Correlations between blowability and uniaxial
stretchability of PET before and after mechanical
recycling

Nathan Sylvestre, Jean-Luc Bouvard, Guillaume Corvec, Mikaël Derrien, Xavier Monnier, Christelle Combeaud

► To cite this version:

Nathan Sylvestre, Jean-Luc Bouvard, Guillaume Corvec, Mikaël Derrien, Xavier Monnier, et al. Correlations between blowability and uniaxial stretchability of PET before and after mechanical recycling. Polymer Testing, 2025, 146, pp.108774. 10.1016/j.polymertesting.2025.108774. hal-05441976

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# Correlations between blowability and uniaxial stretchability of PET before and after mechanical recycling

Nathan Sylvestre<sup>a, b</sup>, Jean-Luc Bouvard<sup>a</sup>, Guillaume Corvec<sup>a</sup>, Mikaël Derrien<sup>b</sup>, Xavier Monnier<sup>b</sup>, Christelle Combeaud<sup>a, * b</sup>

## ARTICLE INFO

**Keywords:**

Mechanical recycling

PET stretching

Free stretch-blowing tests

Time/temperature equivalence

## ABSTRACT

In the literature, previous studies have linked the behaviour of PET under stretch-blowing and under uniaxial stretching thanks to the time-temperature equivalence principle. Other works showed variations of micro-structural, thermal and mechanical properties of PET submitted to mechanical recycling, but blowability was not measured. In this article, the behaviour during free stretch-blowing tests was investigated for a virgin commercial PET before and after mechanical recycling with a step of solid-state post-condensation. For the same physical state, the lower stretch-blown volumes obtained for the recycled PET were correlated with a sooner nucleation leading to a lower Natural Draw Ratio, as observed during uniaxial tensile tests. Our results confirm that the uniaxial stretching tests used in the past to discriminate the behaviours of neat PET resins in blowing, can also be used to predict variations of blowability under mechanical recycling.

## 1. Introduction

During deformation in its rubbery state, the poly(ethylene terephthalate) (PET) can develop local macromolecular orientation and organization, leading to a strain induced crystallization (SIC) and to a structural hardening responsible for good mechanical and barrier properties, which explain the wide use of this material in bottles for beverages [1–3].

It is well known that the Natural Draw Ratio, at which the strain hardening occurs, is function of the (temperature; strain rate) couple which drives the physical state and the molecular mobility of the polymer *via* the gap to its α relaxation [4–6], as described by the time-temperature equivalence principle [7].

In particular, the behaviour of PET during stretch-blowing tests has been very well described by the team of Menary, Nixon and Yan [8–10], and by our research group in Cemef [11–15], which showed that kinematics of blowing of PET is driven by its local structural hardening depending on its physical state.

However, in the works of Deloye et al. [16], it was shown that several copolyesters can exhibit different Natural Draw Ratios for a same physical state, which can lead to variations of the mechanical behaviours of the resins during stretch-blow moulding process.

More recently, the impact of molecular structure on the stretchability of polyesters was described by several papers of Forestier et al. [17–19] comparing the uniaxial stretchability of poly(ethylene furanoate) (PEF) to PET. To be in equivalent physical states, the (temperature; strain rate) couples had to be adjusted to the α relaxation temperatures $T_α$ of each resin. Even in equivalent physical states, the mechanical behaviours of the two polyesters were different, due to the impact of their molecular architecture on the SIC phenomena.

To sum up, if a manufacturer of plastic bottles wishes to get the same final microstructures and properties between two polyesters, it should adjust the blowing parameters according not only to the α relaxation temperatures of each resin, but also according to the variations of behaviours resulting from material variability between the two grades.

The Single-Use Plastics Directives (SUP) [20] sets the minimum incorporation rate of recycled PET in PET bottles at 25 % in 2025, leading to an increasing use of recycled polyesters. In that context, the measurement of the variability of the incoming material becomes a major concern of the manufacturers to adjust their processes to the recycled polyesters. Indeed, the recycled PET are much more variable and unpredictable than a synthetic virgin PET, due to variability of deposit and contaminants [21–23] but also because of the numerous scissions and coupling reactions occurring during PET (re-)processing

<footnote><sup>a</sup> MinesParis, Université PSL, Centre de Mise en Forme des Matériaux (CEMEF), CNRS, UMR 7635, Sophia, Antipolis, France</footnote>

<footnote><sup>b</sup> Sidel Blowing & Services SAS, Avenue de la Patrouille de France, CS 60627 Octeville-sur-Mer, 76059, Le Havre, France</footnote>

<footnote>* Corresponding author.</footnote>

<footnote>E-mail address: christelle.combeaud@minesparis.psl.eu (C. Combeaud).</footnote>
[24,25].

For that reason, protocols to measure variability of behaviour of PET during stretch-blow moulding has to be developed and tested.

In this paper, the method already used in previous studies to characterize virgin polyesters in Cemef [11–15], was applied on a virgin commercial PET before and after a mechanical recycling in conditions representative of industrial recycling. Such approach was similar in some respect to the method of the team of Menary, Nixon and Yan [8–10], and consisting to extrapolate uniaxial tensile tests to stretch-blowing tests thanks to the time-temperature equivalence principle.

The stretchability of the two resins were first characterized in uniaxial conditions for different equivalent strain rates in their rubbery state. The (temperature; strain rate) couples were chosen thanks to the time-temperature equivalence determined by dynamic mechanical analysis (DMA) and extrapolated to high strains. Some microstructural assumptions on the stretchability variations upon recycling were supported by calorimetric analysis.

Then, the equivalent strain rates of PET during the free stretch-blowing tests were deduced from local thermic and kinetic measurements, and compared to the ones get in uniaxial drawing.

Finally, the protocol used in previous studies [11–15] to classify the blowability of several virgin polyesters from uniaxial laboratory tests, thanks to the time-temperature equivalence, was applied here to our virgin and recycled resins. It allowed to evaluate the effects of mechanical recycling on the blowability but also to check if this protocol, valid for virgin polyesters, can be validated for recycled materials.

## 2. Experimental

### 2.1. Materials and processing

The PET resin used in the study is a virgin commercial grade for bottle application, in granulated form, named “V0” in this article. V0 is a copolymer made from terephthalic acid (PTA), isophthalic acid (IPA) and mono-ethylene glycol (EG). Thus it is a poly(ethylene terephthalate – co-ethylene isophthalate). The measurement of comonomers contents were out of the scope of this study, since the goal was to propose a method for prediction of blowability from uniaxial drawing tests.

The recycled PET was obtained from V0 within three stages: injection of preform, blowing of bottles on an industrial line and mechanical recycling. Mechanical recycling of bottles into granulates was done on the pilot line of Sidel as follows: crushing of bottles into flakes, washing with soda, rinsing and separation in water, drying, extrusion-granulation, and solid-state post-condensation (SSP). Recycling parameters are described in Table 1. The recycled PET grade obtained by this way is called “V1” in this article.

<table><caption>**Table 1**

Mechanical recycling parameters used in this study.</caption><thead><tr><th>Step</th><th>Description and parameters setting</th></tr></thead><tbody><tr><td>Washing</td><td>8 min at 85 °C, 0.6 to 2 % soda (12&lt;pH&lt;14), 0.2 % detergent.</td></tr><tr><td>Rinsing and separation</td><td>6&lt;pH&lt;8<br>2-4 g/L wetting agent</td></tr><tr><td>Post-rinse drying</td><td>10 min drying with air circulation to avoid agglomeration Temperatures: 80/90/100/110/ 120/130/140°C.</td></tr><tr><td>Pre-extrusion drying</td><td>140°C,300 min, final water content &lt;50 ppm.</td></tr><tr><td rowspan="3">Extrusion-granulation</td><td>Heating profile<br>70/220/230/240/250/ 260/260°C</td></tr><tr><td>Screw rotation<br>130 rpm</td></tr><tr><td>Vacuum filtration and degassing<br>6h at 140 °C</td></tr><tr><td rowspan="3">Solid State Polymerization (SSP) under vacuum</td><td>Crystallization<br>step<br>700 min (11h40) at 210°C</td></tr><tr><td>Polymerization<br>step</td></tr><tr><td></td></tr></tbody></table>

The granulates of V1 were injected into preforms. Injection parameters were the same for V0 and V1. Then, the preforms were sent on the industrial blowing line of Sidel, where an operator adjusted the parameters according to the state of the art in order to get bottles without defects such as opalescence, cracks or dimensional instability. While V0 was blown at around 106 °C, the same temperature led to a whitening under the neck of the bottles V1, as observed by the operator during industrial tests. This whitening could be due to the thermal crystallization of the preform during heating and/or to the damaging of the bottles during blowing, with spherulites and/or microcavities diffusing the light. By setting the blowing temperature of V1 to 115 °C, the whitening was avoided. This behaviour under industrial conditions is consistent with the results of the free stretch-blowing tests, as explained later in the article. It is known that recycled PET usually show higher absorption coefficient in near infra-red (IR), in the range used for reheating preform before stretch-blow moulding [26]. However, a spectroscopic study done by our industrial partner Sidel showed that IR absorptions of materials V0 and V1 are equivalent. In our case, the differences between blowing temperatures do not come from IR absorption but from the necessity to adjust the processing parameters to avoid the whitening of V1 bottles as explained above.

As shown in Tables 2 and Table 3, intrinsic viscosities, molecular weights and polydispersity of preforms V0 and V1 are equivalent, despite a lower intrinsic viscosity measured on V1 granulates before SSP. Thus, as observed in the literature [27,28], the solid-state post-condensation allows to repolymerize the chains that were cut (chain scission) during extrusion, thanks to a good setting of recycling parameters, especially of SSP parameters.

For uniaxial stretching tests, V0 and V1 were injected into square plates 3 mm thick and 60 mm square. The plates were core surfaced to a final thickness of 1 mm to avoid microstructural gradient. The “hour-glass” stretching specimens were cut from the surfaced plates according to the geometry shown in Fig. 1a. The geometry of the preforms used for stretch-blowing tests is described in Fig. 1b. Stretching specimens and preforms were stored in the open air, which corresponds to the usual conditioning applied by preform blow moulders before blowing.

### 2.2. Methods

#### 2.2.1. Differential scanning calorimetry (DSC)

Differential Scanning Calorimetry (DSC) analyses were carried out with a PerkinElmer 4000 DSC on samples between 2 and 4 mg, taken with a cutter from the entire thickness of the surfaced plates and placed in a sealed aluminium capsule. Three samples were analysed for each grade to check repeatability. The DSC thermal cycle, described in Fig. 2, consisted of one heating stage followed by one cooling step at 10 °C/min between 30 and 280 °C, with a hold at 280 °C for 1 min.

The goal of DSC analysis was double.

Firstly, this technique was used to evaluate the ability of V0 and V1 to thermal crystallization. The first heating does not give access to intrinsic behaviour of the material, because it is related to the thermomechanical history encountered by the polymer during processing. A second heating could also be problematic because it is not guaranteed that both resins V0 and V1 are in the same state of organization and crystallinity at the beginning of the second heating (i.e. at the end of the cooling). Thus, it was chosen to compare the ability of the two resins to the crystallization from their molten state which can considered

<table><caption>**Table 2**

Intrinsic viscosity of V0 and V1 at different stages of processing.</caption><thead><tr><th rowspan="2"></th><th colspan="3">Intrinsic viscosity<sup>a</sup> (dL/g) ±0.02</th></tr><tr><th>Granulates before SSP</th><th>Granulates after SSP</th><th>Injected preforms</th></tr></thead><tbody><tr><td>V0</td><td>–</td><td>0.80</td><td>0.78</td></tr><tr><td>V1</td><td>0.69</td><td>0.83</td><td>0.77</td></tr></tbody></table>

<footnote>a Internal measure, Sidel.</footnote>
<table><caption>**Table 3**
Molecular weights and polydispersity of the preforms V0 and V1. Mean values and standard errors for three measurements per grade.</caption><thead><tr><th rowspan="2"></th><th colspan="3">Molecular weights and polydispersity of the preforms<sup>a</sup></th></tr><tr><th>Mn (kg/mol)</th><th>Mw (kg/mol)</th><th>Mw/Mn</th></tr></thead><tbody><tr><td>V0</td><td>43.7 ± 0.5</td><td>86 ± 2</td><td>1.97 ± 0.1</td></tr><tr><td>V1</td><td>42 ± 1</td><td>84 ± 3</td><td>1.99 ± 0.1</td></tr></tbody></table>

<sup>a</sup> Gel permeation chromatography in PMMA equivalent.

amorphous. More precisely, the comparison was made between the peaks of crystallization from the melt of V0 and V1. Results are presented in section 3.1.

The second target of DSC analysis was motivated by a study of Girard et al. [29] showing that a pre-organization developed during a thermal annealing before stretching promotes the strain-induced crystallization. Thus, it is important to check if even before the stretching, the materials show an organization following the injection process.

Consequently, the crystallinity ratio $\chi_c$ of injected parts (plates) is calculated upon the first heating according to the relationship

$$ \chi_c = \frac{\Delta H_m - \Delta H_{cc}}{\Delta H_m^0} \qquad (1) $$

with $\Delta H_m$ and $\Delta H_{cc}$ enthalpies of melting and cold crystallization (J/g) respectively, and $\Delta H_m^0$ the theoretical melting enthalpy of 100 % crystalline PET ($\Delta H_m^0 = 140$ J/g) [30]. The crystallinity ratio was less than 7 % for all samples before drawing, validating the effectiveness of in-mould quenching at 15 °C during plate injection. More precisely, the crystallinity ratios of the injected parts obtained from the first heating were equal to 3.9 %, 2.1 % and 0.4 % for the three samples V0, and to 1.6 %, 4.6 % and 6.7 % for the three samples V1. Considering the uncertainty of the measurement and the variability of the results, it cannot be claimed that there is a significant gap between the crystallinity ratios of V0 and V1 following the injection process.

### 2.2.2. Viscoelastic measurements

Dynamic mechanical analyses (DMA) were performed in tension mode using a Mettler Toledo DMA, by applying a dynamic deformation of 0.04 % to stay in the linear viscoelastic range. For each test, specimens of 5 x 3 × 1 mm³ were submitted to a preload of 1 N to maintain tension.

Temperature scans consisted of a ramp from 40 to 130 °C at 2 °C/min. Frequency scans were achieved from 0.1 to 15 Hz, for isotherms from 75 to 105 °C by increments of 5 °C. Three samples were analysed for each grade.

### 2.2.3. Uniaxial drawing experiments

The mechanical behaviour was measured on two samples for each grade on a stretching prototype described previously [17,31]. Samples were first heated by convection, then stretched with an exponential traverse speed to get a constant true strain rate (equation (2)).

$$ V(t) = L_0 \times \dot{\epsilon} \times \exp(\dot{\epsilon} \times t) \qquad (2) $$

where V is the arm speed (mm/s), $L_0$ the initial length between the clamps (mm) and $\dot{\epsilon}$ the true strain rate (s⁻¹). Local strains and strain rates at the centre of the processing zone were obtained from image

<figure><img src="image_3.png" /><figcaption>Fig. 2. Thermal cycle applied during differential scanning calorimetry.</figcaption></figure>

<figure><img src="image_4.png" /><figcaption>Fig. 1. Geometry and dimensions (mm) of a) “hourglass” uniaxial drawing specimens (thickness = 1 mm) and b) injected preforms.</figcaption></figure>
correlation using VIC2D® software from Correlated Solutions.

During stretching, molecular mobility and thermomechanical behaviour are driven by the physical state of the material through the gap to its α relaxation. In previous studies, a method based on time-temperature superposition principle to control this gap during stretching has been validated on different polymers: PMMA [32], PET [11–13] or PEF [19]. This method has been used in this context and is described in the Supporting Information.

Finally, stretching conditions were defined according to time-temperature equivalence principle in order to explore the forming range of PET, i.e. for equivalent strain rates from $10^{-2}$ to $10^{-4}$ s$^{-1}$, as shown in Table 4. This table shows that stretching conditions are relatively similar and roughly equivalent between V0 and V1, as α transition temperature of both materials are close.

### 2.2.4. Free stretch-blowing tests

#### 2.2.4.1. Free stretch-blowing prototype.
During the forming of a bottle by stretch-blow moulding of an amorphous preform heated to its rubbery-like state, at a pressure of less than 15 bar during pre-blowing and then 20–35 bar, the pre-blowing step is essential. It ensures most of the axial and orthod radial stretching of the polymer as well as the development of the microstructure and the final properties induced by stretching.

The free (without mould) stretch-blowing prototype from Sidel allows to reproduce this pre-blowing step, during which most of the volume of the bubble and its microstructure are developed, in order to measure the intrinsic mechanical response of the material not constrained by the mould.

The blowability of preforms V0 and V1 was measured on this prototype, as described in the Supplementary Information.

Temperature of the external surface was measured by an infra-red pyrometer at the exit of the oven, 20 mm below the neck of the preform.

To explore all the forming range of PET, three temperatures were investigated, and three preforms (see geometry in Fig. 1b) were tested per temperature and per grade. As shown in Table 5, the good control of heating parameters (i.e. the setting of IR lamps power) allowed to blow the preforms V0 and V1 at the same surface temperatures T1, T2 and T3.

One should notice that the behaviour of the preform is driven by the temperature during stretch-blowing, which is different from the temperature just after heating. A thermographic study, not presented here, showed that the temperature of the preform skin decreases by 10 °C during the transfer from the oven to the blowing module, whatever the resin (V0 or V1) or the temperature (T1, T2 or T3). In the following sections, the temperature of blowing refers to the temperature at the beginning of the stretch-blowing test, as indicated in Table 5.

#### 2.2.4.2. Measurement of free stretch-blowing kinematics.
Before each test, the preforms were marked with white dots (speckle) using a pen. During the stretch-blowing test, images of the preform were saved by a camera

<caption>**Table 4**</caption>

Drawing conditions actually obtained for each grade as a function of the target equivalent strain rate. The experimental values for temperature, true strain rate and equivalent strain rate at a reference temperature $T_α$ are average values before hardening for a true strain between 0.2 and 0.8.

<table><caption>**Table 5**</caption><thead><tr><th>Target ᵈ<sub>eq</sub> (s<sup>-1</sup>)</th><th>Experimental conditions</th><th>V0</th><th>V1</th></tr></thead><tbody><tr><td rowspan="3"><b>10<sup>-2</sup></b></td><td>T (°C)</td><td>86</td><td>84</td></tr><tr><td>芯 (s<sup>-1</sup>)</td><td>0.079</td><td>0.088</td></tr><tr><td>芯<sub>eq</sub> (s<sup>-1</sup>)</td><td>2x10<sup>-2</sup></td><td>5x10<sup>-2</sup></td></tr><tr><td rowspan="3"><b>10<sup>-3</sup></b></td><td>T (°C)</td><td>99</td><td>93</td></tr><tr><td>芯 (s<sup>-1</sup>)</td><td>0.38</td><td>0.11</td></tr><tr><td>芯<sub>eq</sub> (s<sup>-1</sup>)</td><td>1x10<sup>-3</sup></td><td>2x10<sup>-3</sup></td></tr><tr><td rowspan="3"><b>10<sup>-4</sup></b></td><td>T (°C)</td><td>100</td><td>101</td></tr><tr><td>芯 (s<sup>-1</sup>)</td><td>0.078</td><td>0.088</td></tr><tr><td>芯<sub>eq</sub> (s<sup>-1</sup>)</td><td>2x10<sup>-4</sup></td><td>2x10<sup>-4</sup></td></tr><tr><td>Reference temperature T<sub>α</sub> (°C)</td><td></td><td><b>82</b></td><td><b>83</b></td></tr></tbody></table>

Temperatures measured by an infra-red pyrometer at the exit of the oven, 20 mm below the neck of the preforms V0 and V1. Mean values on three preforms per grade and temperature.

<table><thead><tr><th rowspan="2"></th><th colspan="3">Temperature after heating (°C)</th><th colspan="3">Temperature at the beginning of the stretch-blowing test (°C)</th></tr><tr><th>T1</th><th>T2</th><th>T3</th><th>T1</th><th>T2</th><th>T3</th></tr></thead><tbody><tr><td>V0</td><td>107</td><td>120</td><td>128</td><td>97</td><td>110</td><td>118</td></tr><tr><td>V1</td><td>108</td><td>121</td><td>128</td><td>98</td><td>111</td><td>118</td></tr></tbody></table>

Phantom®, every 2 ms (500 frames per second). The images of preforms with white speckle were analysed to follow global and local deformations occurring during the test, by using two protocols.

The first protocol, described in the Supplementary Information, gives the evolution of the volume during the free stretch-blowing test.

The second method of image analysis, also explained in the Supplementary Information, allows to follow the evolution of longitudinal and radial strains in the centre of the preform.

The estimate of the maximum absolute error in the measurement of deformation due to out-of-plane movements is also described in the Supplementary Information.

## 3. Results and discussion

### 3.1. Variations of intrinsic properties with mechanical recycling

A previous study, which is the subject of another article [33], showed that two intrinsic parameters were significantly affected by the mechanical recycling of V0: the crystallization ability in the melt state, and the structural hardening under uniaxial stretching in the rubbery-like state.

In Fig. 3, the thermograms of V0 and V1 under cooling from the melt state show a sooner nucleation for V1 as well as a higher kinetics of anisothermal crystallization.

As discussed in section 2.1, the intrinsic viscosities, molecular weights and polydispersity are equivalent between the two grades and cannot explain their different behaviours upon cooling. Some hypotheses will be discussed later in the paper.

Fig. 4 shows the true stress-true strain curves of V0 and V1 for three equivalent strain rates ᵈ<sub>eq</sub> at their reference temperature $T_α$. The strain

<figure><img src="image_5.png" /><figcaption>**Fig. 3.** DSC thermograms of PETs V0 (blue) and V1 (green) during cooling at 10 °C/min from the molten state. The curves are shifted vertically to make the graph easier to read.</figcaption></figure>
on hardening $\varepsilon_{hardening}$, defined at the intersection of the tangents before and after hardening, was expressed in a more common quantity, the Natural Draw Ratio (NDR) $\lambda_{NDR} = \exp(\varepsilon_{hardening})$, as function of $\dot{\varepsilon}_{eq}$, as shown in Fig. 5.

For high equivalent strain rates (about $10^{-2}$ s$^{-1}$), near from the viscoelastic region of the mastercurves (Fig. 2 in the Supplementary Information), mechanical behaviours of V0 and V1 are close. Near the viscoelastic region, stretching is fast enough to extend and orientate the molecular chains segments efficiently without letting them the time to relax [4,5]. This efficient molecular extension and orientation make easier the nucleation of smectic and nematic mesophases, in which intermolecular interactions lead to a structural hardening of the material [1,17,35–37]. So, nucleation of mesophases and strain hardening are promoted for both materials V0 and V1, and no significant differences are observed.

For lower equivalent strain rates ($10^{-3}$ and $10^{-4}$ s$^{-1}$), further on the rubbery plateau, V1 is initially softer than V0, then V1 presents a sooner hardening (lower NDR) than V0, especially at $10^{-4}$ s$^{-1}$.

According to Figs. 3–5, both thermal and dynamic crystallizations of V1 are promoted compared to V0, as indicated by anisothermal crystallization peak on DSC curves upon cooling and by lower natural draw ratio during stretching.

As a reminder, our study if focused on the prediction of blowability of rPET from uniaxial drawing tests. The microstructural origin of the thermal and dynamic crystallizations promoted for V1 is out of this scope and were not investigated in this work. However, three main hypotheses could be emitted on the base of some articles of the literature.

Firstly, Torres et al. [23] showed a correlation between the rPET contamination and its thermal crystallization. In our case, even if V1 arises from a “clean” recycling closed loop, it cannot be excluded that some impurities were present in the recycling line, in the extruder and/or in the injection machine. Those impurities could act as nucleating agents leading to an earlier crystallization from the melt state [23].

Secondly, Yoo et al. [38] showed that the thermal degradation kinetics at 300 °C of PET increases with the isophthalate (IPA) comonomers content. According to Holland and Hay [39], this higher sensitivity of IPA units to degradation comes from their kinked shape promoting the intramolecular degradation reactions in comparison to linear terephthalate units. According to these studies of the literature, it could be possible that IPA units initially in V0 were removed from molecular chains during mechanical recycling. Besides, as shown by Wu et al. [34], the lower symmetry and regularity of the comonomers chains containing IPA segments leads to a lower kinetics of crystallization in comparison to PET homopolymer. Thus, a hypothesis to explain the easier and faster thermal crystallization of V1 could be a chemical structure closer to homopolymer PET with a lower content of IPA units compared to V0.

A lower IPA content in V1 could also explain its lower modulus and its promoted structural hardening during uniaxial drawing tests.

<figure><img src="image_6.png" /><figcaption>Fig. 5. Uniaxial Natural Draw Ratio (NDR) as a function of equivalent strain rate at the reference temperature $T_α$ for grades V0 (blue) and V1 (green).</figcaption></figure>

Indeed, according to atomistic simulations of Zekriardehani et al. [3], due their kinked configuration, the phenyl rings of poly(ethylene isophtalate) (PEI) show significantly lower relaxation kinetics than the phenyl rings in poly(ethylene terephthalate) (PET), and this difference between PEI and PET increases with temperature. Moreover, as explained by Forestier et al. [19], the increase of the stress at the beginning of the true stress - true strain curve in the rubbery state is due to the extension of the chains. This extension is hindered by higher relaxation kinetics promoting retraction and reptation of the chains [4, 5]. Thus, if V1 contains less IPA units than V0, it could be expected that V1 shows higher kinetics of relaxation and exhibits lower modulus in the first stages of stretching, as shown in Fig. 4.

Also, the lower NDR of V1 during stretching could be coherent with its higher nucleation ability observed during crystallization in quiescent conditions. Indeed, in her thesis, Forestier showed that the more complex structure of poly(ethylene furanoate) (PEF), compared to PET, leads to both slower and less favourable thermal crystallization and later structural hardening (higher NDR) under stretching due to less favourable dynamic crystallization. In the same manner, if V1 contains less IPA units than V0, it could be assumed that the higher linearity and regularity of V1 promotes its thermal crystallization as well as the nucleation of mesophases during stretching leading to a lower NDR.

There is a third hypothesis to explain the lower NDR observed for V1 at lower equivalent strain rates ($10^{-3}$ and $10^{-4}$ s$^{-1}$). This behaviour could be due to the occurrence of thermal crystallization during the 5 min of heating stage. Indeed, Girard et al. [29] showed that a

<figure><img src="image_7.png" /><figcaption>Fig. 4. True stress-strain curves for V0 (blue) and V1 (green) for three drawing conditions corresponding to three equivalent strain rates (from left to right), at a reference temperature $T_α$.</figcaption></figure>
pre-annealing before stretching leads to a preorganization promoting the strain induced crystallization during drawing. An outlook to our work could be to submit PET samples to a thermal cycle equivalent to the one applied during the stretching tests, and to compare by DSC the crystallinity ratios developed in static and in dynamic conditions. Such a protocol should allow to check if the thermal treatment applied during the test promotes or not the crystallization of the sample.

Here, these results are not treated more in details, since they have already been described in a previous study [33].

## 3.2. Kinematics and kinetics of blowing

### 3.2.1. Effect of temperature and physical state on the virgin PET

Fig. 6 shows the evolution of longitudinal and radial stretching ratios and strain rates during stretch-blowing tests for V0, for three temperatures of blowing.

Independently of the temperature, kinematics of stretch-blowing are reproducible from a test to another. This kinematics has already been largely described by the team of Menary, Nixon, Yan [8–10] and Billon [11–15]. The main steps, visible in Fig. 6, are briefly described below as reminder, even if it is not the goal of this article, since all the details can be found in the other articles of literature [8–15].

1. At the beginning, pressure is too low and preform too thick to allow radial inflation. The stretching rod displacement causes a longitudinal stretching.

2. Then, increasing pressure, helped by thinning of the walls of the preform, causes a radial dilatation, i.e. an orthoradial stretching, leading to a local longitudinal contraction.

3. With the thinning of the preform and the increasing pressure, the bubble initiated under the neck “explodes”, i.e. the bubble is submitted to bi-axial stretching and propagates towards the bottom of the preform. During its propagation on the extensometers in the middle of the preform (40 mm under the neck), a bi-axial stretching is measured, as shown by the simultaneous increase of the stretching ratios (Fig. 6 a and c) and by the true strain rates peaks in both longitudinal and radial directions (Fig. 6 b and d).

4. Due to combined action of stretching rod and air injection, the bubble extension is faster in longitudinal direction, thus the structural hardening first occurs in this direction. On the other hand, radial extension continues, until the structural hardening in this second direction. Then, the bubble propagates to not yet hardened regions, easier to deform.

As blowing temperature increases, the biaxiality starts sooner, the maximums of true strain rate peaks are higher, and the final radial stretching ratio increases, while the final longitudinal extension is not significantly affected. The higher strain rates and radial strains are a consequence of the molecular mobility at higher temperature, making conformational changes easier [4]. In the longitudinal direction, the stretching rod displacement is the same for all the blowing temperatures, so the final stretching ratio in the longitudinal direction is imposed by the stretching rod.

As explained above, molecular mobility and stretchability are affected by the physical state of the polymer, via the gap to its $\alpha$ relaxation. This gap is driven by the (temperature; strain rate) couple and can be easily described by the equivalent strain rate at the reference temperature $T_\alpha$, as done for uniaxial tests in section 3.1. of this article.

<figure><img src="image_8.png" /><figcaption>Fig. 6. Evolution with time of strains (graphs a and c) and strain rates (b and d) measured on VIC2D® by virtual extensometers in the radial (a and b) and longitudinal (c and d) directions in the middle (40 mm under the neck) of the V0 preform, for three initial external temperatures: 97 °C (purple plots), 109–110 °C (blue) and 117–118 °C (red). Rod speed = 400 mm/s, pressure setting = 7 bar.</figcaption></figure>
Using the time-temperature equivalence principle determined for V0 (section 2.3.3), the maximum of each strain rate peak (where most of the deformation occurs) was multiplied by the shift factor deduced from blowing temperature and WLF law. The result was the equivalent strain rate at the reference temperature $T_α$ during biaxial stretching, in longitudinal and radial direction, for each blowing condition, as shown in Fig. 7.

When comparing equivalent strain rates in blowing (from $2 \times 10^{-3}$ to $10^{-1}$ s$^{-1}$, Fig. 7) and in uniaxial stretching (from $2 \times 10^{-4}$ to $4 \times 10^{-2}$ s$^{-1}$, Figs. 4 and 5) to the mastercurve of V0 (Fig. 2 in the Supporting Information), it appears that our stretch-blowing and uniaxial tensile tests are achieved for similar physical states which are positioned on the rubbery plateau of V0. Thus, it is possible to measure the influence of the physical state of the polymer on the preform blowability. In this article, the quantity of interest was the final volume developed by stretch-blowing. This final volume is expected to be directly linked to the stretchability and to the strain hardening of the material, with a lower volume obtained for a sooner structural hardening. This quantity gives information about the intrinsic behaviour of the PET, and it is interesting to see the impact of the physical state on this behaviour.

Fig. 8 shows the final volume, measured by the Python® program of contour detection, as a function of radial equivalent strain rate, in the middle of the preform V0. As expected, higher blowing temperature, or lower equivalent strain rate, leads to higher final volume in free stretch-blowing. This result is coherent with our uniaxial stretching tests, where the uniaxial Natural Draw Ratio (NDR) increases when the polymer is deformed further on its rubbery plateau.

### 3.2.2. Variation of blowability under mechanical recycling and correlation with uniaxial stretchability

The previous section showed that it is possible to evaluate blowability of a preform with its final volume, and to correlate this behaviour to the physical state of the polymer thanks to the time-temperature equivalence. The same procedure is used to compare the blowabilities of V0 and V1 according to their physical state, for three blowing temperatures, as shown in Fig. 9.

For high equivalent strain rates, at the beginning of the rubbery plateau, the virgin and recycled preforms develop identical final volumes. When the equivalent strain rate decreases, the final volume increases, due to higher molecular mobility and stretchability. This result is consistent with higher Natural Draw Ratios (NDR) measured in

<figure><img src="image_9.png" /><figcaption>Fig. 7. Longitudinal (circles with vertical line) and radial (circles with horizontal line) equivalent strain rates in the middle of the preform, deduced from the blowing temperatures and from the maximums of true stain rate peaks, as a function of blowing temperature. Mean values for three tests.</figcaption></figure>

<figure><img src="image_10.png" /><figcaption>Fig. 8. Final volume of V0 preform developed by stretch-blowing, measured by the Python® program of contour detection, as function of radial equivalent strain rate deduced from the blowing temperature and from the maximum of the true radial strain rate peak, in the middle of the preform. Mean values for three tests. Blowing temperature is noted next to each point.</figcaption></figure>

<figure><img src="image_11.png" /><figcaption>Fig. 9. Final volume of V0 (blue) and V1 (green) preforms developed by stretch-blowing, measured by the Python® program of contour detection, as function of radial equivalent strain rate deduced from the blowing temperature and from the maximum of the true radial strain rate peak, in the middle of the preform. Mean values for three tests. Blowing temperature is noted next to each point.</figcaption></figure>

uniaxial stretching for lower equivalent strain rates.

Moreover, for equivalent strain rates below $10^{-2}$ s$^{-1}$, V1 reaches smaller volumes than V0. Once again, this result is consistent with the uniaxial drawing tests, which showed that the NDR of V1 was smaller than the one of V0, with a bigger gap between the two resins for lower equivalent strain rates. This behaviour is assumed to be due to the higher ability of the molecular chains of V1 to organize themselves in ordered zones, such as mesophases, responsible for a sooner structural hardening that stops the volume inflation of V1.
## 4. Conclusions

An experimental procedure, based on the time-temperature equivalence principle, was used to measure the intrinsic behaviour during free stretch-blowing tests of a virgin commercial polyester before and after mechanical recycling in conditions representative of industrial recycing. The strain hardening phenomena, already known to drive the kinematics of blowing for different virgin polyesters, is also involved in the blowability of mechanically recycled PET. Moreover, the correlations previously observed between the behaviours of virgin polyesters under uniaxial stretching and stretch-blowing, are verified with recycled polyesters. Indeed, in this article, the lower blowability measured after mechanical recycling was correlated with the smaller Natural Draw Ratios measured under uniaxial stretching. Thus, the strategy to classify the blowability of several polyesters by the mean of uniaxial drawing test, thanks to the time-temperature equivalence principle, works also to classify virgin and recycled grades. However, in our case, the recycled PET is not an unknown resin but was directly obtained from the mechanical recycling of the virgin PET in a closed recycling loop without external contaminations. Two perspectives of this work would be interesting to explore. Firstly, by intentionally adding contaminants in the recycling process, it would be possible to verify if the strategy developed here can also be applied to compare a PET with different natures and contents of contaminants. Secondly, this procedure could be tested on several commercial recycled polyesters, taking into account the differences of infrared absorptions between virgin and recycled commercial PET.

### CRediT authorship contribution statement

**Nathan Sylvestre:** Writing – original draft, Validation, Methodology, Investigation, Data curation, Conceptualization. **Jean-Luc Bouvard:** Writing – review & editing, Validation, Supervision, Methodology, Funding acquisition, Conceptualization. **Guillaume Corvec:** Writing – review & editing, Validation, Supervision, Methodology, Investigation, Data curation. **Mikaël Derrien:** Writing – review & editing, Validation, Supervision, Methodology, Funding acquisition, Conceptualization. **Xavier Monnier:** Writing – review & editing, Validation, Methodology, Investigation, Funding acquisition, Conceptualization. **Christelle Combeaud:** Writing – review & editing, Validation, Supervision, Methodology, Funding acquisition, Conceptualization.

### Declaration of competing interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

## Appendix A. Supplementary data

Supplementary data to this article can be found online at https://doi.org/10.1016/j.polymertesting.2025.108774.

### Data availability

No data was used for the research described in the article.

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