Microstructural reorganisation and mechanical property enhancement in PHAs and their wood biocomposites through post-fabrication heat-treatment in annealing and partial-melting conditions.
Present Address:
VincentMathel1,2✉Emailv.mathel@uq.edu.au
PaulineLeQuellec2
ShazedAziz1,2
DarrenMartin1,2
PeterHalley1,2
MichaelTobiasHeitzmann2,3
Luigi-JulesVandi2,3✉Emaill.vandi@uq.edu.au
1School of Chemical EngineeringThe University of Queensland4072BrisbaneQLDAustralia
2Centre for Advanced Materials Processing and Manufacturing (AMPAM)The University of Queensland4072BrisbaneQLDAustralia
3School of Mechanical and Mining EngineeringThe University of Queensland4072BrisbaneQLDAustralia
Vincent Mathel1,2*, Pauline Le Quellec2, Shazed Aziz1,2, Darren Martin1,2, Peter Halley1,2, Michael Tobias Heitzmann2,3, Luigi-Jules Vandi2,3*
1School of Chemical Engineering, The University of Queensland, Brisbane, QLD 4072, Australia
2Centre for Advanced Materials Processing and Manufacturing (AMPAM), The University of Queensland, Brisbane QLD 4072, Australia
3School of Mechanical and Mining Engineering, The University of Queensland, Brisbane QLD 4072, Australia
*Corresponding authors: Vincent Mathel (E-mail: v.mathel@uq.edu.au); Luigi-Jules Vandi (E-mail: l.vandi@uq.edu.au)
Abstract
Polyhydroxyalkanoates (PHAs) are biodegradable bioplastics with strong environmental benefits, yet their inherent brittleness and high production cost limit broader adoption. Blending PHAs with lignocellulosic biofillers offers a circular and cost-effective pathway but often compromises mechanical performance. This study investigates post-fabrication heat-treatment, e.g. annealing (< 150°C) and partial-melting (> 150°C) conditions, as a scalable strategy to tailor the properties of PHAs such as poly(3-hydroxybutyrate-co-3-hydroxyvalerate) (PHBV), their blends with poly(3-hydroxybutyrate-co-4-hydroxybutyrate) (P34HB), and their wood fibre biocomposites. Annealing improved ductility by enhancing the mobile amorphous fraction (MAF) and reducing the rigid amorphous fraction (RAF), while partial melting promoted crystal perfection but induced thermal degradation, increasing RAF and material stiffness. Optimal mechanical performance was achieved after 30 min at 150°C, with tensile strain at break increasing by ~ 650% for neat PHAs and ~ 200% for wood/PHAs biocomposite variants. This was accompanied by a 20–30% reduction in modulus and ≤ 16% drop in tensile stress for the both materials. Notably, shrinkage above 175°C was significant in neat and blended PHAs but was strongly mitigated by wood biofillers. The results highlight post-fabrication heat treatment as a simple, effective method to enhance the mechanical behaviour and dimensional stability of PHAs-based materials for rigid packaging and other demanding applications.
Keywords:
Natural fiber-polymer composites
Microstructure
heat treatment
Rigid amorphous fraction
Structure-property relations
Toughness
1 Introduction
A
Polyhydroxyalkanoates (PHAs) are a promising family of biodegradable biopolymers that offer a sustainable solution to plastic pollution by replacing persistent, fossil-based plastics that accumulate in landfills and marine environments. Unlike conventional plastics, PHAs are both biodegradable and compostable under a broad range of environmental conditions, including industrial and home composting systems [1]. They can be produced via bacterial fermentation using low-cost, second-generation feedstocks such as non-food crop residues, fruit waste, biogas, animal fats, and other agricultural by-products [2, 3]. This not only lowers production costs but also enhances circularity valorising organic waste streams that would otherwise remain unused [4]. From a life cycle assessment perspective, PHAs demonstrate lower global warming potential and acidification potential compared to both conventional plastics and other biodegradable polymers [5]. Within the PHA family, poly(3-hydroxybutyrate-co-3-hydroxyvalerate) (PHBV) is among the most extensively studied and commercially available bioplastics [6]. Despite these environmental advantages, PHAs remain significantly more expensive than conventional plastics, which continues to hinder their market development.To reduce material costs, blending PHAs with biomass fibres derived from lignocellulosic by-products has been reported as an effective strategy, acting as a reinforcing agent and promoting circularity [7–9]. It also commonly results in enhance biodegradation rate of PHA [10–12], and overall contributes to a reduced carbon footprint of the final product. However, the incorporation of organic fibres trends to increase the brittleness of PHAs [7, 13], a drawback that becomes more pronounced at higher fibre loadings [14, 15] and when using already brittle PHA such as PHBV with low comonomer content. These mechanical limitations restrict the potential applications of PHBV-based biocomposites, particularly in injection-moulded rigid packaging.
Various strategies have been employed to enhance the mechanical performance of post-synthesised PHAs, including the modification of their chemical structure [16, 17]. Moreover, techniques such as reactive extrusion [18], fibre surface treatments [19–21], and the use of compatibilisers or interfacial modifiers [22–24] have been explored to improve interfacial interactions fibre–PHA composites, thereby enhancing the resulting physical properties. However, the drawbacks of these approaches are plenty, including increased material and manufacturing expense, complex process control, potential toxicity and challenging scale-up.
Among the processing techniques, physical blending via twin-screw extrusion remains the most widely used, versatile, low-cost, and scalable technique for controlling the properties of PHAs and PHA-based biocomposites. This is typically achieved through the incorporation of plasticisers [25, 26] or ductile biopolymer [27–30]. While plasticisers can effectively enhance toughness, their addition often leads to a significant reduction in thermal stability due to their lower thermal resistance [31]. In addition, plasticised PHA are more susceptible to physical ageing, resulting in a drop in toughness and strain at break over time [32]. Besides, blend of immiscible, non-compatibilized polymers typically leads to a sea-island morphology with weak interfacial adhesion. As a result, the polymers in blends deform or fracture independently, yielding a strain at break that is often lower than the theoretical weighted average [33, 34]. Furthermore, the incorporation of other biopolymers can compromise the unique biodegradability characteristics of PHAs.
With the recent development and commercial availability of more ductile and flexible PHAs [35], blending strategies that combine strong, stiff, but brittle PHA with more flexible and ductile PHA grades have emerged as a promising approach. Such blends have shown good compatibility, offering the potential to tailor the mechanical properties of PHBV [36–38] and PHBV-based biocomposites. However, miscibility depends strongly on the type of PHA, lend ratio [39] and copolymer content [40]. In our previous work, blending neat PHBV and PHBV-wood biocomposite with the more ductile and flexible poly(3-hydroxybutyrate-co-4-hydroxybutyrate) (P34HB) has demonstrated the potential to produce materials with mechanical properties suitable for replacing conventional plastics such as polypropylene or polyethene [13]. Notably, this study provided evidence of inter-diffusion between the two PHAs after blending, as well as good PHA-wood fibre interactions.
The microcrystalline structure of biopolymers and biocomposites plays a key role in mechanical properties of the materials [41, 42]. Semi-crystalline PHAs are particularly known to undergo physical ageing, driven by secondary crystallisation processes and/or a rigidification amorphous phase (increase in rigid amorphous fraction: RAF) which occur gradually under storage conditions. These microstructural rearrangements can lead to significant changes in the PHA’s properties, notably resulting in increased brittleness over time [43, 44].
Isothermal annealing has been widely reported as a simple and effective method to enhance the ductility of semi-crystalline polymers. This heat treatment promotes structural reorganisation within the microcrystalline regions, notably by increasing the relative mobility of macromolecular segments in the amorphous phase through the relaxation of the RAF [45–48]. For instance, Koning et al. observed an increase in strain at break from 5% to over 30% in PHB homopolymer after annealing at temperatures between 110°C and 150°C [46, 49]. However, in PHBV with a high degree of crystallinity, the restricted mobility of macromolecular chains may limit microstructural reorganisation and resulting in only marginal improvements in mechanical properties. To address this limitation, the less commonly reported partial melting approach can be employed. Under these heating conditions, the thinner and less perfect crystals, formed during the injection moulding process, are melted and subsequently recrystallised, allowing for a more effective rearrangement of the crystalline structure.
A recent study demonstrated an increase in strain at break from 4.9% to over 60% by partially melting PHBV at 176°C, followed by recrystallisation under isothermal conditions between 120°C and 150°C [50]. This post-fabrication heat treatment promoted the homogenisation of the crystalline regions as well as reduced molecular chain density and entanglement within the amorphous phase. Consequently, it generates a more uniform distribution of the stretching degree of the macromolecular segments which contributed significantly to the enhanced ductility observed [51].
A
This study investigates and compares the effects of post-fabrication heat treatments under annealing conditions (below the onset of melting temperature of PHA, approximately 150°C) and partial-melting conditions (above the melting temperature, above 150°C) on PHBV, PHBV/P34HB blends, and their corresponding biocomposites reinforced with wood sawdust flour. Specifically, the evolution of PHA molecular weight was measured by gel permeation chromatography (GPC), the crystalline phase was analysed by differential scanning calorimetry (DSC), and the amorphous phase was characterised using dynamic mechanical analysis (DMA). The resulting mechanical properties were evaluated through tensile testing, as well as the shrinkage ratio obtained after post-fabrication heat-treatment. This work aims to enhance understanding of the impact of post-fabrication heat treatment on microcrystalline structure, thermal degradation, and mechanical performance of PHA biopolymers, as well as the consequences of incorporating biofillers. Furthermore, the study aims to highlight a simple, cost-effective, and scalable complementary approach to blending strategies for tailoring the mechanical properties of PHAs and PHA-based biocomposites. Completing on previous work on PHBV/P34HB/wood blends [13], this paper explores opportunities to develop advanced bio-based and biodegradable materials with adaptable properties, tailored to specific product requirements and performance specifications.2 Experimental Section
2.1 Materials
PHBV contained 2 mol% HV comonomer, 1 wt% nucleating agent (boron nitride) and 0.5 wt% antioxidant, in powder form, called ENMAT Y1010 was supplied by TianAn Biopolymer (Ningbo, Zhejiang, China). The low ratio of comonomer and nucleating agent results in a highly crystalline PHBV, characterised as stiff, strong and brittle PHA. P34HB containing 50 mol% 4HB comonomer, was supplied by CJ Cheiljedang (Seoul, South Korea) in pellet form. The high ratio of comonomer results in mainly amorphous P34HB, characterised as flexible, ductile and weak PHA. Radiata Pine Sawdust flour, referred in this paper as wood biofiller, with average particle size of 75 µm, apparent density of 0.34 g/ml, and an aspect ratio of approximately 6, was supplied by Micro milling Australia.
2.2 PHAs blend and biocomposites preparation
Before compounding, biofillers and PHAs were dried at 60°C for 48 hours to achieve a moisture content below 1%. The dried materials were pre-blended manually to obtain a ratio 75 wt.% PHBV/25 wt.% P34HB for the PHAs blend, 80 wt.% PHBV/20 wt.% wood biofiller and 60 wt.% PHBV/20 wt.% P34HB/ 20 wt.% wood biofiller for biocomposites. Compounding was processed by co-rotative twin-screw extruder (Eurolab 16, Thermo Scientific) with a screw ratio of 1:40. The extruder was fed through Brabender volumetric feeders. The screw profile, composed of only one mixing zone, was designed to melt, blend and convey the materials with minimal thermal degradation of PHAs. Screw design and compounding parameters are described in our previous study [13]. To enable a reliable comparison of the different materials following post-fabrication heat treatment, neat PHBV was subjected to the same thermal process as the other samples.
2.3 Specimen preparation
Babyplast 6/10P micro-injector (Ambaldi Group, Italy) was used to produce tensile test specimens (ASTM D638 type V) and slab specimen (64.5x9.53x4) for DMA analyses. The barrel temperatures were set at 178°C, 180°C, and 185°C (at the die), with the mould temperature maintained at 60°C and the injection pressure fixed at 80 bar.
2.4 Post-Fabrication heat treatment
Post-fabrication heat treatment was performed using two ovens. The first oven was employed for post-fabrication heat treatments across a temperature range from 70°C to 185°C. Part of specimens were removed after 5 min of heat treatment and the others after or 30 min. Following this, a second oven was used to control the cooling and stabilise the microcrystalline structure by holding the specimens at 110°C for 5 minutes. For heat treatments conducted below 110°C, the second oven was not used. All specimens were placed on a non-oven textile inside a stainless-steel tray. Transfers between ovens were carried out manually.
Figure 1a summarises all heat treatment conditions, while Fig. 1b shows the positioning of heat treatment temperatures relative to the DSC heating curves of PHBV. It illustrates that heat treatments conducted below 150°C can be classified as annealing, whereas treatments above 150°C, within the melting peak region, correspond to partial melting processes.
The 5 min heat treatment at 70°C was not performed, as this combination of temperature and duration was judged insufficient to induce observable changes in the microcrystalline structure. Similarly, heat treatment at 185°C for 30 min was not conducted because the specimens began to melt under these conditions. Figure 1c provides a schematic representation of the heat treatment setup. To observe the consistency between temperatures inside the tray and parameters in oven, a heat sensor was placed alongside the specimens. Figure 1d shows a typical temperature–time profile recorded by the sensor during the heat treatment process.
Fig. 1
a) summary of the post-fabrication heat treatment conditions, b) heat treatment temperatures positioned on DSC heating curves of PHBV, c) schematic illustration of heat treatment setup and d) time-temperature data obtained from the sensor.
2.5 Characterisation and Testing
Gel permeation chromatography (GPC) was conducted using an Agilent 1260 Infinity II system to assess the molecular weight of PHAs. To separate PHAs from wood biofillers, biocomposite samples were dissolved in chloroform at 80°C for 120 minutes, then filtered through 0.22 µm polytetrafluoroethylene (PTFE) syringe filters (Kinesis, ESF-PT-13-022). The chromatographic system comprised a guard column (Agilent PL-gel, 10 µm, 7.5 mm × 50 mm) followed by three Agilent PL-gel MIXED-B columns (10 µm, 7.5 mm × 300 mm) connected in series, all maintained at a constant temperature of 30°C. Chloroform was used as the mobile phase at a flow rate of 1.0 mL/min, with detection performed via a refractive index (RI) detector. The system was calibrated using narrow-dispersity polystyrene standards. To account for differences in polymer structure and solvation behaviour, Mark–Houwink parameters specific to PHBV in chloroform (K = 7.7 × 10⁻³ mL/g and α = 0.82) [52] were applied within the Agilent GPC data analysis software to ensure accurate molecular weight determination.
Differential scanning calorimeter, DSC Q2000 (TA Instruments, New Castle, United States), calibrated before measurement with indium, was used for thermal analyses. For each analysis, 4.9 to 5.4 mg of sample were encapsulated in an aluminium T-zero standard pan and lid. During the analysis, samples were first heated from − 50°C to 190°C at a rate of 10°C/min and subsequently cooled to − 50°C at the same heating rate, under a controlled N2 flow of 50 ml.min− 1. Melting temperatures (Tm), temperature of crystallisation (Tc) and melting enthalpy were collected. Melting enthalpy of PHA(s) (
in biocomposite was calculated as:
1
is the enthalpy measured from DSC curve for the biocomposite and
is the weight fraction of wood biofiller in the biocomposite and equal to 20 wt.%. The qualitative degree of crystallinity (Xc) was obtained by dividing
with the theoretical melting enthalpy of 100% crystallised PHB equal to 146 J.g− 1 [53].
Dynamic mechanical analysis (DMA) was performed using a TA Instruments Discovery Series Hybrid Rheometer (HR30) (TA Instruments, New Castle, United States) equipped with a 3-point bending geometry (25 mm span). Relaxation behaviours of materials were investigated over a temperature range of 50°C to 120°C. Measurements were conducted under a nitrogen atmosphere at a heating rate of 5°C/min, using a fixed frequency of 1 Hz and an axial displacement amplitude of 25 µm. Form this analysis, loss modulus, storage modulus and loss factor (tan ẟ) were collected.
Tensile properties of the biocomposites were evaluated using an Instron 5548 Electromechanical Testing System, equipped with a 1 kN load cell, 1 kN pneumatic grips, and a video extensometer to measure accurate strain measurement. Testing was carried out in accordance with ASTM D638 standards. Specimens were subjected to uniaxial loading at a constant crosshead displacement rate of 1 mm·min⁻¹. The mechanical properties measured included tensile strain at break (%), tensile stress at maximum load (MPa), and tensile modulus (GPa). To ensure statistical reliability, five replicates were tested for each sample.
3 Results and discussion
3.1 Molecular weight
Figure 2 illustrates the evolution of PHA molecular weight as a function of heat treatment for neat PHBV, PHBV/P34HB blend, and their corresponding wood biofillers biocomposites. The obtained values are reported in Table 1.
The molecular weights of neat PHBV and PHBV/P34HB blend are similar. Incorporation of wood biofillers leads to a notable decrease in molecular weight of PHA. This reduction is likely due to increased frictional forces during melt processing, due to the physical presence of the wood biofillers. These forces increase the shear rates and mechanical stress, thereby accelerating thermomechanical degradation of the PHA matrix through chain scission.
This effect is more pronounced in the PHBV/P34HB blend, likely due to the greater sensitivity of P34HB to thermo-mechanical degradation in the presence of wood-based biofillers, as previously reported [13]. All post-treated materials exposed to temperatures above 150°C for 30 min exhibit a linear decrease in PHA molecular weight with increasing heat treatment temperature. This reduction is more pronounced (steeper slope) in neat PHBV and the PHBV/P34HB blend, indicating a higher thermal sensitivity of these materials compared to their wood-filled counterparts.
Post-fabrication heat treatment temperature (℃) | Molecular Weight (10-3g.mol-1) | |||||||
|---|---|---|---|---|---|---|---|---|
PHBV | PHBV/P34HB | PHBV-Wood | PHBV/P34HB-Wood | |||||
↓ | ||||||||
Duration (min) → | ||||||||
5 | 30 | 5 | 30 | 5 | 30 | 5 | 30 | |
Non-treated | 282.9 | 277.2 | 233.1 | 199.9 | ||||
70 | 278.3 | 271.5 | 240.2 | 214.8 | ||||
100 | 283.5 | 271.3 | 238.0 | 205.2 | ||||
125 | 280.1 | 267.4 | 222.2 | 198.6 | ||||
150 | 283.8 | 260.9 | 257.4 | 256.8 | 213.1 | 199.5 | 195.7 | 185.3 |
160 | 252.9 | 249.5 | 179.1 | 170.7 | ||||
170 | 230.1 | 225.1 | 159.9 | 152.1 | ||||
175 | 283.3 | 196.0 | 286.0 | 199.3 | 222.9 | 151.6 | 189.5 | 134.7 |
180 | 175.6 | 175.3 | 144.7 | 125.6 | ||||
185 | 270.8 | 210.8 | 218.8 | 195.4 | ||||
This may be attributed to the partial absorption or dissipation of thermal energy by the wood biofillers during heat treatment, acting as “thermal insulator” and limiting the effective calorific energy transferred (or thermal conductivity) to the PHA matrix [54, 55].
Fig. 2
Post fabrication heat treatment vs PHA molecular weight (bleu: 5 min, black: 30 min)
For short times of heat treatment (5 min), the PHA molecular weight remains largely preserved across these investigated temperatures range.
3.2 Thermal analyses
Figure 3 presents the DSC thermograms of neat PHBV, PHBV/P34HB blend, and their corresponding wood biofiller biocomposites.Incorporation of P34HB into PHBV leads to a reduction in its melting temperature (Tm) by approximately 1°C, reduction in crystallisation temperature (Tc) by around 7°C, and reduction of the degree of crystallinity (Xc) by approximately 20%. The decrease in crystallinity is proportional to the amount of P34HB added to the PHBV matrix. The melting temperature (Tm) is typically associated with crystallite stability, lamellar thickness, and, more broadly, crystallite perfection, while the crystallisation temperature (Tc) reflects the crystallisation kinetic. Combined incorporation of P34HB and wood biofillers further reduces the Tm, Tc, and degree of crystallinity of PHBV. As a result, the PHBV/P34HB–wood biocomposite exhibits the lowest Tm and crystallinity values among the samples studied. This reduction can be attributed to molecular interdiffusion between PHBV and P34HB, as well as interactions between the PHA and wood biofiller. These effects disrupt the regular packing of PHBV chains, resulting in less ordered and thermally stable crystalline structures [13]. A shift of Tc in higher temperature observed upon the addition of wood biofiller to neat PHBV may suggest higher PHA-wood interactions in this biocomposite.
Fig. 3
a) heating curves and b) cooling curves obtained by DSC.
Figure 4 show the evolution of Tm and Xc as a function of post-fabrication heat treatment conditions for all materials, based on data extracted from the DSC curves presented in Fig.S1 (Supplementary Data). The obtained values are reported in Table 2.
Crystallinity degree - Xc (%) | ||||||||
|---|---|---|---|---|---|---|---|---|
Non-treated | 64 | 45 | 56 | 36 | ||||
70 | - | - | - | 48 | - | 62 | - | - |
100 | 64 | 64 | 46 | 49 | 59 | 63 | 39 | 45 |
125 | 65 | 68 | 46 | 55 | 60 | 64 | 42 | 49 |
150 | 66 | 71 | 46 | 56 | 55 | 70 | 47 | 52 |
160 | 71 | 74 | 53 | 59 | 61 | 71 | 45 | 54 |
170 | 70 | 76 | 54 | 61 | 65 | 79 | 46 | 55 |
175 | 72 | 74 | 55 | 53 | 72 | 80 | 49 | 52 |
180 | 74 | 78 | 53 | 52 | 66 | 83 | 46 | 48 |
185 | 76 | - | 51 | - | 68 | - | 39 | - |
Below a heat treatment temperature of 150°C, Tm remains unchanged, while the degree of crystallinity increases for all materials with increasing heat treatment temperature. This effect is more pronounced after 30 min of heat treatment compared to 5 min, indicating time-dependent crystalline reorganisation.
Crystallisation that occurs during the injection-moulding process can induce residual stresses and oriented chains within the microcrystalline structure which can be relieved through annealing (relaxation), facilitating macromolecular rearrangement in less constrained amorphous regions and further crystalline reorganisation [45, 46, 48]. The resulting increase in crystallinity is attributed to this re-adjustment [49, 56]. The constancy of the melting temperature (Tm) under annealing conditions has been previously reported, even though an increase in lamellar thickness may occur during this process [49]. This is because the measured Tm primarily reflects the melting of crystallites that have undergone reorganisation during the DSC heating cycle. Therefore, the melting point of crystallites arrangement from annealing may be significantly lower.
Since the crystallisation process of PHBV is initially hindered by the presence of P34HB and wood biofillers, leading to a higher proportion of disordered macromolecular regions, the effect of annealing, driven by relaxation and reorganisation processes, should be more pronounced. As a result, the degree of crystallinity in the PHBV/P34HB blend and its associated biocomposite increases more significantly with annealing, gradually approaching the crystallinity level observed in neat PHBV.
Fig. 4
Evolution of Tm (up) and Xc (bottom) for a post-fabrication heat treatment of 5 min (left) and 30 min (right)
Above a heat-treatment temperature of 150°C, increasing the treatment temperature for a duration of 30 min leads to a significant rise and convergence of the Tm values across all materials. This phenomenon can be attributed to the melting of less stable crystalline phases, which re-crystallise into more thermodynamically stable and structurally perfect crystallites, leading to shift of Tm.
This increase is combined by an increase in the degree of crystallinity for PHBV and its biocomposites. Phenomenon of relaxation and rearrangement of macromolecules around non-melting, more perfect crystallites may also contribute to the observed crystalline properties.
In contrast, the PHBV/P34HB blend and its biocomposites exhibit a reduction in crystallinity above a heat treatment of 165°C. This reduction in degree of crystallinity should be attributed to the thermal degradation of P34HB, which likely generates structural defects and increases chain disorder (entropy) along the P34HB macromolecules [57]. These defects disrupt the crystallisation of PHBV. This reduction is not observed after short heat treatments of 5 min as materials don’t have time for degrading. Regarding evolution of crystalline properties, heat treatment under partial-melting conditions and for extended durations has a more pronounced impact on the microcrystalline structure, promoting formation of more stable crystals with a maximum increase of Tm at 180℃ and degree of crystallinity by about + 12℃ and + 22% for neat PHBV, + 10℃ and + 16% for PHBV/P34HB blend, + 12℃ and + 48% for PHBV-wood biocomposites, + 14℃ and 33% for PHBV/P34HB-wood biocompositesIt has been reported in the literature that these increases in crystallinity and crystallite perfection do not affect the crystallite morphology (crystal lattice) [50, 51]. Perfectly understood the evolution of the microcrystalline structure and its associated thermal transitions can be complex and particularly in injection-moulded samples where crystallinity can vary across the thickness of the specimen and present a skin/core crystalline morphology. While DSC provides valuable insights into bulk crystallinity, it cannot fully elucidate the evolution of the microcrystalline structure, particularly in injection-moulded samples where this structure is not uniform across the specimen thickness. For that, additional crystallographic techniques, such as Small- and Wide-Angle X-ray Scattering (SAXS/WAXS), are required.Dynamic mechanical analyses
At -40°C, PHBV, the PHBV/P34HB blend, and their corresponding biocomposites are in the glassy state, where polymer chains are frozen, resulting in increased material stiffness. As shown in Fig. 5b, the incorporation of P34HB reduces the storage modulus, whereas the addition of wood biofillers significantly increase it. Consequently, the PHBV-based biocomposite containing wood biofillers exhibits the highest stiffness at -40°C.
Fig. 5
DMA curves of a) tan δ, b) storage modulus, c) loss modulus in function of temperature.
Between − 40°C and 45°C, a relaxation peak is observed in both the tan δ (loss factor) and loss modulus curves (Fig. 4a and .4b). This peak can be associated with energy dissipation due to the motion of polymer chain segments within the MAF [47, 58, 59]. Width and intensity of this relaxation peak are more pronounced in the PHBV/P34HB blend, likely due to the combined contribution / convolution of the MAF from both PHBV and amorphous P34HB. It can be attributed to the higher fraction of polymer segments present in this phase and/or their ability to gain mobility as the temperature increases. The greater energy dissipation is also illustrated by the more significant drop in storage modulus observed between − 40°C and 45°C. At these temperatures, the unfreezing of polymer segments within the MAF leads to a reduction in material stiffness. Temperature corresponding to the tan δ peak is commonly identified as the glass transition temperature (Tg) of the polymer. Addition of P34HB to the PHBV matrix results in a 2°C reduction in Tg. Temperature associated with the loss modulus peak, which typically indicates the onset of segmental motion of macromolecules within the MAF, usually occurs at a lower temperature than tan δ peak [60]. For the PHBV/P34HB blend, this peak appears 15°C lower than that of neat PHBV. These reductions are attributed to the increased mobility of macromolecular segments within the MAF, facilitated by the presence of the more flexible P34HB chains.
Addition of wood biofiller to both PHBV and the PHBV/P34HB blend results in a flattening, broadening of the tan δ curve (Fig. 4a) and an increase in storage modulus (Fig. 4b). This behaviour indicates a reduction in macromolecular segment mobility, suggesting a decrease in the MAF, likely due to the presence of the biofiller and potential interactions between the wood fibres and PHAs [13, 23, 61, 62]. In PHBV, addition of wood leads to an increase in Tg by approximately 2.5°C, whereas in the PHBV/P34HB blend, it results in a decrease of about 1.5°C. This contrast highlights the specific interactions between PHBV and wood in the neat matrix, which are diminished or lost upon the incorporation of P34HB into the blend.
Above 45°C, a second relaxation peak is clearly visible in the loss modulus curve of PHBV and its biocomposite. This peak can be assigned to the dissipation of energy of the RAF [47, 63, 64]. This fraction includes all the amorphous chain trapped or constrained (especially at interface with crystallites) occurring in the injection moulding process. This second relaxation is accompanied by a further drop in the storage modulus (E'), reflected by the onset of segmental mobility within the RAF fraction [65–67] .
This relaxation peak disappears upon the addition of P34HB to PHBV and in the corresponding biocomposites, likely due to a reduction in molecular restriction resulting from blending with the more amorphous P34HB. Consequently, a smaller drop is observed in E' within this temperature range. However, this E' drop still suggests the presence of energy dissipation mechanisms, potentially associated with a less constrained rigid amorphous fraction (RAF).
In this paper, it was decided to assess the MAF and RAF contributions (%MAF and %RAF, respectively) on storage modulus (degree of dissipative energy associated to these fractions) following equations and Fig. 5b:
2
3
A
FigureS2 (Supplementary Data) gathered the evolution of storage modulus curves for the different materials as a function of heat treatment. Figure 6a and .6b and Table S1 (Supplementary Data) present, respectively, %MAF and %RAF values extracted from these curves and their evolution with heat treatment conditions. As expected, the PHBV/P34HB blend exhibits a higher %MAF than neat PHBV, while the addition of wood biofillers to the PHAs blend results in a reduction of %MAF and a corresponding increase in %RAF. For PHBV, however, the addition of wood does not show a clear reduction in %MAF.Below heat treatment temperatures of 150°C, %MAF increases with heat treatment temperature, while %RAF decreases. This trend is attributed to annealing conditions, which promote the relaxation of internal constraints in the amorphous phase, thereby enhancing segmental mobility. In this range of heat treatment temperatures, a time-dependent effect is observed, wherein a shorter heat treatment duration (5 min) results in a lower %MAF and higher %RAF, particularly in PHBV and its biocomposites. This behaviour may be attributed to the higher degree of crystallinity and crystallite perfection in PHBV, which is likely to generate greater internal residual stresses and metastable phases. These characteristics increase the sensitivity of neat PHBV to heat treatment compared to the more amorphous PHBV/P34HB blend. As a result, longer heat treatment durations promote relaxation of internal constraints and macromolecular mobility within the amorphous phase.
For heat treatment durations of 5 min above 150°C (partial-melting condition), %MAF continues to increase and %RAF to decrease with rising temperature. However, for 30 min of heat treatment above 150°C, %MAF decreases and %RAF increases. This reversal is likely associated with the formation of a more perfect crystalline structure and a higher degree of crystallinity in comparison to annealing condition. Such phenomena should restrict macromolecular mobility in the amorphous phase [68], despite the concurrent reduction in molecular weight and therefore the entanglement ratio [69], which would usually promote mobility. Figure 6c presents the evolution of the %MAF/%RAF ratio. An trend of increase in this ratio with heat treatment temperatures below 150°C indicates a slight reduction in the RAF contribution, is accompanied by a relatively greater contribution of the MAF to the material properties, even though the overall crystalline fraction in the materials increases. In contrast, above 150°C, an increase in %RAF is accompanied by a greater decrease in %MAF, leading to a reduction in the %MAF/%RAF ratio with increasing heat treatment temperature.
The close %MAF/%RAF ratios observed between PHBV and PHBV-wood biocomposites could be explained by a balance between the less perfect crystalline structure (lower Tm and slightly reduced degree of crystallinity for the biocomposite) and the specific interactions between wood and PHBV, both of which influence macromolecule segment mobility in amorphous phase.
Fig. 6
Evolution of of a) %MAF contribution, b) %RAF contribution, c) ratio %MAF/%RAF as a function of post-fabrication heat treatment temperature and duration: 5 min (blue), 30 min (back)
A
As observed in Fig.S3 and Table S2 (Supplementary Data), although the Tg trends to decrease below 150°C and increase thereafter, its overall evolution appears rather erratic. This highlights the complexity of microcrystalline changes upon heat treatment, particularly upon finer examination. Such behaviour may be attributed to the inherent structural heterogeneity of injection-moulded parts, which typically exhibit a non-uniform skin/core morphology and variations in crystallinity across the part thickness.3.3 Tensile properties
As observed in Fig. 7 and reported in Table 3, incorporation of P34HB into PHBV results in a decrease in tensile modulus by 49% (i.e., reduced stiffness, as also observed in the DMA), a decrease in tensile stress (at maximum load) by 27% and an increase in elongation at break by 339%. The addition of wood biofiller results in an increase of the tensile modulus by 49% for PHBV and by 40% for PHBV/P34HB blend. It also results in a decrease in tensile strain at break by 62% for PHBV and by 75% for PHBV/P34HB blend. The tensile stress slightly increases by 5% with addition of wood in PHBV matrix while it decreases by 14% with addition of wood in PHBV/P34HB. As previously reported, it suggests specific interaction between PHBV matrix and wood that disappear with addition of P34HB in the blend [13].
Fig. 7
Evolution of a) tensile modulus, b) tensile stress of mateirals, tensile strain at break of c) PHBV and biomcomposite and d) P34HB/PHBV and biocomposites as a function of heat treatment temperature and duration: 5 min (blue) and 30 min (black).
Under annealing conditions (below 150°C), both tensile modulus and tensile stress trend to decrease while the tensile strain at break increases (Fig. 7). This change of properties follow the same trend as the increase in %MAF and decrease in %RAF. Even if the degree of crystallinity increases moderately, relaxation of macromolecules in amorphous phase results in a more ductile and flexible materials.
In contrast to the other properties that follow a same range of evolution, the increase in strain at break is more significant in the neat PHAs matrix than in the biocomposites with increasing temperature of heat treatment.
Post-fabrication heat treatment temperature (℃) | Tensile modulus (GPa) | |||||||
|---|---|---|---|---|---|---|---|---|
PHBV | PHBV/P34HB | PHBV-Wood | PHBV/P34HB-Wood | |||||
↓ | ||||||||
Duration (min) → | ||||||||
5 | 30 | 5 | 30 | 5 | 30 | 5 | 30 | |
Non-treated | 3.3 ± 0.1 | 1.7 ± 0.1 | 4.8 ± 0.6 | 2.3 ± 0.2 | ||||
70 | - | 3.2 ± 0.1 | - | 1.7 ± 0.1 | - | 4.7 ± 0.5 | - | 2.5 ± 0.2 |
100 | 3.1 ± 0.1 | 3.1 ± 0.1 | 1.6 ± 0.1 | 1.6 ± 0.1 | 4.2 ± 0.2 | 4.3 ± 0.2 | 2.3 ± 0.3 | 2.5 ± 0.3 |
125 | 3.0 ± 0.2 | 2.8 ± 0.2 | 1.5 ± 0.1 | 1.4 ± 0.1 | 4.0 ± 0.0 | 3.8 ± 0.2 | 2.2 ± 0.1 | 2.1 ± 0.1 |
150 | 2.9 ± 0.2 | 2.5 ± 0.2 | 1.4 ± 0.1 | 1.2 ± 0.0 | 3.9 ± 0.1 | 3.7 ± 0.3 | 2.1 ± 0.1 | 1.8 ± 0.1 |
160 | 2.5 ± 0.1 | 2.7 ± 0.2 | 1.2 ± 0.1 | 1.2 ± 0.0 | 4.0 ± 0.2 | 3.7 ± 0.3 | 2.1 ± 0.2 | 1.8 ± 0.1 |
170 | 2.5 ± 0.2 | 2.7 ± 0.08 | 1.2 ± 0.1 | 1.2 ± 0.1 | 3.8 ± 0.4 | 4.3 ± 0.1 | 2.0 ± 0.1 | 1.9 ± 0.1 |
175 | 2.5 ± 0.2 | 3.2 ± 0.2 | 1.3 ± 0.1 | 1.2 ± 0.1 | 3.9 ± 0.2 | 4.4 ± 0.1 | 2.0 ± 0.1 | 1.9 ± 0.2 |
180 | 2.4 ± 0.2 | 3.0 ± 0.4 | 1.1 ± 0.1 | 1.3 ± 0.2 | 3.8 ± 0.3 | 4.8 ± 0.3 | 1.9 ± 0.1 | 2.2 ± 0.2 |
185 | 2.3 ± 0.2 | - | 1.0 ± 0.1 | - | 3.6 ± 0.3 | - | 1.9 ± 0.1 | - |
Tensile stress (at maximum load) (MPa) | ||||||||
Non-treated | 35.4 ± 1.5 | 25.9 ± 0.2 | 37.2 ± 1.0 | 22.2 ± 0.5 | ||||
70 | - | 35.3 ± 0.9 | - | 25.9 ± 0.7 | - | 35.9 ± 1.5 | - | 22.1 ± 0.7 |
100 | 35.9 ± 0.8 | 34.7 ± 0.7 | 26.0 ± 0.4 | 24.6 ± 0.4 | 34.7 ± 1.2 | 35.6 ± 1.3 | 21.4 ± 0.7 | 22.1 ± 0.6 |
125 | 34.1 ± 1.1 | 33.4 ± 0.8 | 24.8 ± 0.4 | 23.7 ± 0.8 | 35.5 ± 0.9 | 33.8 ± 1.3 | 20.8 ± 0.4 | 20.2 ± 0.3 |
150 | 33 ± 1.0 | 30.6 ± 0.7 | 23.4 ± 0.6 | 21.7 ± 0.6 | 34.9 ± 0.8 | 33.8 ± 0.2 | 20.3 ± 0.2 | 19.1 ± 0.3 |
160 | 33.8 ± 0.8 | 30.5 ± 0.5 | 22.7 ± 0.8 | 21.1 ± 0.3 | 34.1 ± 0.3 | 33.2 ± 0.8 | 20.3 ± 0.3 | 18.4 ± 0.2 |
170 | 32.5 ± 1.0 | 30.5 ± 0.5 | 22.5 ± 0.6 | 20.6 ± 0.5 | 34.2 ± 1.1 | 33.7 ± 0.9 | 19.5 ± 0.5 | 17.7 ± 0.5 |
175 | 32.6 ± 1.2 | 32.4 ± 1.5 | 22.8 ± 0.7 | 20.6 ± 0.6 | 33.8 ± 0.6 | 35.6 ± 0.6 | 19.8 ± 0.5 | 17.4 ± 0.8 |
180 | 31.7 ± 0.4 | 33.1 ± 0.5 | 22.2 ± 0.5 | 20.4 ± 0.8 | 33.9 ± 0.4 | 35.7 ± 0.9 | 19.2 ± 0.8 | 17.5 ± 0.5 |
185 | 30.2 ± 0.6 | - | 21.2 ± 0.4 | - | 33.7 ± 1.4 | - | 18.6 ± 0.5 | - |
Tensile strain at break (%) | ||||||||
Non-treated | 3.3 ± 0.4 | 33.9 ± 17.4 | 1.3 ± 0.1 | 8.5 ± 0.8 | ||||
70 | - | 3.6 ± 0.4 | - | 57.8 ± 19.4 | - | 1.3 ± 0.2 | - | 9.0 ± 0.6 |
100 | 4.9 ± 1.8 | 7.9 ± 2.0 | 61.5 ± 23.1 | 113.7 ± 39.8 | 1.1 ± 0.1 | 1.8 ± 0.2 | 8.4 ± 2.3 | 12.4 ± 2.3 |
125 | 12.4 ± 8.0 | 18.9 ± 6.9 | 107.6 ± 32.9 | 165.4 ± 65.6 | 2.2 ± 0.5 | 2.6 ± 0.4 | 15.1 ± 2.0 | 24.8 ± 2.5 |
150 | 14.0 ± 9.7 | 25.0 ± 7.0 | 135.1 ± 35.7 | 268.4 ± 51.4 | 2.2 ± 0.4 | 3.4 ± 0.9 | 16.9 ± 3.6 | 31.4 ± 5.8 |
160 | 16.4 ± 5.3 | 37.5 ± 10.3 | 151.6 ± 56.7 | 205.0 ± 51.4 | 3.3 ± 0.3 | 3.5 ± 0.4 | 20.5 ± 2.2 | 30.6 ± 7.5 |
170 | 21.8 ± 5.4 | 36.6 ± 10.7 | 188.5 ± 60.6 | 124.7 ± 46.5 | 3.1 ± 0.3 | 3.5 ± 0.5 | 30.0 ± 9.3 | 23.4 ± 3.7 |
175 | 18.0 ± 9.1 | 34.7 ± 14.2 | 182.1 ± 58.7 | 104.5 ± 29.4 | 4.1 ± 0.7 | 2.8 ± 0.3 | 25.7 ± 3.9 | 13.1 ± 2.9 |
180 | 46.5 ± 30.6 | 27.2 ± 24.6 | 233.3 ± 76.0 | 49.7 ± 21.0 | 4.1 ± 0.6 | 2.0 ± 0.5 | 25.7 ± 6.8 | 6.4 ± 1.4 |
185 | 30.4 ± 8.4 | - | 198.7 ± 77.7 | - | 3.7 ± 0.2 | - | 32.9 ± 4.6 | - |
This is likely due to the physical presence of wood biofillers that restrict macromolecules distortion in the constraint direction during the tensile test [70] or/and the initiation of fracture at the interface wood-PHAs [71] that counterbalance with effect of change in the micro-crystalline structure on tensile strain at break.
Figure 7b shows a decrease in tensile stress at maximum load within the same range for all the materials when heat-treatment temperature increases, except for the PHBV-wood fibre biocomposite. For this last one, the decrease is less pronounced, possibly due to specific interactions between the biofiller and the PHBV matrix that counterbalance the effect of change in the micro-crystalline structure. Therefore, the tensile stress difference between the biocomposite and neat PHBV increases with rising heat treatment temperature. Notably, the biocomposite treated at 150°C for 30 min exhibits a tensile stress approximately 10% higher than that of the neat PHBV under the same conditions.
At a 5 min heat treatment, the decrease in tensile modulus and stress of biocomposites with increase in the heat treatment temperature is less pronounced than PHA(s) matrix, which it is not the case at longer heat treatment time. This suggests that the wood biofiller may act as thermal insulator, thereby delaying or reducing the influence of heat treatment, especially when the treatment duration is short.
For the short-term heat treatment, it can also be noticed that both tensile modulus and stress are higher, while the strain at break is lower, compared to a 30 min heat treatment and even if the degree of crystallinity and crystallite perfection are globally higher at longer heat treatment. This behaviour may be attributed to insufficient thermal energy and duration to fully relieve internal constraints within the amorphous regions of the polymer. This interpretation is supported by the overall lower percentage of %MAF and higher percentage of %RAF observed at 5 min. It indicates that the amorphous phase has a more pronounced influence on mechanical properties than the crystalline phase under these conditions.
In partial-melting condition (heat treatment above 150℃), tensile modulus and tensile stress still continue to decrease and tensile strain at break to increase for 5 min of heat treatment. For a longer heat treatment of 30 min, tensile modulus trends to increase while tensile strain at break decreases for all materials. This trend could be the consequence of the higher crystallite perfection and degree of crystallinity with longer heat treatment time, that increases %RAF, decrease %MAF, increase the stiffness and reduce the capability of distortion of macromolecules in the tensile constraint direction (reduce the ductility). In these range of heat treatment temperatures, reduction of molecular weight of PHAs should also affect the tensile strain at break [72]. Two distinct trends can be observed in the tensile stress behaviour. The tensile stress of PHBV and its corresponding biocomposites increases, which may be attributed to a higher prevalence of the crystalline phase under these conditions (Xc > 70%), thereby generate stronger semi-crystalline structure. In contrast, the tensile stress of the PHBV/P34HB blend and its biocomposites decreases. This reduction may result from the greater thermal degradation sensitivity of P34HB, which could lead to a decrease in the degree of crystallinity and, consequently, to the formation of a weaker chemical and/or microcrystalline structure [73].
As a general observation, change in the microcrystalline structure induced by heat treatment, both in terms of temperature and duration, can significantly influence the mechanical properties of biopolymers and biocomposites, potentially resulting in tougher materials. This effect can be illustrated in annealing conditions, where a 30 min treatment at 150°C leads to a substantial increase in ductility, approximately 650% for neat PHAs and around 200% for PHAs blended with wood biofillers. Concurrently, the tensile modulus decreases by 20–30%, indicating enhanced flexibility, while tensile stress shows a moderate reduction of no more than 16%.
These changes, in annealing condition, appear to be more strongly influenced by modifications in the amorphous phase, particularly the contribution between the %MAF and the %RAF, rather than by a direct correlation with increases in crystalline fraction or perfection.
During partial-melting treatment, the more pronounced transformation and predominance of the crystalline structure results in increased rigidification and embrittlement of the material as the heat treatment temperature and duration rise. Under these conditions, it can be assumed that changes in the crystalline structure exert a greater influence than those observed during annealing. This process indirectly affects the relative contribution ratio of %MAF to %RAF, likely due to the restricted mobility of macromolecules between crystalline domains.
Based on the data presented in Table 4, the PHBV/P34HB blend, thermal post-treated at 150℃ for 30 min exhibits properties suitable for substituting conventional HDPE and LDPE in injection-moulded rigid packaging applications.
Properties | PP* | HDPE** | LDPE** | PHBV | PHBV-wood | PHBV/P34HB | PHBV/P34HB-wood | |
|---|---|---|---|---|---|---|---|---|
Post-fabrication heat treatment at 150℃ for 30 min | ||||||||
Tensile Modulus (GPa) | 1.3–1.8 | 0.8–1.5 | 0.1–0.3 | 2.5 ± 0.2 | 3.7 ± 0.3 | 1.2 ± 0.0 | 1.8 ± 0.1 | |
Tensile stress (MPa) | 30–40 | 25–35 | 10–20 | 31 ± 0.7 | 34 ± 0.2 | 22 ± 0.6 | 19 ± 0.3 | |
Tensile strain at break (%) | 10–50 | 100–600 | 200–600 | 25 ± 7.0 | 3.5 ± 0.9 | 270 ± 51.4 | 31 ± 5.8 | |
| *LyondellBasell Moplen & Pro-fax datasheets | ||||||||
| **Fictiv guide; Zetar Mold PE datasheet | ||||||||
Both PHBV and PHBV/P34HB-wood composites in the same heat treatment conditions, demonstrate performance characteristics that are competitive with injection moulding-grade polypropylene. However, while the PHBV-wood composite offers higher stiffness and mechanical resistance compared to traditional polyolefins, it remains relatively brittle.
3.4 Shrinkage
Shrinkage control is essential to produce dimensionally accurate, functional plastic parts and respect the products specification. For most engineering plastics, the typical mould shrinkage acceptable is below 2%, often ranging between 0.5% and 2.5%, depending on the specific polymer and its degree of crystallinity (ISO 20457). In this study, shrinkage was evaluated from percentage change in length of the tensile specimen after heat treatment in function of specimen length from footprint in mould (63.5 cm according to ASTM 638- Type V) and results are summarised in Table 5.
Post-fabrication heat treatment temperature (℃) | Shrinkage (%) | |||||||
|---|---|---|---|---|---|---|---|---|
PHBV | PHBV/P34HB | PHBV-Wood | PHBV/P34HB-Wood | |||||
↓ | ||||||||
Duration (min) → | ||||||||
5 | 30 | 5 | 30 | 5 | 30 | 5 | 30 | |
Non-treated | -1.26% | -1.26% | -0.16% | -0.16% | ||||
70 | - | -1.42% | - | -1.26% | - | -0.16% | - | -0.16% |
100 | -1.26% | -1.73% | -1.10% | -1.42% | -0.16% | -0.31% | 0.00% | -0.31% |
125 | -1.26% | -1.73% | -1.26% | -1.26% | -0.16% | -0.47% | 0.00% | -0.16% |
150 | -1.42% | -1.73% | -1.26% | -1.42% | 0.00% | -0.31% | -0.16% | -0.16% |
160 | -1.57% | -2.05% | -1.26% | -1.57% | -0.16% | -0.31% | 0.00% | -0.16% |
170 | -1.57% | -2.36% | -1.26% | -2.05% | -0.16% | -0.47% | -0.16% | -0.47% |
175 | -1.73% | -2.83% | -1.26% | -2.83% | -0.16% | -0.79% | 0.00% | -0.79% |
180 | -1.89% | -5.35% | -1.42% | -5.83% | -0.31% | -1.89% | -0.16% | -2.36% |
185 | -1.89% | - | -1.42% | - | -0.31% | - | -0.31% | - |
As can be observed on Fig. 8, heat treatment, especially in partial-melting conditions (above 150℃) for 30 min exposition results to significant shrinkage which became critical from 175℃ for neat PHBV and PHBV/P34HB blend. A higher shrinkage of a about 6% is obtained for PHBV/P34HB blend.It can be noticed that addition of wood preventing excessive shrinkage even at the highest heat post-treatment. Finally, Dimensional change is not for short heat treatment of 5 min even at 185℃
Fig. 8
Evolution of shrinkage in materials as a function of post-fabrication heat treatment temperature and duration: 5 min (blue) and 30 min (black).
4 Conclusions
This study demonstrated that heat treatment under annealing and partial-melting conditions can be used to tailor the properties of PHBV, PHBV/P34HB blends, and PHAs-based biocomposites reinforced by wood biofillers. Under annealing conditions, specifically below the onset of the melting peak (150°C in this study), flexibility and ductility of the materials improved. This enhancement is likely due to the reorganisation within the amorphous phase, increased the relative contribution of mobile amorphous fraction (MAF) and reduced relative contributions of rigid amorphous fraction (RAF) (%MAF/%RAF).
In contrast, under partial-melting conditions (above 150°C), the materials exhibited re-embrittlement and increased stiffness compared to those annealed. This behaviour is likely attributed to significant changes in the crystalline structure, including increases in melting temperature (Tm) and degree of crystallinity (Xc). These changes, affect the %MAF/%RAF ratio by restriction of mobility of segemtn in amorphous phases. Reduction in molecular weight due to thermal degradation occurring during the heat treatment can also induce reduction in tensile strain at break.
This work highlights the critical role of microstructural reorganisation in determining the mechanical properties of these materials. These effects are not observed with short-duration heat treatments (e.g., 5 min). Shrinkage induced by heat treatment becomes significant above 175°C for both PHBV and PHBV/P34HB blends. However, the incorporation of wood biofillers into the PHA matrix substantially reduces this shrinkage, even under high heat treatment temperatures.
Among the conditions tested, the optimal mechanical performance for all materials was achieved at 150°C for 30min, with properties that are competitive with conventional polyethylene and polypropylene grades used in injection-moulded rigid packaging applications.
Overall, this investigation demonstrates a simple, cost-effective, and scalable strategy to adapt the mechanical properties of PHAs and their blends, and PHAs-based biocomposites to meet specific product performance requirements.
Further studies should be required to deepen our understanding of the effects of post-fabrication heat treatments on the microcrystalline structure and the resulting physical properties. Such studies should aim to provide a deeper understanding the degree of impact of PHA thermal degradation versus post-fabrication heat treatment conditions on mechanical properties, investigate the evolution of the crystalline microstructure across the specimen thickness using crystallographic techniques or assess long-term stability (e.g. physical aging).
A
Acknowledgement
The authors would like to thank the Queensland Government for funding through the Advance Queensland Industry Research Fellowships Scheme (Grant No. AQIRF065-2019RD2). The authors also wish to thank the Materials Testing Laboratory from School of Mechanical & Mining Engineering.
Conflict of Interest
The authors declare no conflict of interest.
A
Author Contribution
**Vincent Mathel:** Conceptualization, Data curation, Formal analysis, Investigation, Methodology, Project administration, Visualization, Writing – original draft, Writing – review and editing. **Pauline Le Quellec:** Data curation, Formal analysis, Investigation. **Shazed Aziz:** Visualization, Writing – review and editing. **Darren Martin:** Supervision, Writing – review and editing. **Michael Tobias Heitzmann:** Formal analysis, Supervision, Writing – review and editing. **Peter Halley:** Resources, Supervision, Writing – review and editing. **Luigi-Jules Vandi:** Formal analysis, Funding acquisition, Methodology, Project administration, Resources, Supervision, Writing – review and editing.
A
Funding
This work was financially supported by Queensland Government through the Advance Queensland Industry Research Fellowships Scheme (Grant No. AQIRF065-2019RD2).
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Data Availability
The data that support the findings of this study are available from the corresponding author upon reasonable request.
Electronic Supplementary Material
Below is the link to the electronic supplementary material
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