Thermal properties and crystallisation kinetics were studied using a differential scanning calorimeter DSC 1/700 (Mettler Toledo, Greifensee, Switzerland), which was calibrated with indium and zinc standards. Experiments were carried on samples with different recycling cycles (1 to 7) under a constant nitrogen flow of 50 mL/min. Approximately 5 mg of sample specimens were prepared from the cross-section of the injection moulding parts on a rotating microtome Leica RM2255 (Leica Biosystem, Nußloch, Germany) were placed in 40 μL aluminium pans, sealed, and then placed in the DSC chamber. An empty pan was used as a reference. The specimens were heated from 0 to 200 ◦C, maintained at 200 ◦C for 3 min, cooled to 0 ◦C, and finally reheated to 200 ◦C at a heating rate of 10 ◦C/min. Melting peak temperature (*Tp,m*), cold crystallisation peak temperature (*Tp,cc*), premelt crystallisation peak temperature (*Tp,pc*), melting enthalpy (Δ*Hm*), cold crystallisation enthalpy (Δ*Hcc*), and premelt crystallisation enthalpy (Δ*Hpc*) obtained from the first and second heating scans and melt crystallisation peak temperature (*Tp,c*) as well melt crystallisation enthalpy (Δ*Hc*) determined from cooling scans were determined using the STARe software by Mettler Toledo (Greifensee, Switzerland). The samples were characterised at least in duplicate, and the averages were taken as representative values.

The crystallinity degree (*Xc*) of the samples as a function of the recycling cycle was determined according to Equation (1), where Δ*H*0 *m* is the theoretical melting enthalpy of the polymer assumed be 100% crystalline, Δ*H*0 *m* = 106 J/g for PLLA [34], and Δ *H*0 *m* = 146 J/g for PHBV [35].

$$X\_c(\%) = \frac{\Delta H\_m - \Delta H\_{pc} - \Delta H\_{cc}}{\Delta H\_m^0} \Delta 100 \tag{1}$$

In a study of non-isothermal crystallisation kinetics (at a constant cooling rate of 10 ◦C/min), the relative crystallinity (*XT*) was determined according to Equation (2), where Δ*HT* is the heat of melt crystallisation, and *T*, *T*0, *T*∞, and Δ*Hc* represent the temperature at any given moment, initial temperature, final temperature, and heat of melt crystallisation, respectively. The melt crystallisation time (*t*) is given by Equation (3), where *υ* is the constant cooling rate.

$$X\_T(\%) = \frac{\Delta H\_T}{\Delta H\_c} = \frac{\int\_{T\_0}^{T} \left(\frac{dH\_c}{dT}\right) dT}{\int\_{T\_0}^{T\_\infty} \left(\frac{dH\_c}{dT}\right) dT} \Delta 100 \tag{2}$$

$$t(min) = \frac{T\_0 - T}{v} \tag{3}$$

The kinetics of non-isothermal crystallisation was also evaluated according to the method introduced by Kratochvíl and Kelnar [33]. The principle of the method is schematically shown in Figure 1. It is a new simple method for evaluation of non-isothermal crystallisation kinetics at a constant cooling rate. The procedure based on mathematical treatment of the DSC cumulative crystallisation curves (i.e., the dependence of relative crystallinity on temperature) at their inflection point provides four basic parameters: temperature of the start of crystallisation (*Ts*), the temperature of maximum crystallisation rate (*Ti*), the numerical value of the maximum crystallisation rate (*si*) and final crystallinity after cooling from the melt (*Xc*). This approach is particularly convenient for the comparison of the non-isothermal crystallisation kinetics of samples in series with one reference, a virgin polymer. The method provides the temperature of crystallisation start (*Ts*) and maximum crystallisation rate (*Ti*) with standard deviation 0.3 and 0.4 ◦C, respectively. Maximum crystallisation rate (si) and final crystallinity (*Xc*) have coefficients of variation 5.8 and 1.5%, respectively [33]. The repeatability of *Ts*, *Ti*, and *si* improves with decreasing cooling rate. The method does not refer to any crystallisation models that are used in the evaluation of non-isothermal crystallisation (e.g., Ozawa, Nakamura, or Jeziorny extended Avrami kinetic model and others) [36,37] and eliminates the problem of exact setting of the starting time of crystallisation. As stated by Kratochvíl and Kelnar in their study [33], the parameters obtained sensitively describe the crystallisation process at the maximum rate, i.e., at relative crystallinities of about 35–45%. Thus, this method models real conditions encountered during processing polymeric materials well. The only prerequisite of successful application of the proposed method is a smooth cumulative crystallisation curve, i.e., well-developed crystallisation exotherm obtained by the DSC analysis.

### *2.3. Thermogravimetric Analysis (TGA)*

Thermal experiments were performed using TGA2 instrument (Mettler Toledo, Greifensee, Switzerland). The specimens from each recycling cycle were prepared from the cross-section (approximately 5 mg) and heated from 50 to 600 ◦C with 10 ◦C/min heating rate under nitrogen atmosphere (flow: 50 mL/min). The decomposition temperature was determined at 5% weight loss (*T*5) and maximum weight loss (*Tmax*) determined from the derivate thermogravimetry (DTG) curve, which corresponds to the inflection point of the TGA curve. Specimens were subjected to five repetitive tests, and the averages were taken as representative values.

**Figure 1.** Schematic drawing of the non-isothermal crystallisation kinetics method by Kratochvíl and Kelnar [33]: cumulative curve.

### *2.4. Rheological-Flow Properties*

Rheological-flow properties were determined by measuring the melt volume flow rate (MVR) at temperature 190 ◦C and under a load of 2.16 kg using a melt flow tester Ceast (Ceast, Torino, Italy), according to ISO 1133. Before measurement, the material was dried in a vacuum oven Binder VD53 (Binder, Tuttlingen, Germany) to a residual moisture content of less than 0.01%. Residual moisture was controlled on an HX204 halogen analyser (Mettler Toledo, Greifensee, Switzerland) at 130 ◦C.

### *2.5. Mechanical Properties*

Tensile testing, flexural testing, and impact testing were carried on multipurpose test specimens of type A out at standard conditions 23/50, according to ISO 291. Specimens of type A are tensile test specimens, from which specimens with dimensions (80 × 10 × 4) mm have been obtained by machining to determine flexural and impact properties. Tensile modulus, tensile strength, and tensile strain at break values were determined according to ISO 527 standard using LabTest universal electromechanical testing instruments (Labortech, Opava, Czech Republic) with extensometer MFL-300B (Mess- & Feinwerktechnik, Velbert, Germany) with accuracy class ISO 9513, at a crosshead speed of 1 mm/min for determined tensile modulus and 5 mm/min for determined tensile strength and elongation at break, a 10 kN load cell, and gauche length of 50 mm. Flexural modulus and flexural strength values were determined according to ISO 178 (three-point loading test) using a Tinius Olsen H10KT testing machine (Tinius Olsen, Salfords, UK) at a crosshead speed of 2 mm/min. Charpy impact strength values were carried out following ISO 179-1/1eU on the equipment Resil 5.5 (Ceast, Torino, Italy), with a force of 5 J for PLLA and 2 J for PHBV. Mechanical analyses were repeated at least ten-fold, and the averages were taken as representative values.

### *2.6. Heat Resistance Testing*

Heat resistance testing of the recycles biobased polymers was performed with a ZwickRoell's HDT/Vicat 6-300 Allround instrument (ZwickRoell, Ulm, Germany) by determining the Vicat softening temperature (VST) according to ISO 306. The dimensions of the specimens were (10 × 10 × 4) mm, VST were measured under a load of 50 N, and the temperature of the oil bath was raised by 50 ◦C/h. The load was applied to the specimens

after 5 minutes of immersion in an oil bath at an initial temperature of 25 ◦C. Specimens were subjected to five repetitive tests.

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

### *3.1. Thermal Properties and Structure of Recycled Biobased Polymers*

Unlike PHBV, PLLA is a biobased polymer with a characteristic slow melt crystallisation rate. During its subsequent heating, cold crystallisation, and sometimes premelt crystallisation, can be observed. These exothermal reactions have direct impact on the heat quantum necessary for crystal melting. As a result, it is necessary to subtract the values of cold crystallisation (Δ*Hcc*) and premelt crystallisation (Δ*Hpc*) from the final melting enthalpy (Δ*Hm*) when calculating the degree of crystallinity of PLLA (*Xc*) by DSC (Equation (1)). The resulting values of these endothermic and exothermic phase transformations, including the degree of crystallinity, are listed in Table 1 for PLLA and Table 2 for PHBV. The degree of crystallinity is evaluated from the first and second heating phases. The first heating phase reflects the supramolecular structure of the moulded part, which results from the cooling conditions during the moulding process, and the second heating phase reflects the crystallisation of the biobased polymer from the melt at a slow (laboratory) cooling rate of 10 ◦C/min after removing its thermal history. As a result of repeated mechanical recycling, a decrease in the values of premelt crystallisation enthalpy (Δ*Hpc*) and cold crystallisation enthalpy (Δ*Hcc*) was detected in PLLA. The slow cooling rate in the calorimeter resulted in a complete diminishing of the cold crystallisation during the first recycling cycle (Figure 2a). This effect is initiated by the significant increase in Δ*Hc* values during the crystallisation of PLLA from the melt (Figure 2b). Shojaeiarani et al. [20], Pillin et al. [27], and Zaverl et al. [30] found that, for PLA and PHBV with a thermomechanical history of a successive recycling process, the weight average molecular weight, the number average of molecular weight, and the molecular weight distribution values decreased significantly.

**Table 1.** Thermal properties and crystallinity degree of PLLA as a function on the number of recycling cycles.

**Table 2.** Thermal properties and crystallinity degree of PHBV as a function on the number of recycling cycles.

**Figure 2.** PLLA cold and melt crystallisation as a function of recycling cycle: (**a**) cold crystallisation (second heating); (**b**) melt crystallisation.

Chain cleavage induced by thermomechanical degradation during repeated milling and PLLA injection moulding resulted in shorter polymer chains, which acted as nucleation centres. This premise potentially explains the increase in the melt crystallisation enthalpy (Δ*Hc*) and final degree of crystallinity (*Xc*). These results are partially inconsistent with the results reported in the study by Shojaeiarani et al. [20] and Badia et al. [25]. The reason is probably the high optical purity of PLLA, as both studies analysed poly (lactic acid), which contained a minimum of 4% mol of D-lactic enantiomer. The degree of crystallinity increased, on average, by three times in the seventh recycling cycle when PLLA melt was cooled at a rate of 10 ◦C/min. On the other hand, the increase in the degree of crystallinity observed in moulded parts reached only 48%. This difference is probably induced by various thermodynamic cooling conditions during solidification of the melt for moulded part and sample studied under laboratory conditions and is related to the limited movement of the macromolecules.

Chain cleavage is evident from the rheological behaviour of the material (Figure 3). As the number of recycling cycles increases, the values of melt flow volume index (*MVR*) increase. This phenomenon is induced by the decrease in the molecular weight of the analysed biobased polymer. In the seventh PLLA recycling cycle, the melt flow volume index increased six-fold compared to the virgin material (pellets) and 3.6-fold compared to the moulded part made of the virgin material. For PHBV, an increase in *MVR* value was recorded, almost three-fold that of the virgin material in the form of pellets and 2.2-fold that of the moulded part made of the virgin material. Repeated mechanical recycling of biobased polymers causes a significant reduction in viscosity, especially for PLLA, which is an apparent symptom of a lower molecular weight induced by mechanical and thermal degradation. The negative effect of shear stress during the injection moulding process on the degradation of PLLA and PHBV is already evident from the change in the *MVR* values measured on the supplied commercial pellets and the crushed material milled from the moulded parts produced from the same pellets. This finding is consistent with the results of a study published by Pantani et al. [38], which deals with the degradation of PLA caused by different shear rates during its processing (extrusion and injection moulding). Between the pellets and the crushed material milled from the moulded parts produced from the same pellets, an increase in the *MVR* was observed by approximately 68% for PLLA and by 29% for PHBV. Abe [26] stated that a significant rise in melt volume flow rate for PHBV polymer could be attributed to the unzipping reaction in the biobased polymer structure through cis-elimination mechanism (Mclafferty arrangement) initiated

just above its melting temperature. Similarly, the degradation of PLA is induced by the generation of acidic molecules acting as a catalyst to accelerate the degradation. According to Dai et al. [28], the considerable reduction in the molecular weight of PLA was instead associated with the thermal degradation caused by polymer chain scissions into linear and cyclic oligomers.

**Figure 3.** PLLA and PHBV melt volume flow rate as a function of recycling cycle.

From the results of PLLA DSC analysis, is evident that the glass transition temperature (*Tg* ~60 ◦C) did not decrease with increasing level of recycling (repeated thermomechanical loading). This finding is also confirmed by Badia et al. [25] and Zenkiewicz et al. [ ˙ 39]. On the other hand, Pillin et al. [27] found that, during the second recycling cycle, *Tg* of the analysed PLA decreased. According to the Foxe–Flory relationship [40], this indicates that the decrease in molecular weight induced by reprocessing of PLLA was not significant enough to reduce the *Tg* dramatically.

The decrease in molecular weight affects melt crystallisation and, in addition to increasing the melt crystallisation enthalpy (Δ*Hc*), the melt crystallisation temperature (*Tp,c*) also increases, which means that the PLLA melt crystallises faster after repeated mechanical recycling (Figure 2b). A slight increase in Δ*Hc* and *Tp,c* was also observed for highly crystalline PHBV. In the case of the Δ*Hc* value, an increase of only 2.5% was recorded for the seventh level of PHBV recycling, while for PLLA, this increase was almost tenfold (see Tables 1 and 2). The biggest leap in crystallisation enthalpy was observed in the first recycling cycle. In PHBV, chain cleavage did not affect the phase transformation process in the frame of melt crystallisation, as in the case of PLLA. The resulting degree of crystallinity (*Xc*) thus does not change for repeatedly processed PHBV (Table 2). A similar finding was made by Shojaeiarani [20]. On the contrary, this result contradicts the findings of Zaverl et al. [30], who recorded a decrease in the value of *Xc* induced by repeated processing of PHBV (decrease by 20% in the fifth recycling cycle). The proposal can be introduced that this divergence can be caused by the different valerate content in the PHBV copolymer. Shojaeiarani et al. [20] found that chain cleavage occurs mainly in long chain biobased polymers. The higher rate of degradation of PLLA induced by thermomechanical loading during recycling is therefore probably caused by easier cleavage of the chains in the amorphous part of the supramolecular structure, which predominates in PLLA after injection into cold mould (20 ◦C).

From the results of PLLA DSC analysis (see Figure 4), it is evident that the endothermic phase transformation of melting (Peak III) is preceded by a sharper exotherm (Peak II) and a small and narrow endotherm (Peak I) in addition to cold crystallisation. The initial endotherm (Peak I) may be due to the melting of PLLA crystals with low thermal stability, i.e., paracrystalline or microcrystalline structures, formed at *Tp,c* and/or in the temperature range of cold crystallisation. Successive structural recrystallisation and reorganisation

(premelt crystallisation) may result in the appearance of an exotherm (Peak II) and a final endotherm (Peak III) [41,42]. A second alternative explanation for the presence of the endothermic overshoot in PLLA relies on the physical ageing of the rigid amorphous fraction, that is, the glassy fraction at the interface between the mobile amorphous fraction (*Tg*) and the crystalline fraction [43]. Di Lorenzo [44] states that the position and dimension of the various peaks strongly depend on crystallisation temperature.

**Figure 4.** DSC curves (second heating) of PLLA as a function of recycling cycle.

Furthermore, the results showed that repeated mechanical recycling induced a gradual increase in melt crystallisation temperature (*Tp,c*) (see Figure 2b). As the temperature *Tp,c* increased, Peak I and Peak II moved to higher temperatures, Peak I became more pronounced, and Peak II was less intense (see Figure 4). In the seventh recycling cycle, Peak I coincided with Peak III. When crystallisation is performed at higher temperatures, due to repeated mechanical recycling, a single melting peak is always present, and its position is strongly affected by *Tp,c*.

The first DSC temperature cycle, which reflects the macromolecular structure of the moulded parts, shows that the endothermic melting transformation, in addition to cold crystallisation, is preceded by only exotherm (peak II) associated with structural recrystallisation caused by fast cooling rate of the melt after injection of PLLA into the cold mould (Figure 5).

Changes in the macromolecular structure of PLLA and PHBV induced by their gradual recycling were also evaluated by thermogravimetric analysis (TGA). In Table 3, the temperatures corresponding to 5% weight loss and maximum rate of weight loss for each biobased polymers recycling cycle are represented by *T5* and *Tmax*. Degradation of the moulded part made of virgin material (PLLA) began at 343 ◦C, and maximum weight loss occurred at 370 ◦C. In the case of PLLA, after seven recycling cycles, the material started to degrade by 71 ◦C earlier, i.e., at 272 ◦C, and the maximum weight loss occurred at 346 ◦C, i.e., 24 ◦C earlier than in the case neat injection moulded samples. PHBV also accelerated the onset of thermal degradation with increased recycling. However, the temperature changes were significantly lower than for PLLA. In the seventh recycling cycle, the temperature of T5 decreased by 6 ◦C and the *Tmax* by 5 ◦C, compared to the moulded part made of virgin material (PLLA). A similar decrease in degradation temperature, but significantly lower, was observed for PLA by Shojaeiarani et al. [20] and Zenkiewicz et al. [ ˙ 39] and for PHBV by Shojaeiarani et al. [20] and Zaverl et al. [30]. From the results, it can be stated that the higher molecular weight of PHBV increases its thermal stability during repeated processing. Carrasco et al. [45] and Crompton [46] also report that there is a linear relationship between

thermal stability and the average molecular weight, in which the higher molecular weight can result in more thermally stable polymer materials.

**Figure 5.** DSC curves (first heating) of PLLA as a function of recycling cycle.

**Table 3.** Thermogravimetric analysis of PLLA and PHBV as a function on the number of recycling cycles.

### *3.2. Effect of Recycling on Non-Isothermal Crystallisation Kinetics*

The kinetics of non-isothermal crystallisation of PLLA and PHBV affected by their gradual recycling (in the range of seven recycling cycles) was monitored by DSC at a constant cooling rate of 10 ◦C/min. The trends in the behaviour of relative crystallinity (XT) over time was recorded (Figure 6). These graphical dependences were constructed from DSC measurements based on Equations (2) and (3). Table 4 shows the values of the half-time crystallisation (*t*0.5), defined as the time when crystallisation reached 50%. From the results, is obvious that the values of *t*0.5 for PLLA decrease gradually due to recycling; however, for PHBV, they remain unchanged. A decrease in *t*0.5 by 11% was recorded in the first PLLA recycling cycle and by 40% in the seventh recycling cycle. The decrease in *t*0.5 values caused by the gradual recycling of PLLA shows a trend that is typical in the case of faster cooling from the melt, as reported, for example, by Tarani et al. [47]. This effect is induced by the increase in the melt crystallisation temperature (*Tp,c*) [48], which increases with gradual recycling for PLLA, while for PHBV, the change is insignificant (see Tables 1 and 2). The reduction in *t*0.5 caused by repeated mechanical recycling confirms the cleavage of the chains, which act as nucleation centres during the non-isothermal crystallisation of PLLA.

**Figure 6.** Relative crystallinity versus time for non-isothermal melt crystallisation (**a**) PLLA and (**b**) PHBV at cooling rate 10 ◦C/min.

**Table 4.** Crystallisation half-time (*t*0.5), crystallisation starting temperature (*TS*), inflection point temperature (*Ti*), relative crystallinity at the inflection point (*XTi*), and maximum crystallisation rate (*si*) during melt crystallisation of PLLA and PHBV as a function on the number of recycling cycles (cooling rate: 10 ◦C/min).

Table 4 lists several valuable parameters describing the non-isothermal kinetics of PLLA and PHBV crystallisation during their gradual recycling by the method introduced by Kratochvíl and Kelnar [33]. The starting point is the non-isothermal cumulative crystallisation curve (Figure 7), i.e., the dependence of relative crystallinity (*XT*) on temperature (*T*). The method is based on the fact that the slope of the tangent to the cumulative crystallisation curve at any point is directly proportional to the crystallisation rate at the corresponding temperature. The inflection point of the cumulative crystallisation curve (*si*) is the point of maximum crystallisation rate. It is specified by temperature *Ti* and relative crystallinity *XTi* (Figure 1). The tangent slope of the cumulative crystallisation curve at the inflection point, in 1/◦C, is the numerical value of the maximum crystallisation rate. The cross-section of the tangent with the T-axis determines the starting temperature of crystallisation (*Ts*) [33]. From the above-mentioned parameters, the fact arises that the gradual recycling of PLLA increases the temperature of the melt crystallisation process onset and crystallisation rate. For the moulded part made of virgin PLLA, the maximum crystallisation rate (*si*) of 5.6 1/◦C was measured at 100.8 ◦C (*Ti*), i.e., about 8 ◦C after the onset of crystallisation. At the same time, the PLLA after the seventh recycling cycle showed a maximum crystallisation rate 2.7 times higher (15.3 1 /◦C), which already occurs at a temperature of 110.9 ◦C, i.e., approximately 3 ◦C after the onset of crystallisation. In the case of PHBV, the temperatures of *Ts* and *Ti* increase by gradual recycling, but the maximum crystallisation rate (*si*) appears to be independent of the number of recycling cycles. In addition, the shift in temperatures *Ts* and *Ti* is significantly lower than in PLLA. The parameters obtained describing the crystallisation process at its maximum rate correspond to relative crystallinity of approximately 44–54% for PLLA and 37–43% for PHBV, which corresponds to the results of Kratochvíl and Kelnar [33]. Paukszta and Borysiak also observed an increase in the temperature and rate of melt crystallisation with non-isothermal cooling during repeated processing of the PP/PA6 polymer blend [49].

**Figure 7.** Relative crystallinity versus temperature for non-isothermal crystallisation (**a**) PLLA and (**b**) PHBV—cumulative crystallisation curves.

Cold crystallisation occurred in the first phase of PLLA heating at a constant rate of 10 ◦C/min. The first heating reflects the thermal history of the PLLA created during the injection moulding process. The thermal characteristics are given in Table 5 and are based on DSC curves, on the graphical dependence of relative crystallinity on time, and the graphical dependences of relative crystallinity on temperature during cold crystallisation, respectively (Figure 8).

**Table 5.** Cold crystallisation temperature (*Tp,cc*), cold crystallisation starting temperature (*Ts,cc*), inflection point temperature during cold crystallisation (*Ti,cc*), relative cold crystallinity at the inflection point (*XTi,cc*), and maximum cold crystallisation rate (*si,cc*) of PLLA as a function on the number of recycling cycles.

**Figure 8.** Relative crystallinity versus (**a**) time and (**b**) temperature for non-isothermal cold crystallisation PLLA.

Gradual recycling showed that the cold crystallisation peak temperature (*Tp,cc*) and the temperature of cold crystallisation start (*Ts,cc*), which was determined by non-isothermal heating according to the method introduced by Kratochvíl and Kelnar [33], shifted to lower temperatures. The difference between *Tp,cc* and *Ts,cc* between the recycling cycle 0 and 7 is approximately 5 ◦C. The evaluation of the non-isothermal kinetics of cold crystallisation further shows that repeated recycling does not affect the maximum rate of cold crystallisation, in contrast to melt crystallisation. Similar to the beginning of cold crystallisation (*Ts,cc*), the temperature of the maximum rate of cold crystallisation (*Ti,cc*) shifts to lower values. Additionally, the difference in *Ti,cc* values between recycling cycle 0 and 7 is approximately 5 ◦C. The results thus showed that the heating of gradually recycled PLLA, in which the volume of crystalline structure formed during solidification from the melt did not meet the potential of this material (caused by rapid cooling of the material within injection mould), activates conformational processes of macromolecules at lower temperatures. These processes are associated with the cold crystallisation of PLLA. However, the rate of cold crystallisation and the enthalpy of cold crystallisation are consistent for an entire range of reprocessed PLLA that were analysed (see Tables 1 and 5). The nucleation centres created by chain cleavage during repeated recycling of PLLA were thus fully utilised in melt crystallisation. The obtained parameters describing the PLLA cold crystallisation process at its maximum rate correspond to a relative crystallinity of approximately 47–51%.

### *3.3. Effect of Recycling on Mechanical Properties*

The mechanical properties of virgin and recycled PLLA and PHBV biobased polymers were studied through tensile, flexural, and impact tests, and the results are summarised in Figures 9–12. PLA exhibited significantly lower tensile and flexural strength with the increasing number of recycling cycles. In the seventh recycling cycle, an average decrease in tensile strength by 22% was recorded (Figure 9a) and flexural strength even by 46% (Figure 11a). The highest range of changes was recorded in the first recycling cycle, when a decrease in tensile strength reached almost 9%, and flexural strength dropped by 14%. The greatest decrease in tensile strength, comparing the first recycling cycle and the virgin material, was also recorded in PLA by Zenkiewicz et al. [ ˙ 39]. Changes in tensile strength and flexural strength of PHBV are statistically insignificant. In the seventh recycling cycle, an average decrease in flexural strength by almost 3% was recorded for PHBV, and quite contrary, a slight increase by almost 6% was recorded for tensile strength. A very similar

trend, but not with such significant differences, was also observed by Shojaeiarani et al. [20]. A decrease in mechanical strength can probably be attributed to the lower molecular weight in the recycled polymers due to the deterioration of the polymer chains during successive injection moulding cycles. In general, the mechanical properties of biobased polymers are strongly dependent on molecular weight. Therefore, any changes in the polymer chain as a result of degradation induced by high temperature and shear within the injection moulding machine directly impact the mechanical properties. Tensile and flexural strength were significantly lower for reprocessed PLA as compared to their corresponding counterpart made of virgin material (as opposed to reprocessed PHBV), which is probably caused by the trend of decreasing molecular weight in recycled polymers indirectly observed by the change in viscosity and melt volume flow index, respectively (Figure 3).

**Figure 9.** PLLA and PHBV tensile properties as a function of recycled cycle: (**a**) tensile strength; (**b**) tensile modulus.

**Figure 10.** PLLA and PHBV tensile strain at break as a function of recycled cycle.

**Figure 11.** PLLA and PHBV flexural properties as a function of recycled cycle: (**a**) flexural strength; (**b**) flexural modulus.

**Figure 12.** PLLA and PHBV Charpy impact strength as a function of recycled cycle.

The tensile and flexural modulus does not show significant changes in recycled PLLA (Figures 9b and 11b). Considering the scattering of the measured values, these changes appear to be insignificant. These results were also confirmed by the study focused on the tensile modulus published by Pillin et al. [27] for recycled PLA containing 92% of L-lactide and 8% of D-lactide. Conversely, for recycled PHBV, an increase in the mean value of the modulus of elasticity in tension and bending was recorded (by about 6% and 10% in the seventh recycling cycle). The recycling process includes polymer crushing, melting, and its repeated shear stress, which leads to polymer chain degradation. However, thermal degradation in the melting process did not seem to contribute to a significant failure of the intermolecular bonds between the polymer chains, because the stiffness of the polymers did not show a decreasing trend with the increased number of recycling cycles.

The tensile test results showed that PLLA and PHBV exhibited a typical brittle fracture with strain values of less than 10% (around 2% to 4%). This effect occurs as a result of testing PLLA below the glass transition temperature. For PHBV, the brittle fracture is caused by the relatively high degree of crystallinity of the moulded part (about 60%, see Table 2). The brittleness of poly(hydroxybutyrate) (PHB) with the crystallinity degree puts into context a study by El-Taweel et al. [50]. El-Taweel et al. found that, above 40% crystallinity, PHB exhibited brittle fracture. Because the degree of crystallinity of recycled PHBV did not decrease, the tensile strain at break is almost constant in the frame of this study (Figure 10). Conversely, for recycled PLLA, the brittleness increases because of increasing the degree of

crystallinity (Table 1), and another aspect is the material degradation, and that is why the tensile strain at break decreases. Crystalline fractions might favour the crack propagation in the amorphous domain. In the seventh recycling cycle, a decrease in the tensile strain at the break by almost 59% was recorded for PLLA. A significant decrease in tensile strain at break for PLA was also noted by Pillin et al. [27]. In accordance with the findings of Zenkiewicz et al. [ ˙ 39], the results again showed the most significant changes between the moulded part made of material after first recycling cycle and the virgin material, where the tensile strain at break decreased by 28%.

The impact strength of the polymers is shown in Figure 12. It can be seen that recycled PLLA exhibit lower impact strength as compared with their corresponding virgin polymer. The impact strength of PLLA decreased by 9% in the fifth recycling cycle and by 13% in the seventh recycling cycle. The same 9% decrease in impact strength at the fifth recycling cycle for PLA was also observed by Shojaeiarani et al. [20]. However, recycled PLLA shows a higher variance in the measured values. For PHBV, the results showed that the effect of the recycling process on the impact strength was not significant, and the impact strength remained nearly constant after the recycling process. On the contrary, Shojaeiarani et al. [20] observed a decrease by 30% in impact strength in the fifth recycling cycle of PHBV. Based on these results, the assumption can be deduced that in the case of PHBV copolymer, the change in properties during its repeated recycling may be affected by the amount of valerate. The decrease in the impact strength of PLLA can be explained by a significant change in the structure of the recycled PLLA (by the lower molecular weight as indicated by the results of the *MVR* determination) due to the degradation phenomenon caused by heating and shear stress through the recycling process.

The results of the mechanical properties of recycled PHBV show the easy recyclability, which means that the product with required properties can be made of this material.

### *3.4. Effect of Recycling on Heat Resistance Testing*

The heat resistance of virgin and recycled PLLA and PHBV biobased polymers was studied by determining the Vicat softening temperature (*VST*), and the results are summarised in Figure 13. For PLLA, the *VST* values remain constant during reprocessing. Considering that PLLA is a polymer with a lower degree of crystallinity, the *VST* corresponds to the glass transition temperature, which was also constant in the frame of performed analyses focused on the study of recycled PLLA. There was a slight decrease in *VST* for PHBV. In the seventh recycling cycle, the *VST* decreased by only 5 ◦C. This effect can be explained by the stability of the melting point and crystallinity values of the recycled PHBV (see Table 2).

**Figure 13.** PLLA and PHBV Vicat softening temperature as a function of recycled cycle.

### **4. Conclusions**

In this work, the effect of primary mechanical recycling on the physical properties and non-isothermal crystallisation kinetics of reprocessed poly(L-lactide) (PLLA) and poly(3 hydroxybutyrate-*co*-3-hydroxyvalerate) (PHBV) biobased polymers was presented and discussed. Biobased polymers were reprocessed by milling and injection moulding in seven recycling cycles. The physical properties were investigated by differential scanning calorimetry (DSC), thermogravimetry (TGA), determination of melt volume flow rate (*MVR*), Vicat softening temperature (*VST*), and mechanical tests under static tensile and flexural loading and Charpy impact strength. The non-isothermal crystallisation kinetics during melt crystallisation and cold crystallisation were evaluated by the half-time crystallisation, and a relatively new procedure based on mathematical treatment of DSC cumulative crystallisation curves at their inflection point provides three kinetic parameters: temperature of the crystallisation start, temperature of maximum crystallisation rate and the numerical value of the maximum crystallisation rate, and final crystallinity after cooling.

This study confirmed the premise that recycled PLLA and PHBV undergo structural changes, which result in molecular modification of structure initiated by the reprocessing and repeated heating and mechanical stress of the material. This conclusion is based on the increase in the melt flow properties of the biobased polymers and the decrease in the initial degradation temperature determined by the TGA method. In the seventh recycling cycle, the flow properties increased six-fold for PLLA compared to the level reached by the virgin material in the form of pellets (before injection moulding process) and almost three-fold for PHBV. In the case of PLLA, the onset of thermal degradation of recycled material was reduced by up to 80 ◦C comparing the level reached by the material sampled from the moulded part moulded from the virgin material (after injection moulding), while in PHBV, there was a decrease by, at most, 10 ◦C. The structural changes were the most pronounced in PLLA among the studied materials. The results of DSC analysis are revealed as follows. While with reprocessing PHBV, there was only a negligible increase in melt crystallisation temperature and maximum melt crystallisation rate, and the overall degree of crystallinity remained constant; with PLLA, there was a significant change in crystallisation kinetics. As the number of recycling cycles increased, so did the start melt crystallisation temperature, the maximum crystallisation rate (by up to 2.7 times), and the melt crystallisation enthalpy (almost 10 times). Faster crystallisation from the melt reduced the risk of the polymer recrystallisation that is induced by additional heating of the moulded part. The results also provide the findings that the cold crystallisation of PLLA, which occurred after the first recycling cycle, was prevented by slow laboratory cooling. With the gradual recycling, the endothermic transformation associated with the melting of crystals with low thermal stability or with the physical ageing of the rigid amorphous fraction was also reduced. These changes are most likely caused by the polymer chain cleavage, and these chain fragments act furthermore as nucleation centres in the melt crystallisation region but do not affect the cold crystallisation, where the results showed that the maximum rate of cold crystallisation does not change with biobased polymer reprocessing. Still, the onset temperature of cold crystallisation was lowered. Significant structural changes in PLLA were also reflected in its mechanical properties. In the case of reprocessed PLLA (after seven recycling cycles), the tensile strength was reduced by ~22%, the flexural strength dropped even by 46%, the impact strength of Charpy was reduced by ~13%, and the total elongation of this naturally brittle polymer (tensile strain at break) was decreased by 59%. On the contrary, changes in the bending and tensile modulus appear to be insignificant. When measuring the glass transition temperature by DSC, the findings showed that values did not change with repeated recycling of PLLA, resulting in the fact that the Vicat softening temperature was also constant throughout the analysed recycling range of PLLA. In contrast, a slight decrease in VST by 5 ◦C was observed for reprocessed PHBV. The performance of this biopolymer was very consistent throughout the entire analysed range of recycling cycles with minimal changes in mechanical properties. The

results suggest that PHBV processing waste, unlike PLLA, is suitable for recycling as a reusable additive to the virgin biopolymer, at least for the applications with no risk of exposure to significant degradation processes.

**Author Contributions:** Conceptualisation, L.B.; methodology, L.B., J.N. and M.B.; validation, L.B., J.N. and P.L.; formal analysis, L.B., J.H. and J.N.; investigation, L.B., P.B., M.B., and P.L.; resources, P.B. and L.B.; data curation, L.B.; writing—original draft preparation, L.B.; writing—review and editing, L.B. and M.B.; visualisation, J.N.; supervision, P.L.; project administration, L.B. and P.L.; funding acquisition, L.B. and P.L. All authors have read and agreed to the published version of the manuscript.

**Funding:** This work was supported by the Student Grant Competition of the Technical University of Liberec under the project No. SGS-2019-5015 "Research and development for innovation of materials and production technologies with application potential in mechanical engineering". Further funding was supported by the Ministry of Education, Youth and Sports of the Czech Republic and the European Union—European Structural and Investment Funds in the frames of Operational Program Research, Development and Education, project Hybrid Materials for Hierarchical Structures (HyHi, Reg. No. CZ.02.1.01/0.0/0.0/16\_019/0000843).

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** Not applicable.

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

### **References**

### *Article* **A Multifaceted Approach for Cryogenic Waste Tire Recycling**

**Darkhan Yerezhep, Aliya Tychengulova, Dmitriy Sokolov and Abdurakhman Aldiyarov \***

Faculty of Physics and Technology, Al Farabi Kazakh National University, 71 Al-Farabi Ave., Almaty 050040, Kazakhstan<EMAIL_ADDRESS>(D.Y<EMAIL_ADDRESS>(A.T<EMAIL_ADDRESS>(D.S.)

**\*** Correspondence<EMAIL_ADDRESS>
**Abstract:** One of the important aspects for degradation of the life quality is the ever increasing volume and range of industrial wastes. Polymer wastes, such as automotive tire rubber, are a source of long-term environmental pollution. This paper presents an approach to simplifying the rubber waste recycling process using cryogenic temperatures. The temperature of cryogenic treatment is ranged from 77 K to 280 K. Liquid nitrogen was used as a cryoagent for laboratory tests. Experimental and numerical studies have been carried out to determine the optimal conditions for the recycling process. Numerical studies were performed using the COMSOL Multiphysics cross-platform software. The optimal force of mechanical shock for the destruction of a tire which turned into a glassy state after cryoexposure was determined experimentally. The chemical and physical properties of the final product (crumb rubber) have been studied by scanning electron microscopy and energy dispersive Xray spectroscopy. The analysis shows that the morphology and elemental composition of the samples remain practically unchanged, demonstrating environmental friendliness of the proposed process.

**Keywords:** polymer; rubber; recycling; cryoagent; liquid nitrogen; waste tire; thermal conductivity

### **1. Introduction**

World production of natural and synthetic rubber is growing by about 2% annually. Thus, in 2020, the total rubber production amounted to more than 25 million tons [1,2]. More than 60% of them are spent on the production of automobile tires [3]. Consequently, this causes an increase in polymer waste in the form of used tires [4]. Waste tire rubber, which needs to be disposed or recycled, is a source of long-term environmental pollution, since the degradation period of rubber in the ground is more than a hundred years. Furthermore, rubber is flammable, hazardous and is subject to long-term biodegradation, with an extremely harmful effect on the environment, in addition to poisonous substances such as benzene, xylene, styrene, toluene, etc. [5–7], which are released when burning.

Today, about 10 million tons of tires are annually used in the world by recycling or by other methods [2]. However, the accumulated stocks of waste tires requiring recycling or disposal are about 100 million tons. These data indicate a large volume of potentially valuable raw materials that require processing for further use [8,9]. According to preliminary evaluation, only about 10% of waste tires can be recovered, but this only delays the moment when they need to be disposed or recycled [10,11].

Despite its enormous scale, the problem of waste tire rubber recycling is solvable, and there are many ways to do it [12–17]. In principle, all known methods of waste tire rubber recycling can be divided into two main groups: physical and chemical. Chemical processing methods lead to irreversible chemical changes not only in rubber, but also in the constituent substances (vulcanizers, softeners, plasticizers, etc.). These methods are carried out at high temperatures resulting in material destruction [18–22]. Despite the fact that chemical methods of waste rubber recycling make it possible to obtain valuable products and heat, such use is not efficient enough, since it does not allow preserving the original polymer materials.

**Citation:** Yerezhep, D.; Tychengulova, A.; Sokolov, D.; Aldiyarov, A. A Multifaceted Approach for Cryogenic Waste Tire Recycling. *Polymers* **2021**, *13*, 2494. https://doi.org/10.3390/ polym13152494

Academic Editors: Sheila Devasahayam, Raman Singh and Vladimir Strezov

Received: 19 June 2021 Accepted: 26 July 2021 Published: 28 July 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 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 (https:// creativecommons.org/licenses/by/ 4.0/).

Physical methods of processing automobile tires include various methods of grinding aiming to obtain rubber crumb, which most fully preserves the properties of rubber [23,24]. The physical process of rubber grinding is rather complicated, since, due to its high elastic properties, the energy consumed for destruction is spent mostly on mechanical losses. The efficiency of mechanical rubber grinding is largely dependent on temperature and load application rate. If the grinding process occurs at a temperature below the glass transition temperature of rubber, then its deformations are small, and the destruction is brittle [25,26]. At the same time, it can be assumed that among physical methods, the low-temperature method is very promising [27,28].

The cryogenic method has the following advantages over the room temperature methods (i.e., when the rubber is in an elastic state): significantly less energy consumption for mechanical grinding; exclusion of fire and explosion hazard; production of a finely dispersed rubber powder with a particle size of up to 100 microns; elimination of environmental pollution [28]. In general, the effectiveness of cryogenic tire grinding is a consequence of: (a) bond weakening between metal cord and rubber at low temperatures, which leads to partial separation of rubber from metal; (b) sharp decrease in the elasticity of rubber and its brittle destruction even at minor deformations [27].

Bibliographic research has shown that it has not yet been clearly defined which of these methods is more efficient in terms of environmental impact, energy consumption and product application [29,30]. Articles [31,32] mentioned the unpromising nature of the cryogenic method, and also considered such issues as the method application and the economic benefit from cryogenic processing of automobile tires. The authors in [33,34] concluded that the cryogenic method is not economically feasible. These conclusions are mainly based on the cost of the primary material for cold production—liquid nitrogen. The market value of liquid nitrogen in 1980s was 4.05 \$/kg, but today the commercial price varies from 0.2 to 1 \$/kg, depending on the technical and economic capabilities of the region and the manufacturer. It should be noted that to assess the economic feasibility of using the cryogenic method, it is necessary to consider the prime cost of the refrigerant. As an example, we estimated the primary cost of one kilogram of liquid nitrogen produced in the laboratory of cryophysics and cryotechnology of Al Farabi Kazakh National University, which was 0.1 \$/kg. Obviously on an industrial scale this price is expected to be much lower (≈0.05 \$/kg). In addition, an indicator of the liquid nitrogen availability is the fact that it has become widely used in modern companies, ranging from cosmetology to industry [35–42]. Moreover, the use of air turbine refrigeration machines can reduce the cost of cold production by 3–4 times, and specific energy consumption by 2–3 times compared with the use of liquid nitrogen [43].

Thus, the simplification of rubber waste recycling using cryotemperatures is the subject of the present work aiming to reduce the required energy consumption. The object of the study is waste automobile tires. The temperature range of cryogenic treatment varied from 77 K to 280 K. The optimal force of mechanical shock for the tire destruction which turned into glassy state after cryoexposure was determined experimentally. The chemical and physical properties of the crumb rubber have been studied by scanning electron microscopy and energy dispersive X-ray spectroscopy.

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

### *2.1. Materials*

The object of the present research is automobile tires manufactured according to the standard "GOST 4754-97 (Pneumatic tires for passenger cars, trailers for them, light-duty trucks and buses of especially small capacity Specifications)" ("ISO 10191: 2010" Passenger car tires—Verifying tire capabilities—Laboratory test methods). In the experiment, onepiece waste tires were used to determine the parameters of mechanical impact.

### *2.2. Measurement of Thermal Conductivity of Samples*

Knowledge of the thermophysical properties of polymeric materials is important in many industries and technological development. The ability to identify the temperature dependence of polymer properties in a wide temperature range, including low temperatures, can play a particularly decisive role [25,44,45].

For computer simulations of the cooling process of a one-piece automobile tire, we used experimental data of thermal conductivity for a rubber sample. These experimental data were presented in our earlier work [46] in a graphical form. In this work, experimental measurements of the thermal conductivity of various samples were carried out in the low-temperature range from 95 K to 275 K. The samples were tire fragments consisting of pure rubber, rubber with nylon and metal cords. The data in [46] were used only for pure rubber and are shown in Table 1.

**Table 1.** Temperature dependence of thermal conductivity of a rubber sample.

As can be seen from Table 1, the value of thermal conductivity coefficient is maximum at 215 K, equal to 0.331 W/m · K. For the sample under investigation, this temperature is the glass transition temperature, i.e., *T*g = 215 K, the temperature at which there is a phase transition from an elastic to a brittle glassy structure occurs [46–48].

### *2.3. Mechanical Characterization*

To calculate the destruction force of mechanical impact on a glassy tire we performed several experiments using a sledgehammer. The experiment involved a pneumatic sledgehammer of type MA4129A (LLC South Ural Mechanical Plant, Kuvandyk, Russia) with dimensions of 830 × 1560 mm2 and a nominal mass of falling parts 80 kg. The maximum speed of falling parts at the moment of impact (theoretical) was 6.16 m/s. Minimum impact energy, not less than 155 kgf·m. The longest tup stroke was 365 mm. The rated power of the machine is 7.5 kW. The standard hammer strikers were replaced with profiled punches and dies (see Figure 1), especially made for the optimal sizes of automobile tires, in order to apply the mechanical shock entirely to the cooled tire.

The calculation of the impact force was made according to the following considerations: a hammer of mass *m* strikes a frozen sample with a speed *υ*1 and rebounds after an impact with a speed *υ*2. If we take into account that the impulse of the hammer *p*2 after the impact is directed oppositely to the initial impulse *p*1, then the change in impulse will be Δ*p=m* (*υ*1 *+ υ*2). The expression for the force impulse is written as: *F*Δ*t =* Δ*p*. The

average hammer force acting on the sample during the impact was determined, assuming the contact time of the hammer with sample equal to Δ*t* = 10−3 s.

**Figure 1.** Grooved contact part of the sledgehammer.

### *2.4. Mathematical Model*

The COMSOL Multiphysics program was used to simulate and calculate the optimal cooling time for a car tire. The main task of the modeling is the solution of non-stationary thermal fields described by the heat conduction equation, which belongs to the parabolic equations of the second order. The numerical experiment was performed using the "Heat Transfer" and "Computational Fluid Dynamics" modules of COMSOL Multiphysics. The Heat Transfer and Computational Fluid Dynamics (CFD) modules extend the capabilities of the COMSOL Multiphysics numerical simulation environment for the numerical analysis of systems in which hydrodynamic processes are accompanied by various physical phenomena [49].

Waste tire rubber and liquid nitrogen are modeled as two different computational domains, accompanied by the phenomenon of conjugate heat transfer. For each region, the corresponding differential heat conduction equations are used. These equations are shown below:

Heat transfer in the liquid nitrogen model:

$$
\rho\_f c\_f \frac{\partial T}{\partial t} + \nabla \cdot (-\lambda\_f \nabla T) + \rho\_f c\_f T \nabla \mathbf{u} = 0 \tag{1}
$$

where *ρf* is the liquid density, *cf* is the specific heat of liquid, *λf* is the thermal conductivity of liquid, and **u** is the velocity field. Equation (1) represents a homogeneous partial differential equation.

Heat transfer in the simulated tire:

Due to the absence of any mass flow **u** is taken to be zero in the simulated tire. Therefore, the basic heat transfer equation takes the following form:

$$
\rho\_s c\_s \frac{\partial T}{\partial t} + \nabla \cdot (-\lambda\_s \nabla T) = 0 \tag{2}
$$

where *ρs* is the density of the tire material, *cs* is the specific heat capacity of the tire, *λs* is the thermal conductivity of the tire.

The tire model consists of two elements—a rubber component and a metal (cord). The key point in the heat transfer process at the boundary of these elements is the use of the standard finite element method with matched grid cells with common nodal points

at which the solution to the problem is calculated. Grid elements that have common nodes and the continuity condition of the sought field variables are automatically defined. The balance of heat fluxes between elements belonging to different regions is adjusted automatically. Since the elements have common nodes, the temperature fields in these nodes are continuous. The fulfillment of the continuity condition is ensured by duplicating Equation (2) with the corresponding properties.

The CFD module includes special physics interfaces for simulating conjugate heat transfer between a fluid and a solid. For the laminar mode of the convective fluid flow, the condition of temperature continuity at the solid-liquid interface is used (this is the standard setting in non-isothermal flow interfaces). Non-isothermal laminar flow interfaces and conjugate heat transfer are required to conserve energy, mass, and momentum in fluid, and to conserve energy in solid. When solving this problem in COMSOL Multiphysics, the predefined Laminar Flow multiphysics interface in the Conjugate Heat Transfer branch automatically adds the Heat Transfer in Solids and Fluids interface. It is used to simulate heat transfer based on heat conduction and convection.

The following initial and boundary conditions are accepted in the computer model: *p* = 0—constant pressure at the upper boundary, **u** *=* **n***U*0 = 0—on the symmetry axis of the element, **u** *=* **u***f\_finite*—on the lower boundary, **u** = 0—on all other boundaries, initial ambient temperature is 77 K, simulating the temperature of liquid nitrogen.

The model grid was built choosing the optimal number of cells with a minimum loss of computational resolution. In this regard, the model was divided into 67,455 finite elements (triangles). The minimum grid element size is 0.1 mm, and the maximum is 23.5 mm. A fragment of the grid model is shown in Figure 2.

**Figure 2.** Fragment of the grid model for COMSOL computations.

### *2.5. Evaluation of Model Adequacy*

The evaluation of the computer model adequacy was performed by comparing the simulation results with experimental data obtained by registering the temperature change of a rubber sample during its cooling by liquid nitrogen. The sample has a cubic shape with geometric dimensions of 1.5 × 1.5 × 1.5 cm3, and it was derived from a massive piece of tire. A thermocouple (type T) was used as a thermal sensor and was inserted into the center of the rubber sample using a thin tubular catheter.

Furthermore, the sample was cooled by immersion in liquid nitrogen with simultaneous recording of its temperature change over time in an automatic mode using a PC. The process of recording the temperature sensor readings was carried out from the start of cooling until the sample was taken to the temperature of liquid nitrogen. The initial sample

cooling temperature was room temperature (293 K). The experimental results (curve 1) are shown in Figure 3.

**Figure 3.** Comparison of the experimental and simulation data. 1—experiment, 2—model.

The above experiment was simulated using the COMSOL Multiphysics environment with detailed description given in Section 2.4. Geometry in the computer model was taken in accordance with the dimensions of the experimental sample. The temperature range was set from 293 K to 77 K. The results of numerical experiment in comparison with the experimental data are shown by curve 2 in Figure 3. It can be seen from Figure 3 that the results of numerical simulation and experiment are in good agreement, and the deviation of the experimental data from numerical simulation curve does not exceed 10%.

### *2.6. Scanning Electron Microscope and Energy Dispersive X-ray Spectroscopy*

The morphology and elemental composition of the crumbs were studied by scanning electron microscopy (Quanta 200i 3D, FEI Company, Hillsboro, OR, USA) and energy dispersive X-ray spectroscopy (MicroXRF Analysis Report, EDAX int, Berwyn, PA, USA).

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

*3.1. Evaluation of the Mechanical Impact Force*

It is known that, depending on the temperature, natural rubbers can reside in one of the three states: glassy, highly elastic, and viscous-flow [50,51]. The elasticity of rubbers can be completely lost depending on the degree of cooling [52]. If the temperature dependence of elasticity is characterized by mechanical properties, then taking the value of deformation at a given strain and temperature as a characteristic of the rubber state, it is possible to obtain thermomechanical data. This, in turn, can make it possible to compare the value of deformation at different temperatures, and also to find the glass transition temperature corresponding to the transition of rubber to a brittle glassy state [53]. Glass transition depends not only on the temperature, but also on the nature of the mechanical load. Thus, under static loads or dynamic loads of low frequency, the glass transition temperature is lower than under dynamic loads of high frequency [54]. From this perspective, it is clear that the most accurate measurement of the cooling depth of samples and mechanical effect after freezing are necessary to optimize the low-temperature method for rubber waste recycling. As known, the duration of cooling depends on the mass, geometric parameters and thermophysical properties of waste tires [28].

The experiment on cryogenc grinding of car tires was carried out in few steps. On the first step cooling of the waste tire was done by complete immersion in special vessel, made for this purpose (Figure 4a). The time of immersion was determined starting from the immersion moment up to the termination of intense boiling of liquid nitrogen. The cooling time, depending on the mass and average thickness of the tire walls, varied from 200 s to 300 s. Here, the degree of cooling is very important, since it is necessary to have a temperature margin of brittleness, which can be spent due to the heat of dissipation of a mechanical shock.

**Figure 4.** Cooling process of automobile tire in liquid nitrogen (**a**) and mechanical impact on frozen tire (**b**).

Further the frozen waste tire was placed under the hammer (Figure 4b) and subjected to mechanical stress (impact) in a single mode. The strength of the mechanical impact and the number of blows required to completely destroy the tire were determined by varying the speed of the hammer blow. According to the results of tire destruction, the optimal force of one impact was calculated. If necessary, the number of blows can be increased to achieve the required degree of separation of the metal cord and nylon fibers.

The calculation of the force of mechanical impact was carried out experimentally as follows: a hammer of mass m = 80 kg strikes a frozen tire with a speed υ1 ≈ 2.5 m/s, and stops after the impact without rebound, i.e., with υ2 ≈ 0 m/s, spending all the impact energy on the destruction of an embrittled tire. Furthermore, making simple mathematical calculations, according to the considerations of Section 2.3, the required value of the force of mechanical shock was determined.

After the mechanical impact, the crushed rubber crumbs were sorted by size depending on the strength of the impact. The metal and nylon cord elements are almost completely separated from the rubber component (up to 95%), which can significantly reduce further costs for products recycling by other combined methods. It is important to note that this fact is the main advantage of the cryogenic method of recycling waste tires [28,55].

According to preliminary estimates, in one hammer blow, up to 80% of rubber contained in the tire passes into the crumb, with 50% of the crumb with dimensions from 1.25 to 20.00 mm and 25% from 0.10 to 1.25 mm. The evaluation showed that the force with an approximate value of ≈200 kN is a satisfactory result for the destruction of the frozen tire. Figure 5 shows samples of rubber crumb of various fractions with separated metal rods.

**Figure 5.** Typical products of cryogrinding. 1—nylon cord, 2—metal cord, 3—large size fraction of rubber crumb (1–5 mm), 4—mean size fraction (1–2 mm), 5—small size fraction (0.1–1 mm).

### *3.2. Computer Simulation*

Numerical experiment of the tire rubber cooling process was performed using axisymmetric 2D modeling. Many physics interfaces are available in axisymmetric versions and consider the axial symmetry. This allows solving a 2D problem in plane instead of a full 3D model problem, which can significantly save machine memory and computation time. After solving the problem in plane, we can get the result in a three-dimensional form by rotating the two-dimensional axisymmetric solution.

The simulated model of car tire has the following geometrical dimensions: width— 185 mm; tire profile—75%; landing diameter—13 inches. The sketch of the 3D model, which is shown by rotating the cross-section of the 2D model around the axis located at the distance of tire radius, is shown in Figure 6.

**Figure 6.** Computer model illustration.

Experimental data on the temperature dependence of the rubber thermal conductivity, given in Table 1, were used as the values of thermal conductivity *λs* of the simulated tire, and the values of density *ρs* and heat capacity *cs* were taken from tabular data [56,57]. For the coolant and the metal cord, the values of these physical quantities (*ρf*, *cf*, *λf*) were also taken from the tabular data [58,59].

The main pre-definitions of the COMSOL Multiphysics software package for numerical experiments are given in Section 2.4. To be clear, we note that the initial tire temperature was taken equal to T = 293 K, and the temperature of the surrounding liquid medium was set equal to T = 77 K (T = *const*). The process of conjugate heat transfer is carried out between two subsystems, i.e., heat exchange occurs between solid (tire) and liquid (liquid nitrogen).

Please note that the calculation of the heat transfer coefficient which characterizes the intensity of heat transfer on the surface of the simulated tire is essential. It is known that the use of the similarity equation for the dimensionless Nusselt number is a recognized method for the calculation of heat transfer coefficient. These equations make it easy to calculate the heat transfer coefficient for different heat transfer conditions. However, this method is used only for objects of regular geometric shapes, for example, flat surfaces, cylinders and spheres. If the heat transfer surface has a more complex shape, as in the present work, then the heat transfer coefficient is calculated by simulating the conjugate heat transfer. Thus, the present model was numerically solved in COMSOL Multiphysics using the Conjugate Heat Transfer interface, which allows us to calculate the flow and temperature fields in liquid. Based on the simulation results, the program calculates the heat flux density with reference to the corresponding built-in post-processing variable. Furthermore, by dividing the reported value of the heat flux density by the temperature difference (*Ts-TL*), the heat transfer coefficient can be found.

Figure 7 presents the model calculation results. The graph shows the decrease in average temperature of the simulated tire over time. Average temperature is the temperature of the grid nodes divided by the number of these nodes in the selected object. Curve 1 shows the decrease in average temperature of the model rubber component. In addition, curve 2 shows similar decrease for only the metal cord. As expected, the change over time in average temperatures of the subsystems from 300 K to 77 K occurs according to the exponential law (Newton's cooling law [60,61]). It can be seen from Figure 7 that the decline of curve 2 is faster than curve 1. This is explained by the difference in properties of the subsystems set for the simulated materials and the geometric locations of the grid nodes, in which the temperature changes over time are determined. It is clear that the values of the thermal conductivity coefficients of rubbers are two orders of magnitude lower than those of metals [62,63]. Therefore, determination of the average temperature for simulated part of the rubber with the cord at each moment of time occurs faster than the average temperature of the entire model area. In addition, the cord is not in the center, but closer to the border surface. If we assume that cryogenic treatment is performed on both sides of the interface, then the relatively rapid cooling of the rubber with cord in comparison with the entire model is quite obvious. Thus, the simulation is based on the cooling process of investigated object from the known initial temperature to the glass transition temperature [47] and further to temperature 77 K for cold margin at a given time. The duration of time is determined by the series of average values between time t = 0 and the estimated time. These values were calculated in accordance with the given thermal conductivity and temperature of the object using the calculation of moving average. The calculated time value is compared with the experimental cooling time presented in Figure 3. Good convergence between two values can be seen.

**Figure 7.** Tire cooling graph. 1—average temperature change in grid nodes for the entire model cross-section; 2—average temperature change in grid nodes for the cross-section of the metal cord.

Obviously, the thermal conductivity coefficient describes the relationship between the heat flux vector **q** and the temperature gradient Δ*T*, which are included in the Fourier law equation. Since the present model uses the experimental coefficients of rubber thermal conductivity in the low-temperature range (95–275 K), we plot the dependence of temperature gradient on the average temperature at two measured points in order to further check the model adequacy with experiment. In the COMSOL Multiphysics software package, the temperature gradient was determined between two points chosen arbitrarily according to the condition that they lie along the axis normal to the surface of the isotherm. This dependence is shown in Figure 8. In general, the temperature gradient values linearly decrease with decreasing temperature. However, in the vicinity of 220 K, the linear dependence of the temperature gradient undergoes a sharp decline. This is explained by the fact that in the vicinity of this temperature (220 K), the phase transition from the elastic to the glassy state occurs in rubber [46–48,64].

**Figure 8.** Temperature dependence of temperature gradient.

### *3.3. Physical and Chemical Properties*

Morphology and chemical composition of crumb rubber were studied using a scanning electron microscope (Quanta 200i 3D, FEI Company, USA) and energy dispersive X-ray spectroscopy (MicroXRF Analysis Report, EDAX int, USA). The purpose of these studies was to understand the physical and chemical properties of the crumb rubber after cryodestruction and to formulate recommendations for further use of recycled rubber crumbs.

Figure 9 shows the results of microscopic examination of the rubber crumbs. It can be seen that rubber particles have a smooth surface, characteristic of large particles, without microcracks and pores. The ability to produce rubber crumbs without the complex surface (microcracks and pores) is one of the advantages of the cryogenic method of polymers destruction. It is known that fine particles of any solid substance with a complex surface efficiently adsorb ambient gases, particularly atmospheric gaseous oxygen. Adsorbed atmospheric oxygen can lead to rapid oxidation of the polymer, therefore shortening its shelf life. In addition, the oxygenated polymeric material can be extremely flammable.

**Figure 9.** SEM—cryodestructed crumb rubber sample illustration.

Figure 10 shows the results of the elemental composition analysis for the tire sample before (a) and after (b) cryodestruction using energy dispersive X-ray spectroscopy (MicroXRF Analysis Report, EDAX int, Berwyn, PA, USA). A wide range of chemical compounds can be observed in the waste tire rubber, such as natural rubber, SBR (styrene butadiene rubber), butadiene rubber, etc. A large number of other chemicals are also added to tire rubber, such as: vulcanizing agents (sulfur and sulfur compounds), catalysts, inhibitors, pigments, etc. Vulcanizing agents and various additives are used in tires to create a cross-linked structure and provide resistance to various physical and chemical influences [65,66]. Other micro-additives such as calcium, magnesium, sodium, potassium, chloride, etc. are added to improve the mechanical performance. The complex composition of car tires requires performing a chemical content analysis for the presence of toxic elements harmful to the environment. It is also important to accurately identify the characteristics of the recycling products and by-products to assess the environmental footprint of the processing technique.

**Figure 10.** Elemental composition of crumb rubber before (**a**) and after (**b**) cryogenic destruction.

Thus, elemental composition of the crumb rubber after cryodestruction showed no significant difference with the energy dispersive analysis of the original tire sample before the cryogenic treatment.

### **4. Conclusions and Recommendations**

The annual reports of many international companies and scientific journals show that the growth rate of rubber and industrial rubber goods production is increasing annually. Consequently, the growth rates of waste rubber-containing products is also increasing. More than half of this waste consists of the waste tire rubber of various sizes.  Unfortunately, industrial development in its present form is accompanied by violation of the ecological balance as a result of human impact on the nature.

Based on the results of the present work, the following brief conclusions can be drawn: Mechanically stressed glass tire is cooled by full immersion in liquid nitrogen. The duration of tire cooling was determined experimentally. Analysis of the samples on the mechanical impact force showed that the force with approximate value 200 kN is a satisfactory result for the destruction of tire to crumb rubber.

A numerical model of the car tire cooling process showed satisfactory results, correlating with the experiment. These results can be useful in the design of cryogenic reactors for solid household polymer waste recycling.

The results of morphology and chemical composition studies of crumb rubber after cryodestruction can be used when crumb rubber is applied as a filler in various structural materials.

**Author Contributions:** Conceptualization, D.Y., A.T., D.S. and A.A.; methodology, D.Y., D.S. and A.A.; formal analysis, A.T.; supervision, A.A.; investigation, A.T.; data curation, D.S.; writing original draft preparation, D.Y. and A.A.; writing—review and editing, D.Y., A.T., D.S. and A.A.; project administration, A.A.; funding acquisition, A.A. All authors have read and agreed to the published version of the manuscript.

**Funding:** These studies have been carried out with the financial support of the Ministry of Education and Science of the Republic of Kazakhstan under grant AP08855738. Additionally this work was supported by the Postdoctoral Fellowship provided by Al-Farabi Kazakh National University. Responsibility for the information and views set out in this article lies entirely with the authors.

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** The data used to support the findings of this study are available from the corresponding author upon request.

**Acknowledgments:** These studies have been carried out with the financial support of the Ministry of Education and Science of the Republic of Kazakhstan under grant AP08855738. The authors express their appreciation and profound gratitude to Nurlan Tokmoldin, for the fruitful discussion and useful comments during the work on the present manuscript.

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

### **References**

### *Article* **Thermal Degradation Behavior of Epoxy Resin Containing Modified Carbon Nanotubes**

**Xiaohui Bao, Fangyi Wu and Jiangbo Wang \***

School of Materials and Chemical Engineering, Ningbo University of Technology, Ningbo 315211, China<EMAIL_ADDRESS>(X.B<EMAIL_ADDRESS>(F.W.)

**\*** Correspondence<EMAIL_ADDRESS>or<EMAIL_ADDRESS>Tel.: +86-0574-8708-1240

**Abstract:** Via the surface-grafting of carbon nanotubes (CNTs) with a silicon-containing flame retardant (PMDA), a novel flame retardant CNTs-PMDA was synthesized. The flame retardancy was tested by cone calorimeter. Compared with pure epoxy resin, the total heat release (THR) and peak heat release rate (PHRR) of epoxy resin containing CNTs-PMDA were significantly reduced, by 44.6% and 24.6%, respectively. Furthermore, thermal degradation behavior of epoxy resin based composite was studied by the thermogravimetric analysis with differences in heating rates. The kinetic parameters of the thermal degradation for epoxy resin composites were evaluated by the Kissinger method and Flynn-Wall-Ozawa method. The results suggested that activation energy values of epoxy resin containing CNTs-PMDA in thermal degradation process were higher than those of pure epoxy resin in the final stage of the thermal degradation process, which was closely related to the final formation of char layer residues. Finally, the results from Dynamic mechanical thermal analysis (DMTA) and Scanning electron microscopy (SEM) measurements exhibited that the functionalization of CNTs with PMDA obviously improved the dispersion of CNTs in the epoxy resin matrix.

**Keywords:** epoxy resin; carbon nanotubes; silicone; thermal degradation; flame retardancy; activation energy

### **1. Introduction**

Epoxy resins (EP) have been developed for nearly 70 years; their applications cover such varied fields as composite matrices, surface coatings, adhesives, and encapsulation of electronic components, as well as the aeronautical and astronautical industries. Features of epoxy resins such as their excellent mechanical properties, good chemical and moisture resistance, long-term service time, low cost and easy processing are often attractive [1–4]. Nevertheless, the inherent flammability of EP limits its application in many fields due to safety concerns. Therefore, flame retardants have to be used in order to fulfill the requirements of flammability tests [5–7].

In general, carbon nanotubes (CNTs) have been of interest to researchers because they frequently exhibit superior flame retardancy, mechanical and electronic properties. Advanced improvements when compared with conventional flame retardant containing halogen, they are free of toxic smoke or corrosive fumes during combustion [8–10]. By adding 2 wt% carbon nanotube flame retardant to polymer, flame retardancy can be significantly improved without damaging other properties of the polymer. Moreover, carbon nanotubes can also be prepared by extracting lignin from bamboo charcoal, rice straw, coconut fibre and corn straw as carbon sources. Therefore, it has excellent environmentally friendly properties and recyclability [11–13]. However, the agglomeration of CNTs owing to strong van der Waals forces and π–π interactions between carbon nanotubes make it quite impossible to realize the ideal functions of CNTs in polymer nanocomposites [14,15]. To enhance interfacial adhesion between nanotubes and the polymer matrix, chemical functionalization of carbon nanotubes is one of the applicable methods. It can be concluded

**Citation:** Bao, X.; Wu, F.; Wang, J. Thermal Degradation Behavior of Epoxy Resin Containing Modified Carbon Nanotubes. *Polymers* **2021**, *13*, 3332. https://doi.org/ 10.3390/polym13193332

Academic Editors: Sheila Devasahayam, Raman Singh and Vladimir Strezov

Received: 25 August 2021 Accepted: 27 September 2021 Published: 29 September 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 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 (https:// creativecommons.org/licenses/by/ 4.0/).

that the polymers grafted onto CNTs could desirably improve their dispersion but also simultaneously destroy the flame retardancy of the materials due to the polymer's inherent flammability [16,17]. Nowadays, solving these problems usually requires the functionalization of CNTs using some reactive flame retardants such as fullerene, intumescent flame retardant, etc. [17–20]. On the other hand, thermal degradation kinetic methods are widely used to analyze the combustion processes of polymers, including epoxy resin, polycarbonate, polyvinyl chloride, etc. [21–23]. Understanding the role of flame retardants in the combustion process of polymers is helpful to reveal their flame retardant mechanism.

In this article, a novel silicon-containing flame retardant (PMDA), which was composed of 35 mol% methylsiloxane, 60 mol% phenylsiloxane and 5 mol% aminosiloxane units, was prepared by hydrolysis and polycondensation. Then, the PMDA was grafted onto the surface of the CNTs, and the epoxy resin/CNTs-PMDA nanocomposites were subsequently prepared. The flame retardancy and thermal degradation behavior of the flame-retardant epoxy resins were determined by the cone calorimeter and thermogravimetric analysis (TGA), respectively, for the purpose of gaining insight into the possible flame retardant mechanism.

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

### *2.1. Materials*

Methyltrimethoxysilane (MTMS, 98%), tetramethylammonium hydroxide (TMAOH, 97%), pyridine (99.5%), tetrahydrofuran (THF, 99%), concentrated sulphuric acid (H2SO4, 98%), nitric acid (HNO3, 65–70%), (3-aminopropyl)trimethoxysilane (APS, 97%) and N,Ndimethylformamide (DMF, 99%) were all obtained from Alfa Aesar Chemical Reagent Co. Ltd. (Tewksbury, MA, USA). Dimethyldimethoxysilane (DMDS, 95%), phenyltrimethoxysilane (PTMS, 97%) were both reagent grade and provided by Gelest Chemical Reagent Co. Ltd. (Morrisville, PA, USA). Ethyl alcohol (EtOH, 95%) was supplied by Sigma-Aldrich Reagent Co. Ltd. (St. Louis, MO, USA). Thionyl chloride (SOCl2, 98%) and chloroform (CHCl3, ≥99%) were purchased from Fisher Scientific Chemical Co. (Waltham, MA, USA). EPON 826 containing an epoxy equivalent weight of 178–186g was offered by Hexion Specialty Chemicals Inc. (Columbus, OH, USA) and used as received. The hardener with an amine equivalent weight of 60 g, Jeffamine D230, was purchased from Huntsman Corp. (Woodlands, TX, USA) and also used as received. CNTs (length at 10–30 μm, outer diameter at 10–20 nm, inner diameter at 5–10 nm,) synthesized by chemical vapor deposition were supplied by Chengdu Organic Chemistry Co. Ltd., Chinese Academy of Science (Sichuan, China). It was composed of 99.4% carbon nanotubes and 0.6% catalyst Ni, and contained no conventional carbon products such as graphite and carbon black.

### *2.2. Synthesis of Polysilicone (PMDA)*

The polysilicone PMDA was synthesized by the hydrolysis and polycondensation method as shown in Scheme 1. The raw material consisted of 60 mol% phenylsiloxane, 35 mol% methylsiloxane and 5 mol% aminosiloxane as a unit. The ratio of organic groups to silicon atoms (R/Si) was 1.2, which was used to characterize the extent of branches to a polysiloxane structure; the molecular structure of polysilicone is illustrated in Scheme 1. Distilled water (25 mL), EtOH (75 mL) and TMAOH (1 mL) were mixed in a 250 mL flask under stirring, followed by adding the mixture of PTMS, MTMS, DMDS and APS at certainly molar ratios (0.69:0.06:0.20:0.05) and maintaining 10% weight percentage solution. The stirring was maintained for 8 h, and the resulting solution was stored at room temperature overnight. Through decantation of most clear supernatant, precipitated condensate was collected and then washed by vacuum filtration with distilled H2O/EtOH (1/3 by volume), then washed again in pure EtOH. Finally, the acquired rinsed powder (PMDA) was thoroughly dried under vacuum for 20 h at room temperature [24].

**Scheme 1.** Synthesis route of PMDA.

### *2.3. Functionalization of CNTs*

The synthesis steps of CNTs-COCl were as follows: The mixture of HNO3 (30 mL), H2SO4 (90 mL) and CNTs, was sonicated at 50 ◦C for 2 h and, after termination of the reactioncould be cooled to room temperature. The mixture was diluted with deionized water, then underwent vacuum filtering using a nylon film (0.22 μm, Sangon Biotech Co. Ltd., Shanghai, China). The obtained solid was CNTs-COOH, whose polar carboxyl groups were successfully introduced into the convex surface of CNTs, then washed with a large amount of deionized water until the aqueous layer reached neutral. The solid was vacuum dried at 80 ◦C for 12 h. In the next step, the reaction mixture of CNTs-COOH (200 mg), SOCl2 (20 mL) and DMF (1 mL) was sonicated at 50 ◦C for 1 h, after which the reflux was maintained at 70 ◦C for 24 h. Finally, the temperature was raised to 120 ◦C to remove residual SOCl2 via reduced pressure distillation, and CNTs-COCl correspondingly obtained.

In the next step, the synthesis of CNTs-PMDA was as follows (Scheme 2): the suspension of CNTs-COCl (100 mg) and DMF (50 mL) was introduced by PMDA (400 mg) and pyridine (1 mL, as cat.), under the protection of nitrogen. The mixture was reacted at 70 ◦C for 24 h. When cooling to room temperature, the dark solution was filtered and washed to remove unreacted PMDA. The target product, CNTs-PMDA, was obtained after vacuum drying at 80 ◦C for 24 h [24].

**Scheme 2.** Illustration for the functionalization of the CNTs with PMDA.

### *2.4. Preparation of Epoxy Composites*

Briefly, the EP/CNTs-PMDA composite was prepared as follows: the above semiproduct, CNTs-PMDA (2 g), was dispersed in acetone (100 mL) and sonicated for 1 h until forming a uniform black suspension. EPON 826 (73.5 g) was added into the mixture and then mechanically dispersed through stirring for 30 min. The mixture was heated in a vacuum oven at 50 ◦C for 10 h to remove the solvent thoroughly. Subsequently, D230 (24.5 g) was added into the mixture while stirring was maintained for 30 min [25]. The composite

was put into vacuum again tobe degassed for 10 min, in order to remove any trapped air. Finally, the sample was cured at 80 ◦C and post cured at 135 ◦C for 2 h, respectively. For comparison, pure EP was also prepared underthe same processing conditions [26].

### *2.5. Characterization and Measurement*

The Fourier transform infrared spectroscopy (FTIR) spectra was tested using a Digilab Scimitar FTS-2000 IR spectrometer (Digilab Inc., Hopkinton, MA, USA). The surface morphology of carbon nanotubes was observed by JEOL JEM-2100F transmission electron microscopy (TEM, JEOL Ltd., Akishima-shi, Tokyo, Japan). Cone calorimeter measurement was performed on an FTT cone calorimeter (Fire Testing Technology Ltd., East Grinstead, West Sussex, UK) according to ASTM E1354. The heat flux was 50 kW/m2. The dimensions of each specimen were 100 × 100 × 3 mm3. Thermogravimetric analysis (TGA) was carried on a TA instrument Q5000 thermogravimetric analyzer (TA Instrument Corp., New Castle, DE, USA). The sample (about 10 mg) was heated from 50 ◦C to 600 ◦C in a nitrogen atmosphere. Dynamic mechanical thermal analysis (DMTA) was determined using a Rheometric Scientific SR-5000 dynamic mechanical analyzer (Rheometric Scientific Inc., West Yorkshire, UK), and the data were collected from 40 ◦C to 140 ◦C at a scanning rate of 5 ◦C/min. The samples were coated with a conductive gold layer and examined by scanning electron microscopy (SEM) using an FEI Quanta 200 environmental scanning electron microscope (FEI Co., Hillsboro, OR, USA).

### *2.6. Thermal Degradation Theory*

The reaction rate of thermal transformation of a solid state chemical reaction based on the assumption is:

$$r = \frac{da}{dt} = kf(a) \tag{1}$$

where *f*(*a*) is the assumed reaction model, *a* is the degree of conversion, *k* the temperature dependent rate constant, *t* the time and *r* the rate of degradation. *k* is normally presumed to conform to the Arrhenius equation:

$$k = A \exp\left(-E/RT\right) \tag{2}$$

where *E* is the activation energy of the kinetic process, *A* the pre-exponential factor, *T* the temperature and *R* the universal gas constant.

The rate of degradation is dependent on the temperature and the weight change of the sample, and can be expressed as [26]:

$$\frac{da}{dt} = Af(a) \exp(-E/RT) \tag{3}$$

Equation (3) is derived in its integral form, which particularly for isothermal conditions becomes

$$
\ln t = E/RT - \ln[A/\,\mathrm{g}(\mathrm{x})] \tag{4}
$$

Another condition for non-isothermal degradation, Equation (3) becomes

$$\frac{da}{dT} = (A/\beta)f(a)\exp(-E/RT)\tag{5}$$

where *β* is the heating rate (*β* = *dT dt* ) and *g(x)* is the integrated form of mechanism (*g*(*x*) = *a* 0 *da f*(*a*)).

### 2.6.1. Kissinger Method

The Kissinger expression [27] is as follows:

$$\ln\left(\frac{\beta}{T\_{\text{max}}^2}\right) = \ln\left(\frac{AR}{E}\right) - \frac{E}{RT\_{\text{max}}}\tag{6}$$

where *T*max is the temperature of the peak rate.

From the above Equation (6), the peak rate temperatures determined at certain different heating rates allow the activation energy to be calculated by the Kissinger method. Plotting the natural logarithm of ln(*β*/*T*2 max) against the reciprocal of the absolute temperature (1/*T*max), the slope of the resulting line is given by −*E/R*, which allows the value of *E* to be obtained [28].

### 2.6.2. Flynn-Wall-Ozawa Method

The equation of the Flynn-Wall-Ozawa method [29,30] can be expressed as follows:

$$\log(\beta) = \lg AE / \lg(a) \text{R} - 2.315 - 0.457 \frac{E}{RT} \tag{7}$$

The above equation shows that lg(*β*) is linearly proportional to 1/*T*. The activation energy for any particular degree of degradation can then be determined by a calculation of the slope from the lg(*β*) − 1/*T* plots [28].

### **3. Results**

### *3.1. Structural Characterization*

Figure 1 exhibits and compares the FTIR spectra of PMDA, CNTs and CNTs-PMDA. As shown in the pristine CNT spectra at the top, the bands at around 1650 and 3500 cm−1 were attributed to the presence of carbonyl (quinone) and hydroxyl groups, respectively, which means that the CNTs contained some impurities [31,32]. For the middle spectrum, the reaction of CNTs with the PMDA led to the formation of an amide group, which exhibited close wavenumber bands with very close carboxylic groups (the strongpeak at 1653 cm−1 was attributed to the carbonyl in -COOH and -CONH- structures). The band at around 1558 cm−1 characteristic of a monosubstituted amide is evidence of the formation of an amide bond. Moreover, compared with the CNTs spectra other bands at 1105 cm−1 (Si-O-Si) and 2930 cm−1 (-CH2-) in CNTs-PMDA appeared, demonstrating that the PMDA had been grafted to the surface of the functionalized CNTs [24].

The TEM measurement of CNTs and CNTs-PMDA can provide direct morphological evidence for the grafting of PMDA onto the surface of the CNTs (Figure 2). As can be seen from Figure 2a, the carbon tube diameter of pristine CNTs was between 15~30 nm, and the surface was smooth and clean without any adhesion. In contrast, the surface of CNTs-PMDA in Figure 2 was rough, and its tube diameter was between 40~50 nm, which was significantly larger than that of pristine CNTs. Therefore, it can be proven from the appearance of CNTs and CNTs-PMDA that PMDA was indeed grafted onto the surface of the carbon tubes.

### *3.2. Flame Retardancy*

Cone calorimeter measurement can provide a lot of information on the combustion properties of the studied materials. Figure 3 illustrated the heat release rate (HRR) and total heat release (THR) as functions of time for epoxy resin (EP) and EP/CNTs-PMDA (FREP). After ignition, pure EP exhibited a single peak of basic pairs in the HRR curve with a peak height of 1742 kW/m2. Presence of CNTs-PMDA in EP dramatically decreased the peak heat release rate (PHRR) value for EP by 44.6% and increased the time-to-PHRR (tPHRR). In addition, for the FREP the THR was also significantly reduced, by 24.6% compared with that of pure EP. This is attributed to the initial formation of the char layer, which inhibited the heat from pyrolyzing the underlying matrix.

**Figure 1.** FTIR spectra of PMDA, CNTs and CNTs-PMDA.

**Figure 2.** TEM images of CNTs (**a**) and CNTs-PMDA (**b**).

### *3.3. Thermal Stability*

The TGA and DTG curves of EP composites at 10 ◦C/min are depicted in Figure 4, and the relative data summarized in Table 1. Compared with pure EP, the temperatures of 5% mass loss (*T*5wt%) and the maximum mass loss rate (*T*max) of FREP were raised by 15.4 ◦C and 18.7 ◦C, respectively. The use of 5 wt% CNTs-PMDA caused the peak rate to decrease from 2.10 to 1.97 wt%/◦C, which was helpfulfor improving the flame retardancy of EP with CNTs-PMDA. Furthermore, the addition of CNTs-PMDA led to an increase from 8.27 to 10.86 wt% in the residue char amounts. Therefore, the flame retardancy of EP was improved on the basis of CNTs-PMDA. This is further proven by the following activation energy data (Table 1).

**Figure 3.** HRR (**a**) and THR (**b**) curves for EP composites.

**Figure 4.** Thermal stability for EP composites.

**Table 1.** TGA data of EP composites.

### *3.4. Thermal Degradation Kinetics*

As a tool for unraveling the mechanisms of the physical and chemical processes that occur during polymer degradation, the application of non-isothermal TGA methods holds great promise. The TGA and DTG curves of EP composites at different heating rates (5, 10, 20, 40 ◦C/min) in nitrogen were depicted in Figure 5. It can be seen that with increasing heating rate, degradation started at higher temperatures. The reason for this transformation is that higher degradation temperatures at higher heating rates are due to the reduction in residence time, which is not sufficient for heat to permeate to the center of the reactants. Thus, the thermal decomposition process is delayed.

**Figure 5.** TGA and DTG curves of (**a**) EP and (**b**) FREP composites.

As far as the results of the Kissinger method were concerned, the *E* and ln*A* values of degradation for the EP composites were determined from the slope and intercept of the ln(*β*/*T*2 max) vs. 1/*T*max linear equations (Figure 6), respectively, presented in Table 2.

**Table 2.** Kinetic parameters for the degradation of EP and FREP by the Kissinger method.

It was be found that the *E* values of pure EP and FREP were 122.1 and 117.0 kJ/mol, respectively, which indicated that the addition of CNTs-PMDA led to a decrease in the value of the EP composites. The results confirmed that the flame retardant CNTs-PMDA induced the thermal degradation of polymer.

For comparison purposes, the activation energies for the EP composites under thermal degradation were also calculated by the Flynn-Wall-Ozawa method. According to Equation (7), the temperature of specific conversion rates such as a = 2%, 5%, 10%, 20%, 30%, 40%, 50%, 60%, 70%, 80%, 90%, 95%, 98% can be achieved from the TGA curves at various heating rates; the results were then plotted in Figure 7. Figure 7 shows that with the increase of the heating rate at isoconversion rate, the temperature increased significantly.

**Figure 6.** The curves of ln(*β*/*T*2 max) vs. 1/*T*max of EP and FREP.

**Figure 7.** Conversion of (**a**) EP and (**b**) FREP in function of temperature.

Figure 8 illustrates lg(*β*) vs. 1/*T*. The activation energy can be calculated from the slope of the plots of lg(*β*) vs. 1/*T* by Equation (7) with a conversion rate from 2% to 98%. According to the plots, there was a good linear relation between lg(*β*) and 1/*T*.

The values of *E* vs. *α* were exhibited in Figure 9. For pure EP, the changes of activation energies were minimal in the range of conversion rates from 0.02 to 0.80, so that kinetic parameters could be regarded as fixed. When *α* > 0.8, an obvious change of *E* occurred, which was related to the formation of the char layer and to the change in kinetic parameters. Compared with the *E* of pure EP, the variation of FREP had a similar law and the *E* of thermal degradation increased with the increase of the *α*. However, the *E* values of FREP were higher than that of pure EP in the final stage, which suggested that it was closely related to the formation of the final char residue layer.

**Figure 8.** The curves of lg(*β*) vs. 1/T of EP (**a**) and FREP (**b**).

**Figure 9.** Activation energy curves of EP and FREP by Flynn-Wall-Ozawa method.

This is consistent with the reports in the literature. The residue of pure EP had a lot of open holes with different sizes, which provided channels for the combustible volatiles from the inner matrix and for heat feedback from the flame [32]. After the functionalization of CNTs with flame retardants (such as silicone, phosphorus nitrogen compound), the residue char of FREP was very compact and dense, and few open holes were found. This can be explained by the wrapped CNTs forming a continuous network structured protective layer so that the entire structure could then effectively act as a barrier to limit the diffusion of flammable gases to the surface and slow down the combustion and degradation of the polymer [32–34].

### *3.5. Dispersion*

The mechanical properties of CNT-based composites are closely related to the dispersion of carbon tubes in the polymer matrix. Therefore, it is very important to study the dispersion of CNTs in the polymer matrix. DMTA and SEM measurements of composites can provide rich information in this regard. Figure 10 shows the DMTA curve of the EP composite. It can be seen that the tangent delta curves of EP/CNTs and EP/CNTs-PMDA exhibited a double-peak distribution, which indicates that there was an obvious microphase separation between the carbon tubes and polymer matrix. It was found that the distance between the two peaks of EP/CNTs-PMDA (6.0 ◦C) was less than that of EP/CNTs (9.4 ◦C), that is, the degree of phase separation was reduced. This showed that the grafting modification of PMDA on the surface of CNTs is conducive to improving the dispersion of carbon tubes in the polymer matrix.

**Figure 10.** Dynamic storage modulus (**a**) and tan δ (**b**) of EP composites.

Figure 11 shows the SEM images of CNTs and CNTs-PMDA in the EP matrix; the carbon nanotube aggregate are indicated by the box. As shown in Figure 11a, it was evident that many large CNT aggregates were present in the EP nanocomposite, and the pristine CNT bundles or ropes are more obvious under higher magnification. In comparison, the SEM images in Figure 11b clearly displayed that CNTs-PMDA dispersed more homogeneously in the polymer matrix than the pristine CNTs in Figure 11a, although some nanotube agglomerates still appeared. This indicates that after being grafted with PMDA, the interfacial compatibility between CNTs and the EP matrix was remarkably enhanced. Subsequently, the dispersion of CNTs in the polymer matrix was improved, which was completely consistent with the study results of DMTA measurement.

**Figure 11.** SEM images of (**a**) CNTs and (**b**) CNTs-PMDA in the EP matrix.

### **4. Conclusions**

In this paper, the flame retardant CNTs-PMDA was synthesized via the surface grafting of CNTs with PMDA. The results of cone calorimeter measurement indicated that the incorporation of CNTs-PMDA significantly decreased the peak heat release rate (PHRR) and total heat release (THR) of EP composite by 44.6% and 24.6%, respectively. Furthermore, kinetic studies using both Kissinger method and Flynn-Wall-Ozawa method revealed that a higher activation energy of FREP was obtained in the final stage compared with pure EP; this was related to the formation of the final residual char layer. Finally, the results from DMTA and SEM measurements exhibited that the functionalization of CNTs with PMDA obviously improved the dispersion of CNTs in the EP matrix.

**Author Contributions:** Conceptualization, X.B., F.W. and J.W.; methodology, X.B. and J.W.; validation, F.W. and J.W.; formal analysis, X.B.; investigation, X.B.; supervision, J.W.; writing—original draft preparation, X.B.; writing—review and editing, F.W. and J.W.; project administration, J.W.; funding acquisition, J.W. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the Ningbo Natural Science Foundation (2019A610032), National Undergraduate Training Program for Innovation and Entrepreneurship<PHONE_NUMBER>02) and Chongben Foundation. This work was also supported by the Open Fund of Shanghai Key Laboratory of Multiphase Materials Chemical Engineering.

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** The data used to support the findings of this study are available from the corresponding author upon request.

**Acknowledgments:** We gratefully acknowledge the financial support of the above funds and the researchers of all reports cited in our paper.

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

### **References**

### *Article* **Polypropylene Contamination in Post-Consumer Polyolefin Waste: Characterisation, Consequences and Compatibilisation**

**Erdal Karaagac, Mitchell P. Jones, Thomas Koch and Vasiliki-Maria Archodoulaki \***

Institute of Materials Science and Technology, Faculty of Mechanical and Industrial Engineering, Technische Universität Wien, 1060 Vienna, Austria<EMAIL_ADDRESS>(E.K<EMAIL_ADDRESS>(M.P.J<EMAIL_ADDRESS>(T.K.)

**\*** Correspondence<EMAIL_ADDRESS>
**Abstract:** Plastic recycling strikes a balance between functional, mass producible products and environmental sustainability and is pegged by governments for rapid expansion. However, ambitious targets on recycled material adoption across new markets are at odds with the often heterogenous properties of contaminated regranulates. This study investigated polypropylene (PP) contamination in post-consumer low-density polyethylene (PE-LD) and mixed polyolefin (PO) regranulates. Calibration curves were constructed and PP content, its effect on mechanical properties and property recovery in compatibilised material assessed. FT-IR band ratios provided more reliable estimations of PP content than DSC melt enthalpy, which suffered considerable error for PP copolymers. PE-LD regranulates contained up to 7 wt.% PP contamination and were considerably more brittle than virgin PE-LD. Most mixed PO regranulates contained 45–95 wt.% PP and grew more brittle with increasing PP content. Compatibilisation with 5 wt.% ethylene-based olefin block copolymer resulted in PE-LD blends resembling virgin PE-LD and considerable improvements in the properties of mixed PO blends. These results illustrate the prevalence of PP in recycled PE, challenges associated with its quantification, effect on mechanical properties, and compatibilisation viability, thereby representing an important step towards higher quality regranulates to meet the recycling demands of tomorrow.

**Keywords:** post-consumer waste; mechanical recycling; polyethylene; polypropylene; contamination; composition; tensile properties; impact properties; compatibilisation

### **1. Introduction**

With growing emphasis and legislative action on improving environmental sustainability, plastic recycling has become a forerunner in the race to optimise waste management practices, reduce reliance on fossil fuels and adopt closed loop circular economy principles across the globe. The European Commission aims to increase the use of recycled plastics in new products to 10 million tons/year by 2025 [1] and recycle 55% of all plastic packaging waste across the EU by 2030 [2]. Polyolefins (PO), such as polyethylene (PE) and polypropylene (PP), are popular packaging materials and represent more than half of the 29 million tons of plastic waste collected in the EU each year [3,4]. Such numbers very effectively illustrate the importance PO will play in meeting new recycling targets. However, to do so, the use of these recycled plastics will need to be expanded to new applications for which recycled material hasn't traditionally been suitable due to issues with material properties stemming from contamination [4].

Recycled plastics often contain both inorganic and polymer-based contamination, which results in mechanical properties that diverge from application-specific targets and limit their use to sectors utilising lower grade materials, such as agriculture and construction [5]. This contamination often results from flawed sorting practices, which struggle to economically separate materials with very similar characteristics, such as polymers of very similar density [6]. PP is a common contaminant in both low-density polyethylene (PE-LD) and mixed PO regranulates [7]. Several recent studies generically characterise

**Citation:** Karaagac, E.; Jones, M.P.; Koch, T.; Archodoulaki, V.-M. Polypropylene Contamination in Post-Consumer Polyolefin Waste: Characterisation, Consequences and Compatibilisation. *Polymers* **2021**, *13*, 2618. https://doi.org/10.3390/ polym13162618

Academic Editors: Sheila Devasahayam, Raman Singh and Vladimir Strezov

Received: 22 June 2021 Accepted: 4 August 2021 Published: 6 August 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 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 (https:// creativecommons.org/licenses/by/ 4.0/).

selected PE regranulates, providing differential scanning calorimetry (DSC) thermograms and Fourier transform infrared (FT-IR) spectra, as well as thermal degradation and mechanical properties, such as tensile, impact, fracture toughness, and hardness [4,8–10]. The effect of contamination on recycled polyolefins, including that from other polyolefins, and methodologies for the identification of such contaminants are also thoroughly documented in the literature [11–16]. However, a structured investigation centered on the characterisation of the PP constituent of PP contaminated PE blends, its effect on mechanical properties, and improvements possible through compatibilisation is clearly lacking.

This study aimed to investigate PP contamination in post-consumer PE-LD and mixed PO regranulates using the PP content as a base parameter and reference point for analysis rather than arbitrarily characterising these very heterogenous and often incomparable materials. DSC-based melt enthalpy and FT-IR band ratios were used to construct calibration curves to estimate blend PP content across a range of different PP types and their reliability contrasted. Tensile and tensile impact mechanical properties were then assessed and the influence of the PP content on these properties investigated. Finally, the regranulates were compatibilised to improve their mechanical properties and achieve a viable recycled substitute for virgin PE.

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

### *2.1. Materials*

Film grade 290E low density polyethylene (PE-LD) (Dow Chemical Company, Midland, MI, U.S.A), blow moulding grade Hostalen GF4750 high density polyethylene (PE-HD) (LyondellBasell, Rotterdam, The Netherlands) and injection moulding grade HF700SA polypropylene (PP) (Borealis, Vienna, Austria) were purchased as reference materials and used to construct melting enthalpy and FT-IR band ratio calibration curves for calculating PP content in PE-LD and mixed PO blends. HD601CF film grade, HC600TF thermoforming grade and HA104E extrusion grade homopolymer PP, BA202E extrusion grade block- and RD208CF random copolymer PP were purchased from Borealis (Vienna, Austria) to investigate the effect of various types of PP common in regranulates on melting enthalpy and FT-IR band ratio calibration curves. P01-1,2,3, P03-1,2,3, and P05-1,2,3 PE-LD regranulates were provided by Walter Kunststoffe Regranulat (Gunskirchen, Austria). Purpolen PE Grau (PPE), Purpolen PP Grau (PPP), Dipolen H (DPH), Dipolen PP Grau (DPP), Dipolen S Grau (DPS), and Dipolen SP Grau (DSP) mixed PO regranulates were provided by MTM Regranulat (Niedergebra, Germany). INFUSE ethylene-based olefin block copolymer with glass transition and melting temperatures of −65 ◦C and 118 ◦C, respectively, and a tensile elongation at break of 1000% was provided by Dow Chemical Company (Midland, MI, USA) for use as compatibiliser. All materials were used as received.

### *2.2. Prepatation of Virgin, Regranulate and Compatibilised Blends*

Blends of virgin PE-LD and PP used to construct melting enthalpy and FT-IR band ratio calibration curves for PE-LD regranulates were prepared using a HAAKE MiniLab II twin-screw extruder (Thermo Scientific, Waltham, MA, USA) running at 180 ◦C with a screw speed of 100 rpm and dwell time of 5 min. These lower temperature conditions were selected to minimize material degradation during extrusion. PP contents of 2, 5, 8, 10, 12, 15, 20, and 25 wt.% were weighed and hand mixed with PE-LD prior to extrusion. Melt mass-flow rates were calculated in accordance with ISO 1133-1:2011 [17].

Virgin PE-HD and PP blends were used to construct melting enthalpy and FT-IR band ratio calibration curves for mixed PO regranulates. These were prepared using an Extron-Mecanor SWL0914-1 single screw extruder (Toijala, Finland) with a nozzle temperature of 200 ◦C running at 75 rpm. PE-HD blends with a PP content of 2, 8, 10, 20, 30, 40, 50 60, 70, 80, 90, 92, 95, and 98 wt.% were weighed, hand mixed and extruded twice. Pure PE-HD and PP references were also prepared under the same conditions. An overview of the calibration curves used in this study, the materials, and sample compositions used to

generate them and the regranulates subsequently assessed with each curve, is provided in Table 1.

**Table 1.** Overview of the calibration curves utilised in this study, the materials and sample compositions used to generate them and the regranulates that were subsequently analysed using each respective calibration curve.

PE-LD (P01, P03 and P05) and mixed PO (PPE, PPP, DPH, DPP, DPS and DSP) regranulates were again prepared using a HAAKE MiniLab II twin-screw extruder (Thermo Scientific, Waltham, MA, USA). Extruded plastic was collected from the die, cut into small pieces and compression moulded into sheets at 190 ◦C using a Collin P 200 P laboratory press. Preheating was completed at 150 ◦C and 8 bar (hydraulic press pressure) for 10 min, heating from 150–190 ◦C at 22 bar for 8 min, followed by compression at 190 ◦C and 30 bar for 5 min. Samples were then cooled from 190 ◦C to 30 ◦C over 20 min at 10 K/min and 35 bar. Blends compatibilised with 5 wt.% ethylene-based olefin block copolymer were prepared in the same way. The compatibiliser and its content (5 wt.%) were selected based on an extensive study of PE-HD compatibilisation, recommending its use in quantities of 4–8 wt.% [7].

### *2.3. Melting Enthalpy and FT-IR Band Ratio Characterisation and Calibration Curve Generation*

Melting enthalpy was assessed using a TA Instruments Q 2000 differential scanning calorimeter (DSC) (New Castle, DE, USA). An ~8 mg sample mass of each polymer blend was deposited in an alumina testing pan and sealed. Samples were heated to 200 ◦C at 10 K/min, cooled at the same rate to room temperature, and then reheated under the same heating conditions as previously described. A nitrogen atmosphere was maintained at all times using a flow rate of 50 mL/min. The melting enthalpy Δ*Hm* of the second heating run was analysed using TA Instruments Universal Analysis 2000 (v. 4.5A, b. <IP_ADDRESS>). Analysis was based on four replicate specimens for each sample type.

The use of calibration curves based on DSC melting enthalpy to calculate the composition of polyolefin blends is documented in the literature [13–16]. Melting enthalpy-based calibration curves were constructed based on the known PP content of the series of virgin PE-LD/PP and PE-HD/PP blends (X axis) and the melting enthalpy (J/g) defined as the area under the melting peak on the second heating run (Y axis). A linear fit was applied to the resulting points using OriginPro 2019b (v. <IP_ADDRESS>) and the equation noted. The PP content of PE-LD and mixed PO regranulates could then be calculated by measuring their melting enthalpies using DSC and substituting these values into the equation to provide a solution.

IR spectra were recorded using a Bruker TENSOR 27 Fourier transform infrared (FT-IR) instrument in attenuated total reflection (ATR) mode. Three spectra were recorded from different portions of four individual samples to verify homogeneity. Spectra were recorded from 4000 to 400 cm−1.

FT-IR band ratio-based calibration curves enabling the calculation of the PP content in PE-LD regranulates were constructed based on the absorbance bands (amplitude) of the series of virgin PE-LD/PP blends of known composition at 1376 cm−1 and 1461 cm−1. This process is documented in ASTM D7399-18:2018 [18] and the literature [13,15,19]. The known PP content of each virgin blend (X axis) was plotted against the ratio of these bands (1376 cm−1/1461 cm−1) (Y axis), a linear fit applied to the resulting points, the equation noted and used as previously described. The same procedure was completed for the series of virgin PE-HD/PP blends to produce a calibration curve for mixed PO regranulates with the only exception that the absorbance (amplitude) at the bands 720 cm−1 and 1168 cm−1 were used in the ratio 1168 cm−1/(1168 cm−1 + 720 cm−1).

### *2.4. Tensile (Impact) Mechanical Testing of the Virgin, Regranulate and Compatiblised Blends*

Dog bone shaped tensile test specimens were cut from compression moulded sheets (1.8–1.9 mm thick) according to type 5A, ISO 572-2:2012 [20]. Seven replicate tests were performed for each sample type at 23 ◦C and a testing velocity of 10 mm/min using a Zwick 050 universal testing system equipped with a 1 kN load cell and extensometer. Tensile modulus and elongation at break were calculated using the ZwickRoell testXpert III software.

It is important to note that a constant speed of 10 mm/min was used over the entire tensile testing range rather than testing as two distinct segments as suggested in ISO 572-2:2012. The strain rate calculated based on the narrow parallel part of the specimen (0.4 min−1) is subsequently 40 times higher than the strain rate of the special 'modulus segment' described in ISO 572-2:2012 (0.01 min−1).

Tensile impact test specimens were cut from compression moulded sheets (1.1–1.2 mm thick) according to method A, ISO 8256:2004 [21]. Seven replicate tests were performed for each sample type at 23 ◦C using an Instron CEAST 9050 impact pendulum equipped with a 2 J hammer and 15 g crosshead mass. Tensile impact strength (*atN*, kJ/m2) was calculated based on the corrected impact work (*Ec*, J), distance between notches (*x*, mm), and thickness of the narrow parallel test specimen section (*h*, mm) (Equation (1)).

$$a\_{tN} = \frac{E\_c}{\varkappa \cdot h} \cdot 10^3 \tag{1}$$

### *2.5. Thermal Degradation Analysis of the Regranulate*

The thermal degradation properties and inorganic filler content of PE-LD (P05) and mixed PO (PPE, PPP, DPH, DPP, DPS, and DSP) regranulates were assessed using TA Instruments thermogravimetric analysis (TGA) Q500. Regranulate samples of ~10 mg were placed in an alumina crucible and heated from 30 to 600 ◦C at a heating rate of 10 K/min in an air atmosphere.

### *2.6. Morphological and Elemental Analysis of the Regranulate and Compitablised Blends*

Scanning electron microscopy (SEM) imaging and energy dispersive X-ray spectroscopy (EDS) elemental analysis were used to investigate the fracture surfaces of the tensile impact tested PE-LD (P05) and mixed PO (PPE, PPP, DPH, DPP, DPS, and DSP) regranulate specimens and composition of inorganic regranulate residues following TGA, respectively. A Philips XL30 scanning electron microscope was used for the tensile impact tested samples while a ZEISS EVO 10 scanning electron microscope fitted with a ZEISS SmartEDS system was used for the inorganic regranulate residues.

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

### *3.1. Characterisation of Polypropylene Contamination in Post-Consumer Waste*

DSC thermograms of virgin blends of known composition (neat PE-LD, PE-LD with 2, 5, 8, 10, 12, 15, 20, 25 wt.% PP and neat PP) indicated increasing area under the melting peak (melting enthalpy) at 161 ◦C, which is associated with PP, as the PP content increased (Figure 1). P01,03,05 PE-LD regranulates exhibited melting peaks at 109 ◦C and 125 ◦C, indicating that they primarily comprised PE-LD and PE-LLD [22]. An additional melting peak at 161 ◦C suggested a smaller quantity of PP present as contamination. The area under the PP melting peak varied considerably by PE-LD regranulate, ranging from the shortest and narrowest peak associated with P01 to the highest and widest peak for P05. Mixed PO regranulates (PPE, PPP, DHP, DPP, DPS and DSP) exhibited sizable melting peaks between 125–132 ◦C and at 161 ◦C attributable to PE-HD and PP, respectively. PPE exhibited the largest area under the melting peak associated with PE-HD and the smallest associated with PP, while PPP and DPP had the smallest area under the PE-HD melting peak and the largest under the PP peak.

**Figure 1.** Differential scanning calorimetry (DSC) thermograms of (**a**) virgin PE-LD and polypropylene (PP) with (**d)** magnification of PP melting peak (Tm = 161 ◦C) illustrating melt enthalpy differences by blend PP content (2, 5, 8, 10, 12, 15, 20 and 25 wt.%), (**b**) P01,03,05 PE-LD regranulates with (**e**) magnification of PP melting peak (Tm = 161 ◦C) and (**c**) PPE, PPP, DPH, DPP, DPS and DSP mixed PO regranulates with (**f**) magnification of PE-HD (Tm = 125–132 ◦C) and PP (Tm = 161 ◦C) melting peaks.

FT-IR spectra of virgin blends of known composition indicated increasing band intensity at 1376 cm−1, which is associated with −CH3 plane bending, with increasing PP content (Figure 2). The band intensity at 1461 cm−1 associated with −CH2 plane bending simultaneously decreased. P01,03,05 PE-LD regranulates exhibited the same 1376 cm−1 band in addition to a light shoulder at 3200–3500 cm−1 associated with -OH and bands at 1565–1600 cm−1 resulting from–NH stretching. These bands could indicate traces of polyamide, polyester, or low molecular weight contaminants [23]. Mixed PO regranulates exhibited bands at 1168 cm−1 attributable to −CH3 wagging in PP and 720 cm−1 resulting from–CH2−rocking in PE-HD [24].

**Figure 2.** Fourier transform infrared (FT-IR) spectra of (**a**) virgin PE-LD and polypropylene (PP) with (**d**) magnification illustrating differences in the 1376 cm−1 (−CH3 plane bending) and 1461 cm−1 (−CH2 plane bending) bands by blend PP content (0, 2, 5, 8, 10, 12, 15, 20, 25 and 100 wt.%), (**b**) P01,03,05 PE-LD regranulates with (**e**) magnification of 1376 cm−1 (−CH3 plane bending) and 1461 cm−1 (−CH2 plane bending) bands and (**c**) PPE, PPP, DPH, DPP, DPS and DSP mixed PO regranulates with (**f**) magnification of 1168 cm−1 (−CH3 wagging) and 720 cm−1 (−CH2−rocking) bands.

Melting enthalpy-based calibration curves were almost identical for all virgin homopolymer PPs (HD601CF, HC600CF, and HA104E) used to simulate contaminants in PE-LD (Figure 3, Table 2). Block copolymer PP BA202E also exhibited a similar curve. However, random copolymer PP RD 208CF had a radically different gradient to the other curves. This discrepancy can be explained by the differing degree of crystallinity between the homopolymer (~43–45%), block–(~37), and random copolymer PPs (~34%), which affects the melting enthalpy. This makes quantification of the PP content in P01,03,05 PE-LD regranulates, potentially contaminated with any type of PP, challenging using the melting enthalpy. Estimated PP content in P01 PE-LD regranulate ranged from 3.4 to 8.9 wt.% using this method based on the different PP types.

FT-IR-based values were generally slightly higher than those based on the melting enthalpy, a phenomenon also noted in other studies [15] and most likely due to the migration of PP to the material surface during compression moulding. Homopolymer (HD601CF), block–(PP BA202E), and random copolymer (PP RD 208CF) PP calibration curves were much better aligned when using the FT-IR band ratio 1376 cm−1/1461 cm−1. The estimated PP content in PE-LD regranulate ranged from 3.7 to 5.3 wt.% across the different PP types. This makes FT-IR more suitable for estimating the PP content in P01,03,05 PE-LD regranulates. The similarity between the FT-IR band ratio-based calibration curves also allows the provision of a generic equation for calculating blend PP content independent of PP type (Equation (2)), which is useful since the type of PP contamination in regranulates is often unknown. This method is also advantageous as it is faster than DSC-based melting enthalpy experiments and is non-destructive.

$$y = 0.014x + 0.098, \text{ R}^2 = 0.91\tag{2}$$

**Figure 3.** Calibration curves constructed based on (**a**) differential scanning calorimetry (DSC) melting enthalpy and (**c**) Fourier transform infrared (FT-IR) band ratios (1376 cm−1/1461 cm−1) of virgin PE-LD blends of known composition. DSC melting enthalpy-based curves vary considerably by polypropylene (PP) type (especially BA202E block- and RD208CF random copolymers) and crystallinity, while FT-IR band ratio-based curves are less affected. Calibration curves for determining PP content in (**b**) PE-LD and (**d**) mixed PO regranulates are constructed based on DSC melt enthalpy and FT-IR band ratios 1376 cm−1/1461 cm−1 for PE-LD regranulates and 1168 cm−1/(1168 cm−1 + 720 cm−1) for mixed PO regranulates.

**Table 2.** Melt mass-flow rate (MFR, g/10 min) for 2.16 kg at 230 ◦C, degree of crystallinity (%) and calibration curve equations for virgin homopolymer, block- and random copolymer polypropylenes (PP) and calculated PP content for P01 PE-LD regranulate by PP type based on the melting enthalpy and FT-IR band ratio methods.

Generic PP calibration curves for P01,03,05 PE-LD regranulates based on both melting enthalpy and FT-IR band ratios are provided in Figure 3. Similar curves are provided for mixed PO regranulates based on the melting enthalpy and FT-IR bands 720 cm−1 and 1168 cm−1, expressed as 1168 cm−1/(1168 cm−1 + 720 cm−1), of a series of virgin PE-HD (melting enthalpy) and PP (FT-IR band ratio) blends of known composition. These calibration curves were used to calculate the PP content in P01,03,05 PE-LD regranulates and PPE, PPP, DPH, DPP, DPS, and DSP mixed PO regranulates (Table 3).

**Table 3.** Melt mass-flow rate (MFR, g/10 min) for 2.16 kg at a 190 ◦C and b 230 ◦C, and calculated polypropylene (PP) content of PE-LD and mixed PO regranulates based on the melting enthalpy and FT-IR band ratio calibration curves.

Discrepancies between PP content calculated based on melting enthalpy and FT-IR band ratios of PE-LD regranulates were <2 wt.%. PP content varied by up to ~2 wt.% between batches of the same PE-LD regranulate. P01 PE-LD regranulate contained 2.6–3.0% and 2.7–4.6 wt.% PP based on melting enthalpy and FT-IR band ratios, respectively. P03 regranulate contained more PP (5.6–5.9 wt.% and 6.5–7.2 wt.% based on melting enthalpy and FT-IR band ratios, respectively) than P01 but P05 regranulate clearly had the highest PP content with a calculated value of 4.6–6.1 wt.% based on melting enthalpy and 6.9–7.4 wt.% based on the more accurate FT-IR band ratios. P05 PE-LD regranulate was subsequently selected for further mechanical tests.

Discrepancies between calculated PP content based on melting enthalpy and FT-IR band ratios were much higher for mixed PO regranulates than PE-LD regranulates, ranging up to ~24 wt.% for DPH (70.1 wt.% calculated based on melting enthalpy compared to 46.1 wt.% based on FT-IR band ratios). These inconsistencies are attributable to DSC melting curve overlap resulting from the presence of PP block or random copolymer contaminants, which contain an ethylene fraction represented as a low temperature shoulder overlapping the PE peak [16]. PP content calculations based on FT-IR band ratios were accepted as more accurate than values based on melting enthalpy and are subsequently reported here. DPP and PPP mixed PO regranulates comprised almost entirely PP (94.9 wt.% and 87.5 wt.%, respectively), while DSP comprised 60 wt.% PP and DPH and DPS were a little less than half PP (46.1 wt.% and 44.5 wt.%, respectively). PPE contained just 5.5 wt.% PP.

### *3.2. Thermal Degradation Properties and Inorganic Content of Post-Consumer Waste*

P05 PE-LD and PPE, PPP, DPH, DPP, DPS, and DSP mixed PO regranulates exhibited thermal degradation properties typical of PP and PE, indicating negligible aging or well-stabilized material (Supplementary Figure S1). Onset of thermal degradation was determined by the PP phase and further degradation behavior by the dominant blend component. PP rich regranulates were hence readily distinguishable from those which contained large quantities of PE. The inorganic residues of mixed PO regranulates were 0.9 wt.% for PPE, 1.8 wt.% for DPS, 2.1 wt.% for DSP, 2.2 wt.% for DPP, 2.3 wt.% for PPP, and 2.4 wt.% for DPH. P05 PE-LD regranulate had an inorganic residue of 2.8 wt.%. Inorganic residues mainly consisted of Si, Ca, Ti, Mg, Fe, S, Na, and Cl.

### *3.3. Effect of Polypropylene Contamination on Tensile and Tensile Impact Properties*

P05 PE-LD regranulate had a higher tensile modulus (*Et*) than virgin PE-LD (453 MPa compared to 320 MPa), a considerably lower elongation at break (εb) (421% compared to 620%) and tensile impact strength (*atN*) (84 kJ/m2 compared to 115 kJ/m2). These undesirable properties result from the ~7 wt.% PP contamination present in this PE-LD blend, which has a much higher *Et* than PE-LD and hinders miscibility and adhesion between blend components (Figure 4).

**Figure 4.** Effect of polypropylene (PP) content (hollow markers) and melt mass-flow rate (MFR, solid markers) on the (**a**) tensile modulus (*Et*), (**b**) elongation at break (*εf* ) and (**c**) tensile impact strength (*atN* ) of PE-LD and mixed PO regranulates. Effect of 5 wt.% ethylene-based olefin block copolymer on the (**d**) *Et*, (**e**) *εf* and (**f**) *atN* of PE-LD and mixed PO regranulates (solid markers prior to compatibilisation and hollow markers after).

Compatibilisation with 5 wt.% ethylene-based olefin block copolymer enhanced interfacial adhesion and provided an 81% improvement in the εb of P05 (421% to 762%), a value 23% higher than even virgin PE-LD (762% compared with 620%). The compatibiliser also improved stress transfer between the phases resulting in a 39% increase in the *atN* of P05 (84 kJ/m2 to 118 kJ/m2) and a comparable value to that of virgin PE-LD (115 kJ/m2). SEM micrographs indicated that the already present necking in P05 was increased postcompatibilisation with shear yielding the primary deformation mechanism in both samples (Figure 5). The *Et* of P05 simultaneously decreased with compatibilisation to 381 MPa, a 16% reduction which resulted in an *Et* just 19% higher than virgin PE-LD. This is due to the low *Et* of ethylene-based olefin block copolymer, which encapsulates the dispersed phase in the matrix and reduces the *Et* of the blends. Significantly, these results indicated that compatibilised P05 PE-LD regranulate exhibits competitive or better tensile properties than even virgin PE-LD films and can be considered a viable recycled substitute.

All mixed PO regranulates other than PPE exhibited high (≥~45 wt.%) PP contents and high melt mass-flow rates (MFR) (2.5−20.0 g/10 min) better suited to injection than compression moulding, which resulted in brittle tensile properties. Literature does, however, indicate similar properties in injection moulded samples suggesting a more general, moulding method independent sensitivity of εb to the presence of mixtures of polymers with different molar mass distributions in recycled materials [25]. Blend *Et* increased with MFR and PP content, while εb decreased violently at elevated MFRs and PP contents, indicating the sensitivity of this parameter to these factors. *atN* was also sensitive to increasing MFR and PP content, dropping very quickly as PP content increased, but then recovering slightly at MFR ≥ 5 g/10 min and PP contents ≥ 60 wt.%.