant material shows a macromolecular structure with a higher degree of entanglements to resist mechanical deformation. A similar mechanical improvement was observed, for instance, by the addition of diisocyanates, in which the ε<sup>b</sup> value notably increased from approximately 5% to 300% [34]. Similarly, maleinized hemp seed oil (MHO) and acrylated epoxidized soybean oil (AESO) have been used to increase the mechanical performance of PLA materials [52].

**Figure 1.** Proposed mechanism of chain extension of polyethylene terephthalate (PET) with poly(styrene -*co*-glycidyl methacrylate) (PS-*co*-GMA).

#### *3.2. Morphologgy and Appareance of the Bio-PET*/*r-PET Blends*

Figure 2 shows the appearance of the injection-molded neat bio-PET, r-PET, and their blends. Color changes of the bio-PET/r-PET pieces were quantified by the L\*a\*b\* coordinates and are reported in Table 3. PET is a semi-crystalline polyester that can be manufactured during injection molding into articles of different transparencies by selecting the appropriate cooling conditions, being very transparent when full amorphous or opaque when it is highly crystallized [53]. The neat bio-PET piece presented a natural bright color, but it was opaque, indicating that the biopolymer developed a certain crystallinity during cooling in the injection mold and in the subsequent annealing. The r-PET pieces were also opaque, but slightly grey and less bright. One can observe that the bio-PET pieces developed a yellowish color (b\* > 0) as the weight percentage of r-PET increased, being also opaque. This effect has been described by Torres et al. [29], who attributed it to the presence of spherulitic crystallization as a result of the addition of r-PET. One can also observe that the introduction of PS-*co*-GMA reduced the yellow tonality in the pieces, though the effect of the random copolymer on their visual appearance was relatively low. Therefore, although a different color was noticed (ΔE\* ≥ 5) in all the pieces, the color differences were significantly less noticeable after the incorporation of PS-*co*-GMA.

**Figure 2.** Visual aspect of the injection-molded pieces of the partially bio-based and recycled polyethylene terephthalate (bio-PET/r-PET) blends corresponding, from left to right, to: B100, R100, B85-R15, B70-R30, B55-R45, B55-R45-X1, B55-R45-X3, and B55-R45-X5.



Figure 3 presents the FESEM images corresponding to fracture surfaces of the injection-molded pieces after the impact tests. All morphologies showed the formation of microcracks that were responsible for the final fracture of the material. Figure 3a, which corresponds to the FESEM micrograph of the neat bio-PET piece, shows that the size of the cracking steps was larger with somewhat more rounded edges compared with those of the r-PET piece shown in Figure 3b. This morphological change is due to more energy being required during the breakage process so that the fracture produced a rougher surface. In both cases, the cracking steps were uniform. In contrast, the FESEM micrographs corresponding to the bio-PET/r-PET blends, shown in Figure 3c–e, yielded very heterogeneous fracture surfaces. This type of fracture can be mainly ascribed to the hardness increase of the injection-molded piece [54]. Furthermore, although the morphologies revealed the presence of a single phase, this type of surface can also be ascribed to the lack of chemical interaction between the matrix, that is, bio-PET, and the dispersed domains of r-PET. One can observe in Figure 3f that a similar morphology was obtained when the polyester blend was melt-processed with 1 phr of PS-*co*-GMA. However, as seen in Figure 3g,h, the addition of the 3 phr and 5 phr of PS-*co*-GMA produced pieces with fracture surfaces very similar to those observed for neat bio-PET. In particular, one can observe that larger cracks with rounded edges were produced due to the aforementioned increase in toughness. Therefore, the fracture surfaces correlated well with the mechanical strength and impact strength performance attained during the mechanical analysis. A similar morphology was

reported by Badia et al. [55], who showed a smoother surface loss in the surface fractures of virgin PET after the addition of r-PET. This behavior was also noted in other recycled polymers [56,57].

**Figure 3.** Field emission scanning electron microscopy (FESEM) images, taken at 500×, corresponding to fracture surfaces of the injection-molded pieces of the partially bio-based and recycled polyethylene terephthalate (bio-PET/r-PET) blends: (**a**) B100; (**b**) R100; (**c**) B85-R15; (**d**) B70-R30; (**e**) B55-R45; (**f**) B55-R45-X1; (**g**) B55-R45-X3; (**h**) B55-R45-X5. Scale markers of 10 μm.

#### *3.3. Thermal Characterization of the Bio-PET*/*r-PET Blends*

The thermal transitions and degree of crystallinity are parameters of great technological importance that reflect the chemical structure of the polymer and they can be correlated with its mechanical properties [58]. Figure 4 shows a comparative plot of the DSC thermograms of the bio-PET/r-PET blends during the cooling step (Figure 4a) and second heating step (Figure 4b). The main thermal properties obtained from the DSC analysis are summarized in Table 4. In relation to the bio-PET sample, the glass transition region was seen as a step in the baseline located around 82 ◦C in the heating thermogram. Then, one can observe in the cooling thermogram that the biopolyester did not crystallize from the melt but developed a cold crystallization process during heating in the thermal range between 140 ◦C and 180 ◦C, showing a maximum exothermic peak, the so-called Tcc, at 161 ◦C. Finally, shown in the heating thermogram, the endothermic peak between 220 ◦C and 260 ◦C corresponds to the melting process of the total crystalline fraction in the biopolyester, showing a Tm value of nearly 245 ◦C. The thermal parameters attained for bio-PET were very similar to those obtained in our previous work [45]. A similar thermal profile was also observed for r-PET, with the most significant difference being the lower Tcc value observed, which was located at approximately 150 ◦C. This observed value can be ascribed to the lower MW of r-PET, which compromises shorter chains that can more easily cold crystallize due to the fact that the polyester was subjected to a second round of processing [59].

Interestingly, the bio-PET/r-PET blends showed a different thermal profile during DSC analysis. One can observe that the polyester blends crystallized in all cases during cooling, showing values of Tc in the 180–190 ◦C range. This observation suggests that crystallization was thermodynamically favored in the PET blends. This phenomenon can be ascribed to the potential role of r-PET as a nucleating agent in bio-PET. In this regard, the nucleating effect of r-PET has been previously related to the presence of impurities [29]. Therefore, the spherulitic crystallization of r-PET occurs at lower temperatures and the crystals formed can thereafter promote the homogenous crystallization of PET chains at higher temperatures. This effect was further confirmed by the increase in the Tc values and the higher intensity of the exothermic peaks observed during crystallization from the melt with higher r-PET contents. Thus, χ<sup>c</sup> value increased from 10.6% in the neat bio-PET sample to 29% for the bio-PET blend containing 15 wt% of r-PET, whereas these values reached 32.3% and 37% for the blends containing 30 wt% and 45 wt% of r-PET, respectively. It is also worth mentioning that while the melting process of the neat bio-PET sample occurred in a single melt peak, all the bio-PET/r-PET blends showed two overlapping peaks during melting. This double-melting peak phenomenon can be explained by the presence of dominant and subsidiary crystals [60]. The two melting peaks (Tm1 and Tm2) were observed at approximately 238 and 248 ◦C, with the latter temperature being assigned to primary folded chain crystals with a higher melting point than those observed for the neat bio-PET. This observation further confirms that the presence of r-PET in bio-PET favors the formation of more perfect crystalline structures, which were obtained from the melt of imperfect structures or crystals with lower lamellae thicknesses [61]. Similar results were observed by Spinacé and De Paoli [62], who indicated that the chain scission, occurring during the processing of PET, improved chain packing, increasing the crystallite size and consequently shifting Tc and Tm to higher values.

The incorporation of PS-*co*-GMA barely modified the Tg, Tm1, and Tm2 values, though the use of contents of 3 phr and 5 phr decreased the crystallinity to 17.7% and 25.9%, respectively. The reactive random copolymer also caused a slight reduction in the Tc values, showing values around 185 ◦C, indicating that the structural disorder introduced by the chain extender potentially hindered both nucleation and crystallization. These observations agree with the aforementioned assumption that the induced MW increase could decrease the molecular mobility, resulting in a low crystallization rate and lower crystallinities [63]. Moreover, Kiliaris et al. [64] also concluded that the ability of the PET chains to crystallize in folded lamellae was limited after chain extension, which led to the formation of smaller and less perfect crystallites. Furthermore, if one considers branching as the prevailing chain-extension reaction, the presence of the resultant branching points could effectively hinder chain symmetry, restricting the segment movements and yielding crystal defects.

**Figure 4.** Differential scanning calorimetry (DSC) curves during cooling (**a**) and second heating (**b**) of the injection-molded pieces of the partially bio-based and recycled polyethylene terephthalate (bio-PET/r-PET) blends.

**Table 4.** Summary of the main thermal properties of the injection-molded pieces of the partially bio-based and recycled polyethylene terephthalate (bio-PET/r-PET) blends in terms of glass transition temperature (Tg), cold crystallization temperature (Tcc), crystallization temperature (Tc), melting temperature (Tm), and degree of crystallinity (χc).


One of the main problems related to the mechanical recycling of thermoplastic materials is the loss of thermal stability. Figure 5 gathers the TGA curves (Figure 5a) and the first derivatives of the curves (DTG) (Figure 5b) of the bio-PET/r-PET blends, while Table 5 presents the thermal values obtained from the TGA curves. It can be observed that bio-PET degradation occurred in two stages, as previously reported by Vannier et al. [65]. The first stage, occurring from 350 ◦C to 460 ◦C, corresponds to the main degradation of the biopolymer backbone and the formation of char with a mass loss of ~83%. The second step, which was associated with a loss of nearly the totality of the remaining biopolyester, was related to the thermo-oxidative degradation of the char. This occurred from 490 ◦C to 560 ◦C. It has also been reported that the decomposition mechanism of PET consists of an hererolytic scission via a six-membered ring intermediate, where the hydrogen from a β-carbon to the ester group is transferred to the ester carbonyl, followed by scission at the ester links [66]. Dzie¸cioł and Trzeszczy ´nski [67,68] also reported that the degradation of PET leads to the formation of acetaldehyde, oligomers with terminal carboxyl groups, CO, and CO2, among other compounds. In all cases, the bio-PET/r-PET samples resulted in a residual mass of 1–2 wt% at 700 ◦C.

One can also observe that the thermal degradation profile of r-PET and the bio-PET/r-PET blends was similar to that of the neat bio-PET. The addition of PS-*co*-GMA yielded an increase in the onset temperature of degradation that was measured as the degradation temperature at 5% of mass loss, that is, T5%. In particular, the beginning of the thermal degradation was delayed from approximately 393 ◦C for the blends containing 30 wt% and 45 wt% of r-PET to nearly 403 ◦C. The thermal degradation peaks (Tdeg1 and Tdeg2), determined when the maximal degradation rates were produced, remained nearly constant during the first and main mass loss, while they slightly increased during the second one in the PET blends processed without PS-*co*-GMA. In particular, the Tdeg2 values were increased by up to nearly 19 ◦C, which suggests a delay in the thermo-oxidative degradation of the char. In any case, the influence of the reactive copolymer on the thermal degradation of the bio-PET/r-PET blends was relatively low, since the thermal stability of the polyester blends was already high.


**Table 5.** Summary of the main thermal properties of the injection-molded pieces of the partially bio-based and recycled polyethylene terephthalate (bio-PET/r-PET) blends in terms of the degradation temperature at 5% of mass loss (T5%), degradation temperature (Tdeg), and residual mass at 700 ◦C.

**Figure 5.** (**a**) Thermogravimetric analysis (TGA) and (**b**) first derivative thermogravimetric (DTG) curves of the injection-molded pieces of the partially bio-based and recycled polyethylene terephthalate (bio-PET/r-PET) blends.

#### *3.4. Thermomechanical Properties of the Bio-PET*/*r-PET Blends*

The prepared bio-PET/r-PET blends and the effect of the reactive random copolymer on these polyester blends were analyzed in terms of their thermomechanical properties by means of DMTA and TMA. Figure 6 shows the DMTA curves of the bio-PET/r-PET blend pieces. Figure 6a shows the storage modulus as a function of temperature. The evolution of the storage modulus of the bio-PET was characterized by a single thermal transition with a temperature increase in the 50–130 ◦C range. Up to approximately 70 ◦C, the variation in the storage modulus values was nearly negligible since the biopolyester was in a glassy state. At room temperature, all the pieces showed similar storage modulus values, being around 1 GPa and slightly lower for the neat bio-PET due to its reduced

crystallinity. Thereafter, one can observe a remarkable decrease in the storage modulus values in the temperature range from 70 ◦C to 85 ◦C. This process notably involved a decrease of more than two orders of magnitude in the storage modulus, which is representative for the glass-to-rubber transition of bio-PET. At higher temperatures, the variation in the storage modulus was nearly negligible, indicating that the biopolyester reached its rubber state and that further temperature increase did not affect the mechanical properties. Table 6 presents the values of the storage modulus obtained at 60 ◦C and 100 ◦C since these temperatures are representative of the mechanical rigidity of the bio-PET pieces before and after the glass transition region of bio-PET. While at 60 ◦C, the value of the storage modulus of the neat bio-PET piece was nearly 900 MPa, this value was reduced to approximately 3 MPa at 100 ◦C. It can also be observed that the r-PET piece and all the bio-PET/r-PET pieces showed similar thermochemical performance with slightly higher storage modulus values. The incorporation of PS-*co*-GMA led to an increase in the storage modulus values, in both the glassy and the rubber states. The highest enhancement was obtained for the bio-PET/r-PET blend piece processed with 5 phr of PS-*co*-GMA, in which it reached values of close to 1.2 GPa and 30 MPa at 60 ◦C and 100 ◦C, respectively. This thermomechanical increase can be ascribed to the aforementioned formation of a chain-extended macromolecular structure with higher MW that showed increased resistance to flow during the application of external forces as a result of the large number of entanglements associated with the long chains or branches [69].

Figure 6b shows the evolution of the damping factor (*tan* δ) as a function of temperature. The maximum peak of the *tan* δ curves corresponds to alpha (α)-transition of bio-PET, which is related to its Tg. These values are also given in Table 6, where it can be seen that for neat bio-PET, a Tg value of approximately 81 ◦C was observed, being nearly the same as that obtained above by DSC. A slight lower value of Tg, that is, ~80 ◦C, was observed for r-PET, with this potentially being related to its lower chain lengths and/or the presence of plasticizers. The bio-PET/r-PET blends showed intermediate values of Tg, ranging between 80 ◦C and 81 ◦C, being also similar to those obtained by DSC. The biopolyester blends melt-processed with PS-*co*-GMA showed nearly the same Tg values but with a remarkable reduction in the α-peak intensities. This observation supports the higher crystallinity and the large number of entanglements achieved in the polyester blends based on the fact that the amorphous phase content was reduced due to the presence of r-PET, which partially suppressed the relaxation of the bio-PET chains and lowered the number of molecules undergoing α-transition [70].


**Table 6.** Storage modulus measured at 60 ◦C and 100 ◦C, glass transition temperature (Tg), and coefficient of linear thermal expansion (CLTE) of the injection-molded pieces of the partially bio-based and recycled polyethylene terephthalate (bio-PET/r-PET) blends.

**Figure 6.** (**a**) Storage modulus and (**b**) damping factor (*tan* δ) of the injection-molded pieces of the partially bio-based and recycled polyethylene terephthalate (bio-PET/r-PET) blends.

Finally, the dimensional stability the bio-PET/r-PET blend pieces was determined by measuring the CLTE values below and above the glass transition region. As also shown in Table 6, below Tg, bio-PET showed a lower value than r-PET, that is, 78.9 <sup>μ</sup>m·m−1· ◦C−<sup>1</sup> versus 90.9 <sup>μ</sup>m·m−1· ◦C−1. This thermomechanical difference can be ascribed to the lower MW attained in the recycled polyester due to the aforementioned phenomenon of chain scission during reprocessing. Then, the bio-PET/r-PET blends presented intermediate values according to their r-PET content. The incorporation of PS-*co*-GMA induced a reduction in the CLTE values, increasing with the reactive random copolyester content. In particular, a value of 70.6 <sup>μ</sup>m·m−1· ◦C−<sup>1</sup> was reached for the bio-PET/r-PET blend containing 5 phr PS-*co*-GMA. This observation further supports the formation of a branched and larger macromolecule that reduced the effect of temperature on the dimensional stability, which is also in agreement with the DMTA results. A similar trend with more significant differences was observed for the CLTE values measured above Tg.

#### **4. Conclusions**

PS-*co*-GMA, a multi-functional random copolymer, was melt-mixed with bio-PET/r-PET blends at contents of 1–5 phr of polyester blend and shaped into pieces by injection molding. The resultant pieces were characterized to ascertain the potential use of PS-*co*-GMA as a chain extender for the mechanical recycling of these polyester blends. The results showed that while the incorporation of 1 phr of PS-*co*-GMA had a slight influence on the mechanical and thermal performance of the bio-PET/r-PET blend pieces, the contents of 3 phr and 5 phr successfully yielded a significant increase in their ductility and toughness. In particular, ε<sup>b</sup> increased from 10.8%, for the blend piece containing 45 wt% of r-PET, to 312.9% and 378.8%, for the same pieces that were melt-processed with 3 phr and 5 phr of PS-*co*-GMA, respectively. In addition, the impact strength values increased from 1.84 kJ·m−<sup>2</sup> to 2.43 kJ·m−<sup>2</sup> and 2.52 kJ·m−2, respectively. This mechanical improvement was ascribed to the chain extension mechanism of PET by reactive extrusion due to the reaction of the multiple GMA groups that are present in PS-*co*-GMA with the –OH and –COOH terminal groups of both bio-PET and r-PET. Furthermore, PS-*co*-GMA reduced the crystallinity of the PET blends, suggesting that the structural disorder introduced by the chain extender potentially hindered both the nucleation and crystallization of r-PET on the bio-PET chains, contributing to increase material flexibility. In addition, the influence of the reactive copolymer on the thermal stability of the bio-PET/r-PET blends was low, since the blends were already thermally stable up to nearly 400 ◦C. The thermomechanical properties of the bio-PET/r-PET blends also improved after the addition of PS-*co*-GMA due to the formation of a branched and larger macromolecule.

Bio-PET can be effectively mixed with its recycled petrochemical counterpart, that is, r-PET, and then mechanically recycled in existing recycling facilities by means of reactive extrusion with PS-*co*-GMA. Therefore, the original biopolymer properties can be successfully restored and the ultimate performance of bio-PET articles would be retained for a given number of reprocessing cycles. This will permit the recovery of upcoming bio-PET streams with current r-PET waste to manufacture the same or similar products. According to this scenario, mechanical recycling for bio-based but non-biodegradable polymers will be appropriate from both an economic and environmental point of view. This will potentially contribute to accelerating the transition of the plastic packaging industry from its traditional linear model to a more valuable and sustainable circular model.

**Author Contributions:** Conceptualization was devised by N.M., S.F.-B., and S.T.-G.; data curation and writing-original draft by S.M.-J., L.Q.-C., and T.B.; methodology, validation, and formal analysis by S.M.-J., D.L., and A.V.M.-S.; Investigation, resources, Writing—Review and editing by S.T.-G.; project administration, S.F.-B. and S.T.-G. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research work was funded by the Spanish Ministry of Science, Innovation, and Universities (MICIU) project numbers RTI2018-097249-B-C21 and MAT2017-84909-C2-2-R.

**Acknowledgments:** L.Q-C. wants to thank Generalitat Valenciana (GVA) for his FPI grant (ACIF/2016/182) and the Spanish Ministry of Education, Culture, and Sports (MECD) for his FPU grant (FPU15/03812). D.L. thanks Universitat Politècnica de València (UPV) for the grant received through the PAID-01-18 program. S.T.-G. is recipient of a Juan de la Cierva contract (IJCI-2016-29675) from MICIU. S.R.-L. is recipient of a Santiago Grisolía contract (GRISOLIAP/2019/132) from GVA. J.I.-M. wants to thank UPV for an FPI grant PAID-01-19 (SP2019001). Microscopy services at UPV are acknowledged for their help in collecting and analyzing FESEM images. Authors also thank Polyscope Polymers B.V. for kindly supplying XibondTM 920.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


© 2020 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

*Article*

## **Functionalization of Partially Bio-Based Poly(Ethylene Terephthalate) by Blending with Fully Bio-Based Poly(Amide) 10,10 and a Glycidyl Methacrylate-Based Compatibilizer**

#### **Maria Jorda 1, Sergi Montava-Jorda 2, Rafael Balart 1, Diego Lascano 1,3,\*, Nestor Montanes <sup>1</sup> and Luis Quiles-Carrillo <sup>1</sup>**


Received: 19 July 2019; Accepted: 9 August 2019; Published: 10 August 2019

**Abstract:** This work shows the potential of binary blends composed of partially bio-based poly(ethyelene terephthalate) (bioPET) and fully bio-based poly(amide) 10,10 (bioPA1010). These blends are manufactured by extrusion and subsequent injection moulding and characterized in terms of mechanical, thermal and thermomechanical properties. To overcome or minimize the immiscibility, a glycidyl methacrylate copolymer, namely poly(styrene-ran-glycidyl methacrylate) (PS-GMA; Xibond™ 920) was used. The addition of 30 wt % bioPA provides increased renewable content up to 50 wt %, but the most interesting aspect is that bioPA contributes to improved toughness and other ductile properties such as elongation at yield. The morphology study revealed a typical immiscible droplet-like structure and the effectiveness of the PS-GMA copolymer was assessed by field emission scanning electron microcopy (FESEM) with a clear decrease in the droplet size due to compatibilization. It is possible to conclude that bioPA1010 can positively contribute to reduce the intrinsic stiffness of bioPET and, in addition, it increases the renewable content of the developed materials.

**Keywords:** bio-based; poly(ethyelene terephthalate)—PET; poly(amide) 1010—PA1010; mechanical properties; morphology; compatibilization; Xibond™ 920

#### **1. Introduction**

In the last decade, there has been a noticeable increase in the sensitiveness and concern about environment. Topics such as sustainable development, circular economy, carbon footprint, petroleum depletion, among others are gaining relevance [1–3]. Therefore, many research works are focused on the development of environmentally friendly materials to positively contribute to a sustainable development. This situation is particularly aggravated in the polymer industry which accounts for the use of large amounts of petroleum-derived plastics with the subsequent environmental impact both at the origin (petroleum) and at the end of the life cycle or disposal (most of the petroleum-based polymers are not biodegradable). For these reasons, the polymer industry is demanding continuously environmentally friendly polymers It is worthy to note the important role that some petroleum-based polymers have acquired in the last decade. In particular, aliphatic polyesters such as poly(ε-caprolactone) (PCL) [4] poly(butylene succinate) (PBS) [5], poly(glycolic acid) (PGA) [6], poly(butylene succinate*-co-*adipate) (PBSA) and their blends/composites with other polymers and lignocellulosic fillers have gained interest in several industrial sectors, despite being petroleum-based, as they can undergo degradation under controlled compost soil [7–9]. Another promising group of environmentally friendly polymers includes polysaccharides (and derivatives), protein-based polymers and bacterial polymers. Poly(lactic acid) (PLA), together with thermoplastic starch (TPS), are perhaps the most studied polymers in this group that can be derived from polysaccharides [10], in particular, from starch-rich materials, i.e., potato, corn, bagasse, and so on. PLA is commercially available at a competitive price. Protein-based polymers include some interesting materials as gluten, soy protein, collagen, rape-seed protein, among others, that find applications in the form of film, parts, fibers, and so on [11–14]. Finally, bacterial polymers include all poly(hydroxyalkanoates) (PHAs) which are expected to invade the market soon [15,16]. Some of the most interesting PHAs include poly(3-hydroxybutyrate) (PHB), poly(3-hydroxybutyrate*-co-*valerate) (PHBV), poly(3-hydroxybutyrate*-co-*hexanoate) (PHBH) [17,18].

Despite all of these materials representing a clear environmental efficiency, in general, their properties are far from those of petroleum-derived commodity and engineering plastics. For this reason, many studies have been focused on obtaining commodity and engineering plastics from renewable resources. These show identical properties to their petroleum-based counterparts, but they offer interesting environmental efficiency as they can be totally or partially derived from renewable materials, usually bio-products coming from the food industry and agroforestry. Bio-based poly(ethylene) (bioPE), is a commodity that is synthesised from bioethanol from sugarcane and can reach almost 95% bio-based content. This shows a clear positive environmental efficiency compared to poly(ethylene) from crude oil [19–21]. Currently, bioPE is available worldwide at a relatively cost competitive price and it has been recently used as a base material for 3D printing [22]. In regard to engineering plastics, it is worthy to note the increasing consumption of bio-based poly(ethylene terephthalate) (bioPET) and bio-based poly(amides) (PAs) [23–26]. In recent years, poly(ethylene furanoate) (PEF) has generated great expectations as it can be fully bioderived and could potentially substitute poly(ethylene terephthalate) (PET) polymers [27]. Although in the future bioPET can reach 100% renewable source since there is a bio-route to synthesise terephthalic acid (TA) [28,29], currently its bio-based content is related to the ethylene glycol which can give approximately 30% bio-based content. Regarding bioPAs, castor oil plays a key role as a starting material for PA synthesis [30]. It is worthy to note bioPAs are engineering plastics with different bio-sourced content. Thus, bioPA610 typically offers 60–63% renewable content [31]; bioPA1012 usually offers a renewable content of 45% and bioPA1010 can be 100% derived from renewable resources from sebacic acid and 1,10-decamethylene diamine (DMDA), both derived from ricinoleic acid [32]. Some of these bioPAs have alternative eco-routes and could be fully bioderived. PET and, recently bioPET are widely used in the packaging industry for bottles. Despite this, some beverages (especially oxygen-sensitive beverages) require the use of scavengers that usually are derived from poly(amides), so that poly(amides) are increasingly present in the PET bottle-to-bottle cycle [33].

This work explores the potential of high bio-based content blends of partially bio-based poly(ethylene terephthalate) (bioPET) and fully bio-based poly(amide) 10,10 (bioPA1010) up to 30 wt %. Although bioPET and bioPA1010 show similar properties to their corresponding petroleum-derived counterparts, currently bioPET only contains approximately 30 wt % of biobased content while bioPA10120 can be fully bioderived from castor oil. The production of these partially or totally biobased polymers is increasing in a remarkable way and new biobased routes are being developed for partially biobased polymers to achieve 100% biosourced materials. This can have a positive effect on sustainable development and circular economies as most of the biobased building blocks could be obtained from by-products of the food or agroforestry industries. For these reasons, blending these two polymers is attractive from an environmental standpoint as these blends could reach high biobased contents without compromising other properties, thus leading to engineering blends with potential in the packaging industry. Due to their immiscibility, a glycidyl compatibilizer, namely a poly(styrene-ran-glycidyl methacrylate) copolymer (PS-GMA) Xibond™ 920 was used. The effect

of both bioPA1010 and the PS-GMA compatibilizer are evaluated on their mechanical, thermal and thermomechanical properties. The novelty of this work is the high renewable content that can be obtained by these blends together with improved toughness.

#### **2. Experimental**

#### *2.1. Materials*

The partially bio-based poly(ethylene terephthalate), bioPET and fully bio-based poly(amide) 1010 were supplied by NaturePlast (Ifs, France). Table 1 summarizes the main properties of these commercial grades.

**Table 1.** Commercial grades and main properties of partially bio-based poly(ethylene terephthalate)—bioPET and fully bio-based poly(amide) 1010, supplied by NaturePlast.


\* measured between 230–240 ◦C.

The selected compatibilizer was a poly(styrene-glycidyl methacrylate) random copolymer (PS*-*GMA) Xibond™ 920 and was kindly provided by Polyscope (Geleen, The Netherlands). The GMA functionality has excellent affinity with polycondensates which can result in compatibilization, chain extension and/or branching. Figure 1 shows a scheme of the different materials used in this research.

**Figure 1.** Schematic representation of the chemical structure of bio(polyethylene terephthalate), bio-based poly(amide) 10,10 and glycidyl copolymer compatibilizer Xibond™ 920.

#### *2.2. Manufacturing of Binary BioPET*/*BioPA Blends*

Initially, all materials (see Table 2 for code and composition) were dried at 60 ◦C for 24 h to remove residual moisture. After this, the corresponding amounts of each material were mechanically mixed in a zipper bag, and then were fed into the hopper of a twin-screw co-rotating extruder from DUPRA S.L. (Castalla, Spain). The screw diameter was 30 mm and the temperature profile was set to four different barrels as follows (from the hopper to the die): 250 ◦C, 260 ◦C, 260 ◦C and 260 ◦C. The rotating speed was set to 20 rpm. After this initial compounding stage, the obtained blends were cooled down to room temperature and subsequently pelletized for further processing by injection moulding. The injection moulding machine used was a Mateu & Solé mod. Meteor 270/75 (Barcelona, Spain). The temperature profile was 240 ◦C (feeding Hopper), 245 ◦C, 250 ◦C and 255 ◦C (injection nozzle). The filling time was set at 1 s and the cooling time was 5 s.

**Table 2.** The compositions and labeling of binary bioPET/bioPA1010 blends. The bio-based content is calculated considering that bioPET contains an average bio-based content of 30 wt %, bioPA1010 is 100% bio-based and Xibond™ 920 is petroleum-derived (0 wt % bio-based).


\* phr: weight grams of Xibond™ 920 per one hundred grams bioPET/bioPA blend.

#### *2.3. Mechanical Characterization*

The tensile properties were obtained through ISO 527-2:2012 standard on injection moulded dog-bone samples using an electromechanical machine ELIB-50 from S.A.E Ibertest (Madrid, Spain). All tests were run at a cross-head speed of 10 mm·min<sup>−</sup>1, using a 5 kN loadcell. Regarding the impact strength, it was estimated through a Charpy test using a 6-J pendulum from Metrotec S.A. (San Sebastián, Spain) on the unnotched rectangular samples, following indications of ISO 179-1:2010. Finally, the hardness was obtained by using the Shore method in a 673-D durometer from J. Bot Instruments (Barcelona, Spain) as suggested by ISO 868:2003. All mechanical tests were run at room temperature and at least five different samples were tested to obtain the average characteristic parameters.

#### *2.4. Thermal Characterization*

Differential scanning calorimetry (DSC) was used to study the main thermal transitions of the manufactured materials. DSC tests were carried out on a Q200 calorimeter from TA Instruments (New Castle, DE, USA). A dynamic thermal program was scheduled in three different stages using standard aluminium crucibles. The first heating from 30 ◦C up to 280 ◦C was followed by a cooling down to 0 ◦<sup>C</sup> and a second heating up to 350 ◦C. The heating/cooling rate was set to 10 ◦C·min−<sup>1</sup> with a constant nitrogen flow rate of 50 mL min<sup>−</sup>1. The maximum degree of crystallinity was calculated for both bioPET and bioPA (see Equation 1) by comparing the melt enthalpy (Δ*Hm*) with the corresponding melt enthalpy of a theoretical 100% crystalline polymer (*Hm*<sup>0</sup> for PET <sup>=</sup> 140.1 J·g−<sup>1</sup> [34], and 244.0 J·g−<sup>1</sup> for PA1010 [35]), and considering the weight fraction of each polymer in the blend (*w*).

$$\% \chi\_{\mathcal{L}} = \frac{\Delta H\_m}{\Delta H\_m^0 \cdot w} \tag{1}$$

Additional thermal characterization was carried out by thermogravimetry (TGA) in a TGA/SDTA 851 thermobalance from Mettler-Toledo (Schwerzenbach, Switzerland). A dynamic heating program from 20 ◦<sup>C</sup> to 700 ◦<sup>C</sup> at a heating rate of 20 ◦C·min−<sup>1</sup> was applied to an average sample weight of 8 mg in an air atmosphere in alumina crucibles.

#### *2.5. Morphology Characterization*

Field emission scanning electron microscopy (FESEM) was used to reveal the morphology of the fractured surfaces blends after the impact tests. A FESEM microscope from Oxford Instruments (Abingdon, UK) was used working at an acceleration voltage of 1.5 kV. As the polymer blends were not electrically conducting materials, a metal sputtering process was carried out to provide conducting properties to the samples and to avoid sample charge. All fractured samples were coated with an ultrathin gold-palladium alloy in a Quorum Technologies Ltd. EMITECH model SC7620 sputter coater (East Sussex, UK).

#### *2.6. Thermo-Mechanical Characterization*

The effect of the temperature on the mechanical properties and dimensional stability was studied by dynamic-mechanical thermal analysis (DMTA) and thermomechanical analysis (TMA) respectively. Thermomechanical analysis was carried out in a Q400 thermoanalizer from TA Instruments (New Castle, DE, USA). The particular conditions for this test were a dynamic thermal sweep from 0 ◦C up to 140 ◦C at a constant heating rate of 2 ◦C·min−<sup>1</sup> with an applied load of 20 mN. Regarding dynamic-mechanical thermal characterization, an oscillatory rheometer AR-G2 from TA Instruments (Delaware, USA), equipped with a special clamp system for solid samples (working in a combination of shear-torsion stresses) was used. The maximum shear deformation (%) was set to 0.1% at a frequency of 1 Hz. The thermal sweep was scheduled from 20 ◦C up to 160 ◦C at a heating rate of 5 ◦C·min<sup>−</sup>1.

#### **3. Results and Discussion**

#### *3.1. Mechanical Properties and Morphology of Binary BioPET*/*BioPA Blends*

It has been reported that mechanical properties of PET are highly dependent on the processing conditions [36,37]. In addition, the mechanical properties of PET polymers are also dependent on the thermal treatment, quenching, annealing, etc. As it can be seen in Figure 2, the mechanical and thermal properties of bioPET are highly dependent on the annealing time. To assess this, bioPET has been subjected to different annealing times and studied by dynamic-mechanical thermal analysis (DMTA). Figure 2a shows the evolution of the storage modulus (*G* ) as a function of the increasing temperature. DMTA is based on the use of a dynamical time-dependent stress function, σ = σ<sup>0</sup> sin(ωt) [σ<sup>0</sup> is the maximum stress and ω represents the frequency] which produces a sinusoidal strain (ε) given by ε = ε<sup>0</sup> sin(ωt-δ) where ε<sup>0</sup> is the maximum strain and δ is the phase angle which represents the delay (viscous) properties of the material. As the modulus represents the ratio between the maximum stress to the maximum strain, then it is possible to define the complex modulus (*G*\*) as σ<sup>0</sup> = ε<sup>0</sup> *G*\* sin(ωt + δ). This expression can be expanded to give σ<sup>0</sup> = ε<sup>0</sup> *G*\* sin(ωt) cos(δ) + ε<sup>0</sup> *G*\* cos(ωt) sin(δ), which can be expressed as σ<sup>0</sup> = ε<sup>0</sup> *G* sin(ωt) + ε<sup>0</sup> *G*" cos(ωt) with *G* = *G*\* cos(δ) and *G*" = *G*\* sin(δ), thus leading to an elastic response related to *G* (storage modulus) and a viscous response related to *G*" (loss modulus). As it can be derived, the ratio between the loss modulus (*G*") to the storage modulus (*G* ) represents the damping factor or tanδ, which is directly related to the lost energy due to viscoelastic behaviour.

It can be seen that, neat bioPET shows a characteristic DMTA curve characterized by different zones. Below 60 ◦C, the storage modulus remains almost constant at a temperature range comprising between 60 ◦C and 80 ◦C, and a remarkable decrease in *G* occurs. This decrease of near three orders of magnitude is representative of the glass transition process. In addition, information is provided about the high amorphous structure due to this three-fold change. Then, at the temperature range comprising between 106 ◦C and 125 ◦C, it is possible to observe an increase in *G* which is directly related to the cold crystallization process which involves packing of polymer chains in an ordered way which gives increased stiffness. After 15 min annealing time at 110 ◦C, the morphology of the DMTA curve has changed in a remarkable way. As it can be seen, the decrease in *G* is remarkably lower. The glass transition process has shifted to higher temperatures (probably due to the restriction of chain mobility in the crystalized structure) but it is still possible to find a slight increase in *G* in the temperature range comprising between 106–120 ◦C. Nevertheless, with an annealing of 30 min, it is possible to conclude that the maximum crystallinity is achieved. The curves for 30 and 45 min are shifted to the right (higher temperatures) and the cold crystallization process has completely disappeared. Similar findings have been reported by A. Bartolotta et al. [38] who showed a remarkable change in the glass transition onset from 40 ◦C (cold-drawn PET) up to 90 ◦C in highly crystalline PET. A. Bartolotta et al. attribute this phenomenon to an increase in density on glassy domains related to the presence of more crystal-packed domains and conclude that there is a link between the chain stiffness since there is a connection between the bulk glass to the ordered structures. Z. Chen et al. [39] have also reported different crystallization mechanisms depending on the annealing temperature with remarkable changes, not only in the glass transition process but also on the melt temperature. In addition, the *G* values are higher with increasing annealing time.

Figure 2b shows the evolution of the dynamic damping factor, tan δ. Notably, the maximum damping factor value decreases with the increasing annealing time. This is consistent with the damping factor definition which shows the ratio between the loss modulus (*G*") and the storage modulus (*G* ). As the material becomes stiffer by the cold crystallization process, the denominator is higher, and this leads to lower damping factor values. It is worthy to note that the damping factor is related to the lost energy to stored energy ratio. As the annealing time increases, the material becomes stiffer and this is responsible for less lost energy. There are several methods to assess the glass transition temperature (*T*g) by using several methods from DMTA, i.e., the onset of the *G* decrease, the peak maximum of *G*" or the peak maximum of the damping factor. The peak temperature of tan δ is widely used to give accurate values of *T*g. The un-annealed bioPET shows a *T*<sup>g</sup> of 79.9 ◦C which increases progressively with increasing annealing time at 110 ◦C resulting in values of 84.8 ◦C (15 min annealing), 93.7 ◦C (30 min annealing) and 96.3–96.4 ◦C for 45 and 60 min annealing. These results are in total accordance with those reported by A. Bartolotta et al. [38], who showed a change in the onset of *T*<sup>g</sup> (using the *G* method) from 40 ◦C up to 90 ◦C.

**Figure 2.** *Cont.*

**Figure 2.** The effect of annealing on the dynamic-mechanical properties of neat bioPET subjected to different annealing times, (**a**) storage modulus, *G* and (**b**) dynamic damping factor (tan δ).

In this work, bioPET and its blends have been characterized without any annealing process, just as obtained by the injection moulding. Neat bioPET showed a tensile strength (σt) of 46.7 MPa and a relatively low elongation at yield (εy) of 3.87% which leads to a stiff material. The effect of the addition of bioPA to bioPET produces two important effects. On the one hand, the tensile strength decreases as expected in an immiscible blend as reported by K.C. Chiou et al. [40], in PA6/PBT binary blends, but the decrease is not so pronounced. In fact, the maximum percentage decrease is close to 11% for the uncompatibilized blend containing 30 wt % bioPA. It is important to remark that the bio-based content of the blend containing 30 wt % bioPA is above 50% which is a positive property from an environmental point of view. On the other hand, the addition of a flexible polymer such as bioPA, provides improved elongation at yield up to values of 4.8% which represents a percentage increase of 24%, thus leading to improved ductile behaviour. The effect of the poly(styrene*-ran-*glycidyl methacrylate) copolymer (PS*-*GMA) Xibond™ 920 gives interesting results. It is worthy to note that the addition of 3 phr Xibond™ 920 leads to higher tensile strength regarding neat bioPET, reaching values of 47.1 MPa with a parallel increase in elongation up to values of 6.09% (57% increase compared to neat bioPET and 27% the same composition without compatibilizer). This indicates that Xibond™ 920 is providing somewhat compatibilization properties to this binary system. Similar findings have been obtained using a Zn2<sup>+</sup> ionomer on PET/PA6 blends with improved elongation at break and toughness compared to an uncompatibilized blend [41]. Y. Huang et al. [42], reported the exceptional compatibilizing effect of the glycidyl group by using an epoxy resin (E-44) as a compatibilizer in PET/PA6 blends. C.T. Ferreira et al. [43] reported the potential of reactive extrusion of recycled PET and recycled PA by a reaction of the carboxyl end-chain groups in PET and the amine end-chain groups in PA with a noticeable improvement in mechanical properties using tin(II) 2-ethylhexanoate as a trans-reaction catalyst. Regarding the hardness, as both polymers show similar Shore D values, it is not possible to observe a tendency with varying composition as shown in Table 3.


**Table 3.** The mechanical properties of binary bioPET/bioPA blends obtained from tensile, hardness and Charpy tests.

Regarding the impact strength which is measured through the Charpy impact test, it is worthy to note that all the developed materials have increased impact strength in comparison to neat bioPET. Neat bioPET offers a quite brittle behaviour with an impact-absorbed energy of 23.1 kJ·m<sup>−</sup>2. The uncompatibilized binary blend with 10 wt % bioPA1010 offers an increased impact strength of 27.0 kJ·m−<sup>2</sup> (which represents a percentage increase of approximately 16.9%). This result is in total agreement with other ductile properties such as elongation at yield (εy). Y. Yan et al. [44] reported the lubricant effect of PA56 on PET blends at a molecular level with the subsequent effect on mechanical properties. This phenomenon has been observed in other binary blends composed of a brittle polymer matrix in which a rubber-like material is finely dispersed, even with poor miscibility between them. J. Urquijo et al. [45] reported a remarkable increase in both elongation at the break and impact strength in binary blends of poly(lactic acid) (PLA) and different loadings of poly(ε-caprolactone) (PCL). J. Urquijo et al. demonstrated the relevance of the elongation rate on the final elongation and, regarding the impact strength, they attributed the improvement to the small particle size of PCL-rich domains embedded in the brittle PLA-rich matrix which positively contributed to absorb energy in impact conditions. Similar findings have been reported with PLA/PCL, PHB/PCL, PLA/PBS binary blends [46–49], ternary PLA/PHB/PCL blends [50], and some poly(ester) copolymers [51]. This improved toughness is more evident in an uncompatibilzed blend containing 30 wt % bioPA1010 reaching an impact strength of 40.5 kJ·m−<sup>2</sup> (75.3% increase). The effect of the compatibilizer has a positive effect on improved toughness as it can be seen in Table 2. Xibond™ 920 is a random copolymer of poly(styrene-glycidyl methacrylate) (PS*-*GMA) and gives excellent results in compatibilizing condensation polymers. This is because of the glycidyl methacrylate group which can interact with both hydroxyl end-groups in bioPET and amine (primary or secondary) in bioPA1010, thus leading to somewhat compatibilization with a marked effect on impact strength. The addition of 5 phr Xibond™ 920 gives an impact strength of 44.6 kJ·m−<sup>2</sup> (93.1% increase compared to neat bioPET and an additional 10.1% compared to the uncompatibilized blend containing 30 wt % bioPA1010). It is evident the positive effect of the compatibilizer in improved toughness. With regard to the bio-based content, the blends with 30 wt % bioPA with compatibilizer show a bio-based content of approximately 50%. The GMA-based copolymers have been reported as good compatibilizers in different blends containing PAs or polyesters due to their reactivity with both polymers, such as PA6/PP [52,53], PA6/PVF [54], PET/PP [55], HDPE/PET [56].

The improved toughness is directly related to the morphology of the obtained materials. Figure 3 gathers FESEM images of uncompatibilized bioPET/bioPA blends. Figure 3a shows the fracture surface of neat PET with a typical rigid and brittle material, that is, very smooth surfaces resulting from microcrack appearances and their growth without plastic deformation. This brittle behaviour for neat PET has been reported by A.R. McLauchlin et al. [57] in PET/PLA blends. Figure 3b shows the morphology of the uncompatibilized binary blend containing 10 wt % bioPA. This morphology is remarkably different to neat bioPET. In particular, it is possible to observe a brittle fracture surface with a noticeable increase in roughness due to presence of small-sized bioPA immiscible droplets embedded into the bioPET matrix. As the bioPA wt % increases, the roughness is more evident, and the characteristic brittle fracture disappears. In Figure 3d, corresponding to the binary blend with 30 wt %

bioPA, it is possible to observe the characteristic droplet-like structure of an immiscible polymer blend in which bioPA appears in the form of spherical droplets with an average size of approximately 3.9 ± 1.1 μm. This size is higher than the average size observed for lower bioPA content. It is evident that by increasing the bioPA content, the average particle size increases due to the particle coalescence as reported by A.M. Torres-Huerta et al. [58] on PET/PLA and PET/chitosan blends. Y. Yan et al. [44] reported the high immiscibility of PET blends with PA56 (up to 30 wt %) and used dissipative particle dynamics (DPD) to assess the immiscibility of both polymers and how the PA56 domains increase with increasing PA56 content. The poor compatibility between these polymers can be observed at a 30 wt % bioPA in the blends as the morphology shows the typical spherical bioPA droplets embedded in the matrix (sea-island morphology) as well as some voids with the same average diameter consisting of some bioPA droplets on the holes which have been produced after being pulled out during the impact test. This pulling out occurs because of the poor polymer-polymer interactions.

**Figure 3.** Field emission scanning electron microscope (FESEM) images of the fractured surface from an impact test at 1000x corresponding to uncompatibilized bioPET/bioPA blends with different bioPA content, (**a**) 0 wt % (PET100), (**b**) 10 wt % (PET90), (**c**) 20 wt % (PET80) and (**d**) 30 wt % (PET70).

These results suggest that the poly(styrene*-ran-*glycidyl methacrylate) copolymer (PS-GMA) can positively contribute to a partial compatibilization effect due to the reaction of the glycidyl groups with both hydroxyl terminal groups in bioPET and amine groups in bioPA [54,55]. Y. Huang et al. [42] reported the reactivity of the glycidyl group of an epoxy resin (E-44) in a PET/PA6 binary blend. As the nature of both bioPET and bioPA is the same as petroleum-derived PET and other semicrystalline polyamides, it is possible to assume similar reactions as described in by Y. Huang et al. That study reported the greater tendency of the glycidyl group to react with polyamide due to the presence of many hydrogen bonding (together with carboxylic and amine end-chains) while the reaction of the glycidyl group with PET is restricted to hydroxyl and carboxyl end-chains. Moreover, Y. Huang et al. further reported evidences of these reactions by extracting the polyamide fraction by formic acid which was subjected to FTIR analysis. This analysis showed a shift of the N–H bending from 1560 to 1544

cm−<sup>1</sup> and shift of the C=O stretching from 1662 to 1642 cm<sup>−</sup>1, both changes indicating the reaction of PA6 with the glycidyl group in E-44 epoxy resin. In addition, the characteristic peaks of PET at 1730, 1104 and 730 cm−<sup>1</sup> were also detectable by FTIR thus giving consistency to the grafting process.

The indirect effects of these reactions can also be detectable by a remarkable change in surface morphology as it can be seen in Figure 4. The addition of PS-GMA Xibond™ 920 gives a noticeable decrease in the droplet size of bioPA-rich domains. There is not a great difference between the images corresponding to the compatibilized bioPET/bioPA blends containing 1, 3 or 5 phr PS-GMA. At this magnification (1000×), it can be realized that the droplet diameter has been reduced down to values under 1 μm. This situation can be clearly observed in Figure 5 which shows a comparative FESEM image of the uncompatibilized blend with 30 wt % bioPA and the same blend with 3 phr PS-GMA Xibond™ 920.

**Figure 4.** Field emission scanning electron microscope (FESEM) images of the fractured surface from an impact test at 1000x corresponding to compatibilized bioPET/bioPA blends with different loadings of Xibond™ 920 (in phr), (**a**) 1 phr, (**b**) 3 phr and (**c**) 5 phr.

As can be seen in Figure 5, the uncompatibilized blend (Figure 5a) shows a characteristic morphology, typical of poor polymer-polymer interactions. The particle droplet average diameter is 3.9 ± 1.1 μm as indicated previously. It is possible to observe this lack (or poor) interaction between both polymers. In fact, it is evident that some bioPA droplets have been removed (white rectangles) and there is a small gap between the bioPA spheres and the surrounding bioPET matrix (white arrows). Nevertheless, this gap is relatively low compared to other immiscible systems (it is in the nanoscale range), and this contributes to improved tensile properties and toughness as indicated previously. Furthermore, when observing Figure 5b is that the droplet size has been reduced in a remarkable way in the same blend with 3 phr PS-GMA compatibilizer. The new droplet size for the compatibilized blend is 0.62 ± 0.27 μm which is remarkably lower than the average size of bioPA-rich domains in the uncompatibilized blend. In addition, the surface morphology of the polymer matrix is different. As can be seen in Figure 5a for the uncompatibilized blend, the bioPET matrix is quite smooth while this surface is completely different in the compatibilized blend (Figure 5b) since it is remarkably rougher as reported by Y. Pietrasanta et al. [56] on HDPE/PET blends compatibilized with glycidyl polymers. Regarding the

gap between the embedded bioPA droplets, the morphology is also different since these embedded bioPA domains seem to be more embedded in the compatibilized blend. Similar findings have been reported for other PET-based immiscible blends such as those developed by C. Carrot et al. [59] (PET/PC) with a clear change in morphology in the presence of compatibilizers, O.M. Jazani et al. [60] (PET/PP), A.M. Torres-Huerta et al. [58] (PET/PLA) and (PET/chitosan), among others.

**Figure 5.** Detailed FESEM images corresponding to (**a**) uncompatibilized bioPET/bioPA blend with 30 wt % bioPA (PET70) and (**b**) compatibilized bioPET/bioPA blend with 30 wt % bioPA and 3 phr Xibond™ 920 (PET70Xibond3).

#### *3.2. Thermal and Thermo-Mechanical Properties of Binary BioPET*/*BioPA Blends*

The main thermal transitions of the developed materials are gathered in Figure 6. The neat bioPET shows in a clear way the main transitions. The step change in the 70–80 ◦C range corresponds to the glass transition phenomenon (*T*g) and it is 75.2 as shown in Table 4. Then, a peak located in the 120–140 ◦C range corresponds to the cold crystallization process which involves crystallization of the fraction that has not been able to crystallize because of the cooling rate. This process shows a characteristic peak temperature (*T*cc\_PET) of 133.2 ◦C. Finally, the melt process can be observed at higher temperatures of 225–260 ◦C with a peak temperature of 248.2 ◦C. The addition of bioPA up to 30 wt % on uncompatibilized blends does not provide any remarkable change in the *T*<sup>g</sup> with values of approximately 75–76 ◦C, very similar to neat bioPET. Although these *T*<sup>g</sup> values cannot be clearly seen in Figure 6a,b, the *T*<sup>g</sup> values were obtained from the zoomed DSC thermograms in the 65–85 ◦C leading to the values shown in Table 4. Regarding the cold crystallization process, bioPA plays a key role in this process. By the addition of 10 wt % bioPA, the peak temperature moved down to values of 121.9 ◦C, thus indicating bioPA enables crystallization of bioPET. Above 10 wt % bioPA, the cold crystallization process disappears and a slight decrease in the maximum crystallinity of bioPET (calculated with the obtained melt enthalpy values, Δ*H*m\_PET) can be detected as seen in Table 3. In fact, neat bioPET shows a maximum degree of crystallinity of 22.7% and it is slightly reduced to the values of 19.9% for the uncompatibilized blend containing 30 wt % bioPA. The melt peak temperature does not change in a remarkable way for all the developed materials and moves between the 247–248 ◦C narrow range. The effect of the PS-GMA compatibilizer is interesting. As can be seen in Table 4, a clear decrease in the crystallinity is detected from 19.9% (uncompatibilized blend with 30 wt % bioPA) to 17.3% in same composition with 3 phr Xibond™ 920. These results are in total agreement with those reported by Y. Huang et al. [42] who indicated a key role of the interface on crystallization as the interface is directly related to two relevant phenomena: Crystal nucleation and crystal growth. Y. Huang et al. report the use of an epoxy resin (E-44) as a compatibilizer in PET/PA6 blends and they conclude that although the epoxy resin can positively contribute to improve mechanical properties, a decrease in crystallinity is observed with increasing E-44 content due to the formation of less perfect crystals as a consequence of the increased interactions. Moreover, this study confirmed independent crystallization of PET and PA6 as suggested by the wide angle of the x-ray diffraction spectroscopy

(WAXD). In fact, Y. Huang et al. also report a different effect of epoxy compatibilization on hindering crystallization on both PET and PA6. The glycidyl group has more reactive points with PA6 due to the high number of hydrogen bonding in the structure while the reaction of the glycidyl group with PET is restricted to hydroxyl and carboxyl groups located at the end-chains. Y. Huang et al. reported a percentage decrease in the melt enthalpy of PET of approximately 25.9% while the decrease for PA6 is close to 40%. Due to the nature of both bioPET and bioPA, the same behaviour with the glycidyl compatibilizer is expected as can be seen in Table 4 with a clear large decrease in the melt enthalpy of bioPA compared to bioPET with increasing Xibond™ 920. On the other hand, Quiles-Carrillo et al. [61] reported a clear decrease in crystallinity by reactive extrusion of PLA with maleinized hemp seed oil (MHO). This decrease was attributed to the high reactivity of the maleic anhydride groups towards the hydroxyl groups in PLA which can give chain extension, branching and even, some crosslinking, all these phenomena having a negative effect on crystallization and formation of imperfect crystals. Quiles-Carrillo et al. [62] also reported a similar effect on PLA by using another reactive compatibilizer derived from soybean oil, namely, acrylated epoxidized soybean oil (AESO).

**Figure 6.** A plot comparative of the differential scanning calorimetry (DSC) thermograms corresponding to binary bioPET/bioPA blends.

The immiscibility of both polymers is also evident from DSC as two melt peaks are obtained with very slight changes in their corresponding peak temperature values. BioPA shows a melt peak located at 202–203 ◦C with whatever the composition may be. Nevertheless, the crystallinity is also affected by the presence of bioPET as the major component. As expected, the crystallinity of bioPA increases with increasing bioPA content from 9.8% (10 wt %) up to 16.9% (30 wt %) since the presence of higher bioPA loadings promote more intense and independent nucleation and crystal growth processes. Nevertheless, the effect of PS-GMA is the same as in the case of bioPET. The reaction between glycidyl groups in PS-GMA with both bioPET and bioPA polymer chains leads to the formation of imperfect crystals which is responsible for a decrease in the overall crystallinity as seen in Table 4, down to values of 10.2% for the compatibilized blend with 30 wt % bioPA and 5 phr Xibond™ 920. Another interesting finding is that the compatibilizer leads to a slight increase in the *T*<sup>g</sup> of bioPET up to values of 78 ◦C which is indicating that chain mobility is restricted. Despite this, the determination of *T*<sup>g</sup> by DSC is sometimes complex and inaccurate due to the problems related to the base line and the dilution effect in polymer blends. Similar findings have been reported by D. Garcia-Garcia et al. [63], using reactive extrusion of PHB and PCL with different dicumyl peroxide (DCP) loadings. The reaction of the free radicals generated by DCP can react with both PCL and PHB thus leading to partial compatibilization. These reactions reduce chain mobility as observed by the dynamic mechanical-thermal analysis (DMTA).

**Table 4.** A summary of the main thermal parameters of binary bioPET/bioPA blends obtained by differential scanning calorimetry (DSC).


Regarding thermal stability (degradation at high temperatures), Table 5 shows a summary of some thermal degradation parameters obtained by thermogravimetry (TGA). Two different characteristic temperatures are gathered in this table, the temperature required for a weight loss of 5% which is representative for the onset degradation (*T*5%) and the maximum degradation rate temperature (*T*max) which corresponds to the peak maximum of the first derivative TGA curve (DTG). As can be seen, the *T*5% for the neat bioPET is 382.6 ◦C and the addition of bioPA contributes to delay the onset degradation process as the *T*5% characteristic temperature is moved up to 397.4 ◦C for the uncompatibilized blend containing 30 wt % bioPA. This is because PA1010 has more thermal stability than PET. The effect of the PS-GMA compatibilizer is a slight increase in the onset degradation temperature up to values close to 404 ◦C with 3 phr Xibond™ 920. Regarding the maximum degradation rate, it is worthy to note a decreasing tendency with increasing bioPA loading on blends. This could be related to the fact that PA1010 is more thermally stable than PET but the degradation rate of PA1010 (change in weight loss with temperature) is higher than PET. For this reason, the *T*max shows a decreasing tendency. S. Jiang et al. [64] reported that the onset degradation temperature of PA1010 is located at 419.2 ◦C which is remarkably higher than PET, thus contributing to improved thermal stability.


**Table 5.** A summary of the thermal degradation of binary bioPET/bioPA blends obtained by thermogravimetry (TGA) analysis.

Regarding the effect of bioPA and the PS-GMA copolymer on mechanical-dynamical thermal properties, Figure 7 gathers some characteristic curves corresponding to the neat bioPET and the uncompatibilized and compatibilized (5 phr Xibond™ 920) blend with 30 wt % bioPA. Two main effects can be observed on the storage modulus, *G* . On the one hand, bioPET is stiffer than its blends with bioPA independently of the PS-GMA compatibilizer. T. Serhatkulu et al. [65] showed this flexibilization phenomenon on PET/PA6 blends. On the other hand, the presence of bioPA minimizes the cold crystallization process as observed in DSC. In fact, some cold crystallization occurs in bioPET/bioPA blends but DSC is not sensitive enough to detect it. However, these slight changes can be observed by DMTA as seen in Figure 7a. Another interesting phenomenon is the shift of the cold crystallization process towards lower temperatures which is in total accordance with the results obtained by DSC. The intensity of the cold crystallization can be observed in Figure 7b as the shoulder located to the right side. The *T*g values follow a similar tendency as that observed with DSC but DMTA seems to be more accurate to obtain these parameters. In particular, the *T*<sup>g</sup> for neat bioPET is 79.9 ◦C while the binary blend with 30 wt % bioPA shows a *T*<sup>g</sup> of 81.1 ◦C and the compatibilized blend (PET70Xibond5) offers a *T*<sup>g</sup> of 80.5 ◦C.

**Figure 7.** *Cont.*

**Figure 7.** The dynamic mechanical thermal behaviour (DMTA) of binary bioPET/bioPA blends in terms of increasing temperature (**a**) storage modulus, *G* and (**b**) dynamic damping factor, tan δ.

In addition to the dynamic mechanical-thermal analysis (DMTA), the dimensional stability has been studied by thermomechanical analysis (TMA). Figure 8 shows the TMA profiles of neat bioPET as well as the uncompatibilized and compatibilized (5 phr Xiboond™ 920) blend containing 30 wt % bioPA. From these TMA curves, it is possible to see the thermal behaviour of these materials. Below 60 ◦C, all three materials show a linear expansion (see Table 6 for values of the coefficient of linear thermal expansion, CLTE). Below this temperature, the slope is low compared to the slope above 120 ◦C. The glass transition temperature (*T*g) can be observed in the temperature range of 65–80 ◦C as the onset of a change in the slope. The slope is very high in the rubbery state from 80 ◦C up to 100 ◦C. Then, the slope is negative which is indicating increased dimensional stability. This is caused by the cold crystallization process. As seen previously by DSC, the cold crystallization peak is clearly detectable for neat bioPET and it is almost negligible for blends with high bioPA content. These results are in accordance with those shown in Figure 6 as the highest change in the dimensions can be seen for neat bioPET due to the cold crystallization process. Nevertheless, this change is very short for the other developed materials. Finally, above 120 ◦C, the linear tendency stabilizes, therefore indicating the cold crystallization has finished. Regarding the CLTE values (Table 6), it is worthy to note they follow the same tendency observed for ductile properties. The CLTE for neat bioPET is 152.4 μm m−<sup>1</sup> K<sup>−</sup>1, and it increases with increasing bioPA loading up to values of 347.3 μm m−<sup>1</sup> K−1. The effect of the compatibilizer is that expected since the reaction between the PS-GMA and bioPET and bioPA produces a restriction on chain mobility and this has a positive effect on dimensional stability. Notably, the CLTE value for the blend with 30 wt % bioiPA is compatibilized with 5 phr Xibond™ 920. All these results are in total agreement with the mechanical properties above-mentioned.

**Figure 8.** A comparative plot of thermomechanical behaviour of binary bioPET/bioPA blends in terms of increasing temperature.

**Table 6.** The calculated coefficient of linear thermal expansion (CLTE) of bioPET/bioPA blends obtained by thermomechanical analysis (TMA).


\* The CLTE has been calculated form the slope below 60 ◦C.

#### **4. Conclusions**

This work reports the viability of binary blends of partially bio-based poly(ethylene terephthalate) (bioPET) and fully bio-based poly(amide) 10,10 (bioPA1010) up to 30 wt % bioPA1010. Due to their immiscibility, a poly(styrene*-ran-*glycidyl methacrylate) (PS-GMA) copolymer (Xibond™ 920) is used to provide enhanced interaction. These blends can reach up to 50 wt % bio-based content without compromising other mechanical and thermal properties. The effectiveness of the PS-GMA has been corroborated with an increase in toughness, elongation at yield and tensile strength for a Xibond™ 920 loading of 3 phr. A FESEM study revealed a clear droplet-like structure with a bioPET matrix embedding bioPA-rich spherical (droplets) domains. The exceptional compatibilization effect of Xibond™ 920 in this binary blend is assessed by a remarkable decrease in the droplet diameter changing from almost 4 mm (uncompatibilized blend with 30 wt % bioPA) down to values lower than 1 mm (compatibilized blend with 30 wt % bioPA and 3 phr Xibond™ 920). Regarding the thermal properties, bioPA inhibits cold crystallization and a decrease in the degree of crystallinity of bioPET due to the formation of imperfect crystals. Xibond™ 920 also gives improved dimensional stability to blends thus leading to a new series of binary blends with balanced properties and a clear positive environmental impact since the bio-based content of these blends is close to 50 wt %.

**Author Contributions:** Conceptualization was devised by L.Q.-C. and R.B.; main experimental procedures were developed by M.J.; methodology, validation, and formal analysis was carried out by M.J., D.L., and S.M.-J.; investigation, resources, data curation, and writing-original draft preparation was performed by R.B. and N.M.

**Funding:** This research was funded by the Ministerio de Economía, Industria y Competitividad (MICINN) project number MAT2017-84909-C2-2-R. L. Quiles-Carrillo wants to thank GV for his FPI grant (ACIF/2016/182) and MECD for his FPU grant (FPU15/03812). D. Lascano wants to thank UPV for the grant received though the PAID-01-18 program. Microscopy services at UPV are acknowledged for their help in collecting and analyzing FESEM images. Authors thank Polyscope for kindly supplying Xibond™ 920 to carry out this study.

**Conflicts of Interest:** The authors declare no conflicts of interest.

#### **References**


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