Anomalous Mechanical Response of Stretched Film of Poly(3-Hydroxybutyrate-co-3-Hydroxyhexanoate)

Anomalous Mechanical Response of Stretched Film of Poly(3-Hydroxybutyrate-co-3-Hydroxyhexanoate) | Research Square window.SnipcartSettings = { analytics: { enabled: false } }; (function() { var accessVector = localStorage.getItem('access_vector') || ''; window.dataLayer = window.dataLayer || []; if (accessVector) { window.dataLayer.push({ user: { profile: { profileInfo: { snid: accessVector } } } }); } })(); (function(w,d,s,l,i){w[l]=w[l]||[];w[l].push({'gtm.start':new Date().getTime(),event:'gtm.js'});var f=d.getElementsByTagName(s)[0],j=d.createElement(s),dl=l!='dataLayer'?'&l='+l:'';j.async=true;j.src='https://www.googletagmanager.com/gtm.js?id='+i+dl;f.parentNode.insertBefore(j,f);})(window,document,'script','dataLayer','GTM-K279D39R'); Browse Preprints In Review Journals COVID-19 Preprints AJE Video Bytes Research Tools Research Promotion AJE Professional Editing AJE Rubriq About Preprint Platform In Review Editorial Policies Our Team Advisory Board Help Center Sign In Submit a Preprint Cite Share Download PDF Research Article Anomalous Mechanical Response of Stretched Film of Poly(3-Hydroxybutyrate-co-3-Hydroxyhexanoate) Yuta Fukuda, Khunanya Janchai, Takenobu Sunagawa, Masayuki Yamaguchi This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-4463452/v1 This work is licensed under a CC BY 4.0 License Status: Published Journal Publication published 05 Aug, 2024 Read the published version in Journal of Polymers and the Environment → Version 1 posted 11 You are reading this latest preprint version Abstract The mechanical responses during loading, unloading, and reloading cyclic tensile tests of a tubular blown film of poly(3-hydroxybutyrate- co -3-hydroxyhexanoate) are studied. Although the stress–strain curve recorded during the initial stretching process is typical for a crystalline polymer, the stretched film behaves like a rubber during the reloading process; that is, low modulus with a small residual strain after unloading. Furthermore, the stress–strain curves during the reloading process are an inverted “S” shape. During the first stretching process of the polymer film, small crystals are destroyed without reorganization into a crystalline structure, leading to the observed decrease of crystallinity. In contrast, well-developed crystals that orient to the machine direction of the film do not disappear during the first stretching and act as crosslink points during reloading. As a result, a rubber-like response is detected. This mechanical response during reloading is considerably different from those of conventional crystalline plastics such as polyethylene and polypropylene. Poly(3-hydroxybutyrate-co-3-hydroxyhexanoate) Mechanical property Tensile test Residual strain Figures Figure 1 Figure 2 Figure 3 Figure 4 Figure 5 Figure 6 Figure 7 Figure 8 Figure 9 Figure 10 Figure 11 Introduction Poly(3-hydroxybutyrate) (PHB) and its copolymers show good biodegradability even in marine environments [ 1 – 3 ], which makes them attractive to replace conventional plastics; e.g., isotactic polypropylene (PP). To date, various basic properties of PHB and its copolymers have been clarified, such as rheological properties in the molten state [ 4 – 6 ], thermal properties including crystallization behavior [ 1 , 2 , 7 – 10 ], processability [ 2 , 11 , 12 ], mechanical properties in the solid state [ 1 , 2 , 13 – 17 ], structure and properties of blends and composites with other materials [ 15 , 18 – 23 ], and degradation behavior [ 1 – 3 , 24 , 25 ]. The melting point ( T m ) of PHB is approximately 180°C and decreases with increasing comonomer content [ 1 , 2 ]. The glass transition temperature ( T g ) of PHB is slightly lower than room temperature [ 1 , 2 ]. Therefore, the mechanical toughness of PHB is usually poor especially at high strain rates, where craze formation initiates failure [ 13 – 17 ]. Various methods have been proposed to improve the mechanical toughness of PHB, such as adding a high-molecular-weight fraction [ 14 ] and mixing with rubbery materials [ 15 , 20 ]. Copolymers of PHB show ductile behavior to some degree because their crystallinity decreases with increasing comonomer content. In fact, copolymers with no or low crystallinity behave like a rubber at room temperature [ 19 ]. These properties of PHB including T g and T m are similar to those of PP. Therefore, the mechanical properties of PHB and its copolymers are often compared with those of PP [ 17 ]. The marked difference between PHB and PP is their crystallization rate; PHB shows very slow crystallization compared with PP [ 7 ]. For the industrial application of PHB and its copolymers, their crystallization rate needs to be increased by doping with nucleating agents and/or plasticizers [ 8 ]. Flow-induced crystallization [ 9 , 10 ] and the melt memory effect [ 10 ] can also be used to enhance the crystallization rate of polymers. Slow crystallization of PHB has also been detected during post-processing annealing [ 19 , 26 ]. Consequently, the structure and properties of PHB and its copolymers depend on their storage conditions after processing, which is also the case for PP. Another important but unfavorable property of PHB is its severe thermal degradation [ 25 ]. To avoid molecular scission during processing, the resin temperature must be low. Despite these limitations, some copolymers of PHB have been successfully commercialized. One example is poly(3-hydroxybutyrate- co -3-hydroxyhexanoate) (PHBHHx), which has been already used for various applications [ 20 , 27 , 28 ]. Here, we investigated the mechanical properties of stretched films of PHBHHx, which differ considerably from those of stretched PP films. In general, polymer solids show complicated mechanical responses including viscoelastic properties and plasticity. In particular, plastic deformation strongly affects mechanical toughness. Therefore, various approaches have been developed to clarify the plastic deformation behavior of polymers. One of the phenomenological methods to characterize plastic deformation is loading, unloading, and reloading cyclic tests [ 29 – 41 ], which are also used to investigate the Bauschinger effect [ 30 ]. In glassy plastics, the residual strain after an unloading process, i.e., plastic strain, is proportional to the applied strain beyond the yield point [ 29 , 34 , 37 ]. Similar behavior has been reported for crystalline polymers with T g below ambient temperature, such as PP [ 29 , 31 – 33 , 35 , 36 , 39 – 41 ]. Therefore, the residual strain increases monotonically with the pre-strain level. Beyond the yield point, the initial modulus upon reloading increases in general because of the molecular orientation present from the stretching history. The initial modulus upon reloading is not affected by the magnitude of pre-strain when the applied pre-strain is in the neck-formation region of the original stress–strain curve. In this study, we investigated the mechanical responses of a tubular blown film of PHBHHx during cyclic tensile deformations and its structure before and after testing were characterized. Experimental Procedure Material and sample preparation The PHBHHx (Kaneka Corporation, Tokyo, Japan) was composed of two fractions; one contained 28 mol.% of 3-hydroxyhexanoate (HHx) and the other contained 2 mol.% of HHx. Because both fractions had the same weight ratio, i.e., 50 wt.%, the average HHx content of PHBHHx was 15 mol.%, which was confirmed by gas chromatography measurement after methanolytic degradation [ 42 ]. The molecular weight and distribution of PHBHHx were characterized by size-exclusion chromatography (SEC; TSK-GEL GMHXL 16141; Tosoh, Tokyo, Japan) at 40°C using chloroform as a solvent. The number- and weight-average molecular weights of PHBHHx were M n = 380,000 and M w = 690,000 (Da), respectively, which were measured using a polystyrene standard. The SEC curve of PHBHHx is shown in Fig. 1 . Both fractions had a similar molecular weight and distribution, so the curve was narrow with M w / M n of 1.8. After drying the PHBHHx pellets, a blown film was prepared by tubular extrusion using a single-screw extruder (Hokushin Sangyo, Furukawa, Japan) at a screw rotation speed of 30 rpm. The out-put rate was 25 kg h − 1 . The diameter of die was 100 mm and the die gap was 1 mm. Therefore, the shear rate at the die wall \(\dot {\gamma }\) , calculated by Eq. (1) [ 43 ], was 40 s − 1 , assuming that the melt density was 900 kg m − 3 [ 44 ]. The temperature at the screw top and the die was controlled at 165°C, and the air temperature was 35°C. The take-up speed was 7.5 m min − 1 and the blow-up ratio was 2.55. Measurements Various mechanical tests were performed using a tensile testing machine (EZ-LX HS; Shimadzu, Kyoto, Japan) at 25°C. Rectangular samples with a width of 2 mm were cut from the blown film. The initial distance between the clamps was 8 mm. The whole stress–strain curves were measured in both the machine direction (MD) and transverse direction (TD). One of the clamps was moved at a constant speed of 1.6 mm s –1 . Therefore, the initial strain rate was 0.2 s –1 . Measurements were carried out five times and the average values of initial modulus, yield stress, ultimate stress, and elongation at break were calculated along with their standard deviations. For the multiple-cycle deformation tests, the measurements were performed under MD stretching at a constant crosshead speed of 0.16 mm s –1 during loading, unloading, and reloading processes. The first loading was performed at a gauge length of 16 mm, i.e., engineering strain of 2, which was immediately followed by the unloading process until the load was 0 N. Then, reloading was immediately performed until the gauge length was 16 mm longer than the minimum gauge length of the previous unloading process. For the single-cycle deformation tests, the clamp was moved at 0.16, 0.8, and 1.6 mm s –1 during all processes; i.e., the initial strain rates were 0.02, 0.1, and 0.2 s –1 , respectively. First, the sample was stretched to a gauge length of 32 mm, i.e., engineering strain of 4, in the MD. Then, the crosshead was immediately moved in the opposite direction, i.e., unloading, until the load became 0 N, followed by the reloading process. Structural characterization was performed using the original blown film and the necked region of the stretched film, which is denoted as the stretched film. To obtain the stretched film, the original film was stretched in the MD at a stretching speed of 1.6 mm s –1 (strain rate of 0.2 s –1 ). The initial distance between clamps was 8 mm and the final distance was 32 mm. After removal of the sample from the tensile testing machine, the necked region cut out for subsequent measurements. The optical retardation of the original and stretched films was measured by an optical birefringence analyzer (KOBRA-WPR; Oji Scientific Instruments, Amagasaki, Japan) at a wavelength of 589 nm, i.e., the Fraunhofer D-line. Two-dimensional wide-angle X-ray diffraction (2D-WAXD) patterns of the original and stretched films were obtained using an X-ray diffractometer (SmartLab; Rigaku, Tokyo, Japan) with a camera length of 45 mm. The diffractometer was operated using CuKα radiation (45 kV and 200 mA) with a charge-coupled device detector (HyPix-400; Rigaku). The exposure time was 15 min. The data were converted into 2 θ profiles by integrating the intensity along 2 θ from 10° to 30°. The 2D-WAXD crystallinity χ XRD was calculated by the following equation: $${\chi _{XRD}}=\frac{{\sum\limits_{i} {{A_{{c_i}}}} }}{{\sum\limits_{i} {{A_{{c_i}}}} +{A_a}}}$$ 2 , where A ci and A a are the integrated areas of the crystalline diffraction peaks and amorphous halo, respectively. The values of A c and A a were determined by fitting the sum of the Gaussian functions to the 2D-WAXD profile along 2 θ . Thermal properties of films were evaluated by differential scanning calorimetry (DSC) using a calorimeter (DSC8500; PerkinElmer, Waltham, MA, USA). The sample films encapsulated in an aluminum pan were heated from 20 to 165°C at a heating rate of 10°C min –1 under a nitrogen atmosphere. The DSC crystallinity c DSC was calculated by the following equation: $${\chi _{DSC}}=\frac{{\Delta {h_F}}}{{\Delta {h_F}^{0}}}$$ 3 where Δh F is the heat of fusion and Δh F 0 is the heat of fusion of a perfect crystal of PHB, which was reported to be 146 J g –1 [ 45 ]. The temperature dependence of the dynamic tensile moduli of the original and stretched films was measured at 10 Hz using a dynamic mechanical analyzer (Rheogel-E4000; UBM, Muko, Japan) over the temperature range from − 100 to 180°C at a constant heating rate of 2°C min –1 . The gauge length was 10 mm and width was 5 mm. The oscillatory strain was applied in the MD. Results and Discussion The whole stress–strain curves measured during MD and TD stretching are shown in Fig. 2 . Both stress and strain were the engineering values. The initial distance between the clamps was 8 mm and the crosshead speed was 1.6 mm s –1 (strain rate of 0.2 s –1 ). Figure 2 reveals that both curves were composed of elastic deformation, yielding, neck-propagation, and strain-hardening regions. This is typical behavior for semi-crystalline polymers. The Young’s modulus and stress values of the film that underwent MD stretching were higher than those of the film subjected to TD stretching. The average values of Young’s modulus, yield stress, strain at break, and ultimate stress along with their standard deviations are summarized in Table 1 . Table 1 Tensile properties of PHBHHx films stretched in machine and transverse directions Young’s modulus (MPa) Yield stress (MPa) Strain at break (-) Ultimate stress (MPa) MD 211 (13) 12.0 (0.8) 8.9 (0.5) 21.5 (1.2) TD 206 (22) 10.4 (0.1) 9.5 (0.8) 18.7 (1.5) * Standard deviations are shown in parentheses. Figure 3 shows the stress–strain curves of the films during multiple-cycle deformation. The crosshead speed during all processes was 0.16 mm s –1 and the initial distance between the clamps was 8 mm. Each reloading process was carried out until the gauge length became 16 mm longer than the final one (where the load became zero) of the previous unloading process. As shown in Fig. 3 , the residual strain after unloading processes and initial modulus of the reloading processes decreased as the number of cycles increased. As a result, almost the same curves were obtained after ten cycles. The stress–strain curve eventually became an inverted “S” shape, which was resembling to the stress–strain curve expressed by the non-Gaussian statistical model of rubber elasticity [ 46 , 47 ]. This result differed markedly from those reported previously for various plastics [ 29 , 31 – 37 , 39 – 41 ]. The present result suggested that rubber-like deformation occurred during the reloading process. It was also demonstrated that stress/strain-induced crystallization barely occurred for PHBHHx at 25°C, which is quite different from the case for conventional crystalline polymers such as polyethylene and PP [ 48 – 51 ]. Figure 4 shows the results of single-cycle deformation tests at strain rates of 0.02, 0.1, and 0.2 s –1 , which correspond to crosshead speeds of 0.16, 0.8, and 0.16 mm s –1 , respectively. The unloading process started at an engineering strain of 4. The yield stress during the first loading process decreased at a low strain rate; i.e., 0.02 s –1 . The reloading curves were almost independent of the applied strain rate; i.e., the stretched film showed rubber-like behavior. Based on the experimental results, we then prepared stretched films to examine their mechanical responses during reloading. The original blown film was stretched at 0.2 s –1 in the MD with an engineering strain of 4 and then removed from the tensile machine to give the stretched film. Figure 5 shows the 2D-WAXD profiles of the original and stretched films. The 2 θ profiles of the original and stretched films are presented in Fig. 6 (a). Some diffraction peaks attributed to α-crystals were observed. In this figure, the Millar indices are denoted in parentheses [ 52 ]. The χ XRD values calculated using Eq. ( 2 ) were 0.42 for the original film and 0.33 for the stretched film. These results indicated that the crystallinity of the film was decreased by the stretching process. The azimuthal angle distributions of the (020) plane of crystallites in the original and stretched films are shown in Fig. 6 (b). Intense diffraction peaks were detected on the equator, suggesting that the α-crystals were orientated along the MD. This orientation must be induced by the flow history during film processing. It should be noted that a similar distribution was detected for the stretched film, demonstrating that some oriented crystals were not destroyed by the stretching treatment. The birefringence of the original and stretched films was evaluated from the optical retardation and film thickness. The birefringence at 589 nm was − 2.5 × 10 –4 for the original film and − 2.6 × 10 –4 for the stretched film. Because α-crystals of PHB show negative birefringence [ 53 ], this was a reasonable result and corresponded to the azimuthal angle distribution in Fig. 6 (b). The DSC heating curves of the original and stretched films are depicted in Fig. 7 . Although both films showed almost the same T m of 154°C, their heat of fusion, i.e., crystallinity, was different. The χ DSC values calculated using Eq. ( 3 ) were 0.19 for the original film and 0.17 for the stretched film. Figure 8 presents the temperature dependence of the tensile storage modulus E’ and loss modulus E” of the original and stretched films when oscillatory strain was applied in the MD. Both films exhibited typical viscoelastic properties of crystalline polymers. Although the E’ values of both films were almost the same below T g , those of the stretched film decreased greatly beyond T g . The E’’ peak ascribed to T g was located at approximately − 10°C. The E’’ peak of the stretched film was sharper than that of the original film, which was attributed to the decrease in crystallinity after stretching. These results demonstrated that stretching decreased the crystallinity of the polymer film, as revealed by the 2D-WAXD profiles. Therefore, the stretched film exhibited a lower modulus than that of the original film, as shown in Figs. 3 and 4 . Considering that the modulus decrease was detected only in a relatively low temperature region, i.e., from 0 to 80°C, only small crystals with low T m must be destroyed by the stretching process; these small crystals do not reorganize into well-oriented crystals with high T m after stretching. This result should be noted because conventional crystalline plastics including PP show higher modulus after stretching because of the stress/strain-induced crystallization in the solid state. The DSC results (Fig. 8 ) also indicated that well-developed crystals with relatively high T m were not destroyed by the stretching process. The optical retardation during the reloading process of the stretched film was measured. As shown in Fig. 9 , some strains were applied to the stretched films by the tensile machine, and then images under crossed polars with/without a quarter-wave plate were captured. The slow axis of the quarter-wave plate was oriented along the MD and the polarizer and analyzer were tilted at 45° with respect to the MD, as illustrated in Fig. 9 . As mentioned, the stretched film showed negative birefringence at a strain ( ε ) of 0. Then, the birefringence became almost 0 at approximately ε = 0.3. Finally, positive birefringence was detected under further stretching, e.g., ε = 0.7, demonstrating that the sign of the birefringence changed from negative to positive during the reloading process. We also evaluated the 2D-WAXD profiles at the same strain levels as those used in Fig. 9 to reveal the deformation behavior of the stretched film. Figure 10 shows the 2 θ profiles and azimuthal angle distribution of the (020) plane for the stretched film at different ε . The crystallinity and crystallite orientation were not affected during the reloading process, i.e., the α-crystals oriented to the MD even after stretching. Considering that a simple addition rule is applicable for birefringence [ 46 , 54 , 55 ], the growth of positive birefringence must be attributed to the orientation of amorphous tie chains under the applied ε . Figure 11 shows a schematic illustration of crystals and amorphous tie chains in the stretched film during tensile deformation. As revealed in the azimuthal angle distribution of the 2D-WAXD profile, the crystals oriented to the MD with no/low orientation of amorphous chains before reloading. Therefore, the birefringence of the stretched film without applied stress was negative. During deformation, i.e., the reloading process, the amorphous tie chains were orientated along MD while keeping the orientation of α-crystals. Considering that the stretched film showed rubber-like behavior with a small residual strain, amorphous tie chains between crystals must be responsible for the mechanical response during reloading. Conclusion The mechanical properties of a tubular blown PHBHHx film were studied by cyclic tensile tests. During the first stretching process, the film showed a typical stress–strain curve for a crystalline polymer. During stretching, the film lost small crystals with low T m and did not show stress/strain-induced crystallization in the solid state. As a result, the crystallinity and thus the modulus of the PHBHHx film were decreased by the stretching history. Moreover, well-developed crystals oriented in the MD were not destroyed during stretching. These crystals acted as crosslink points during the reloading process. Therefore, the stress–strain curve during reloading was an inverted “S” shape, and the residual strain after unloading was minimal. During the reloading process, the amorphous chains connecting crystallites oriented, which was supported by the results of birefringence measurements. Declarations Acknowledgment We thank Natasha Lundin, PhD, from Edanz (https://jp.edanz.com/ac) for editing a draft of this manuscript. Funding This work was supported by JST SPRING, Japan Grant Number JPMJSP2102. Author Contributions Data curation and Investigation, Y.F.; Methodology, Y.F. and K.J.; Writing Original Draft, Y.F.; Samples, T.S.; Writing Review and Edit, M.Y. All authors read and approved the final manuscript. Conflict of interest: The authors declare no conflict of interest. References Doi Y (1990) Microbial Polyesters. VCH Publishers, New York Doi Y, Steinbüchel A Editors (2001) Biopolymers: Polyesters I and II Vol 3. 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Cite Share Download PDF Status: Published Journal Publication published 05 Aug, 2024 Read the published version in Journal of Polymers and the Environment → Version 1 posted Editorial decision: Revision requested 06 Jul, 2024 Reviews received at journal 19 Jun, 2024 Reviewers agreed at journal 14 Jun, 2024 Reviewers agreed at journal 13 Jun, 2024 Reviews received at journal 05 Jun, 2024 Reviewers agreed at journal 28 May, 2024 Reviewers agreed at journal 26 May, 2024 Reviewers invited by journal 26 May, 2024 Submission checks completed at journal 24 May, 2024 Editor assigned by journal 24 May, 2024 First submitted to journal 22 May, 2024 You are reading this latest preprint version Research Square lets you share your work early, gain feedback from the community, and start making changes to your manuscript prior to peer review in a journal. As a division of Research Square Company, we’re committed to making research communication faster, fairer, and more useful. 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Also discoverable on Platform About Our Team In Review Editorial Policies Advisory Board Help Center Resources Author Services Accessibility API Access RSS feed Manage Cookie Preferences © Research Square 2026 | ISSN 2693-5015 (online) Privacy Policy Terms of Service Do Not Sell My Personal Information {"props":{"pageProps":{"initialData":{"identity":"rs-4463452","acceptedTermsAndConditions":true,"allowDirectSubmit":false,"archivedVersions":[],"articleType":"Research Article","associatedPublications":[],"authors":[{"id":309965009,"identity":"cf00010b-37af-41a9-8e3c-2e25495e6712","order_by":0,"name":"Yuta Fukuda","email":"","orcid":"","institution":"Japan Advanced Institute of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Yuta","middleName":"","lastName":"Fukuda","suffix":""},{"id":309965011,"identity":"f73f68da-7838-469c-9e00-f85846a1b158","order_by":1,"name":"Khunanya Janchai","email":"","orcid":"","institution":"Japan Advanced Institute of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Khunanya","middleName":"","lastName":"Janchai","suffix":""},{"id":309965013,"identity":"f9006f41-e554-4181-b4de-20e1e5e1c5a9","order_by":2,"name":"Takenobu Sunagawa","email":"","orcid":"","institution":"Green Planet Technology Laboratories, Kaneka Corporation","correspondingAuthor":false,"prefix":"","firstName":"Takenobu","middleName":"","lastName":"Sunagawa","suffix":""},{"id":309965015,"identity":"194be1a2-c57f-45f1-9847-bb146e9b84e0","order_by":3,"name":"Masayuki Yamaguchi","email":"data:image/png;base64,iVBORw0KGgoAAAANSUhEUgAAAZAAAAAyAQMAAABI0h/eAAAABlBMVEX///8AAABVwtN+AAAACXBIWXMAAA7EAAAOxAGVKw4bAAAA0UlEQVRIiWNgGAWjYPACG2YGZiiTsYE4LWkgLRDFxGo5TIJi3QYew083as6z87ezP3/AUGPHwDybgE6zAzzG0jnHbjNLHOYxbGA4lszAOOcAQS0G0jlst5kZDvMAHcZ2gIFxRgJhW37n/DvHLH+Y/WEDwz/itJhJ57YdYDY4zGDYwNhGlBa2MuvcvmRmQ6BfZiT2JfMQ4Rfmzbdzvtkly50//uDDh292coaEQoxB/oUBiEoGc4BOAtpFQAcDA/sDEGmHMEOCoJZRMApGwSgYYQAAXU9AlB25omgAAAAASUVORK5CYII=","orcid":"","institution":"Japan Advanced Institute of Science and Technology","correspondingAuthor":true,"prefix":"","firstName":"Masayuki","middleName":"","lastName":"Yamaguchi","suffix":""}],"badges":[],"createdAt":"2024-05-23 01:14:22","currentVersionCode":1,"declarations":"","doi":"10.21203/rs.3.rs-4463452/v1","doiUrl":"https://doi.org/10.21203/rs.3.rs-4463452/v1","draftVersion":[],"editorialEvents":[{"content":"https://doi.org/10.1007/s10924-024-03370-x","type":"published","date":"2024-08-05T15:58:13+00:00"}],"editorialNote":"","failedWorkflow":false,"files":[{"id":57787735,"identity":"59528c07-52ab-4fe4-ba95-2b2ab6f76cf0","added_by":"auto","created_at":"2024-06-05 16:39:28","extension":"png","order_by":1,"title":"Figure 1","display":"","copyAsset":false,"role":"figure","size":6355,"visible":true,"origin":"","legend":"\u003cp\u003eMolecular weight distribution of PHBHHx measured using a polystyrene standard.\u003c/p\u003e","description":"","filename":"1.png","url":"https://assets-eu.researchsquare.com/files/rs-4463452/v1/99065d1783f7f0a05da1b4be.png"},{"id":57787739,"identity":"6fdba39f-c640-4788-a111-b24e105fe7a9","added_by":"auto","created_at":"2024-06-05 16:39:28","extension":"png","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":13957,"visible":true,"origin":"","legend":"\u003cp\u003eStress–strain curves of PHBHHx films during MD and TD stretching at 25 °C; (a) magnified view at low strain and (b) whole curves.\u003c/p\u003e","description":"","filename":"2.png","url":"https://assets-eu.researchsquare.com/files/rs-4463452/v1/7dcd315497684d0d16bd0908.png"},{"id":57787736,"identity":"2c6904f0-73cb-414a-93ac-e02119e2964e","added_by":"auto","created_at":"2024-06-05 16:39:28","extension":"png","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":17191,"visible":true,"origin":"","legend":"\u003cp\u003eStress–strain curves of a PHBHHxfilm during multiple-cycle deformation testing at 25 °C.\u003c/p\u003e","description":"","filename":"3.png","url":"https://assets-eu.researchsquare.com/files/rs-4463452/v1/e472fed529acf99acc087a2e.png"},{"id":57787741,"identity":"3e436b76-2c5d-4533-989b-a0f72b56514e","added_by":"auto","created_at":"2024-06-05 16:39:28","extension":"png","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":14001,"visible":true,"origin":"","legend":"\u003cp\u003eStress–strain curves of a PHBHHxfilm during single-cycle deformation tests at different strain rates at 25 °C.\u003c/p\u003e","description":"","filename":"4.png","url":"https://assets-eu.researchsquare.com/files/rs-4463452/v1/f41a193aceab5bc0aa589975.png"},{"id":57787737,"identity":"402e5025-5bd9-4a8d-aefa-efc4c273e035","added_by":"auto","created_at":"2024-06-05 16:39:28","extension":"png","order_by":5,"title":"Figure 5","display":"","copyAsset":false,"role":"figure","size":237285,"visible":true,"origin":"","legend":"\u003cp\u003e2D-WAXD patterns of (a) original and (b) stretched films.\u003c/p\u003e","description":"","filename":"5.png","url":"https://assets-eu.researchsquare.com/files/rs-4463452/v1/a2bd011e738c23221b087ebb.png"},{"id":57788155,"identity":"57e2d00c-578d-451f-a80f-16e1536144b2","added_by":"auto","created_at":"2024-06-05 16:47:28","extension":"png","order_by":6,"title":"Figure 6","display":"","copyAsset":false,"role":"figure","size":19147,"visible":true,"origin":"","legend":"\u003cp\u003e(a) 2\u003cem\u003eθ\u003c/em\u003e profiles and (b) azimuthal angle distribution of the (020) plane for the original and stretched films.\u003c/p\u003e","description":"","filename":"6.png","url":"https://assets-eu.researchsquare.com/files/rs-4463452/v1/4054aaa0fe8e4f5663e02142.png"},{"id":57787744,"identity":"af25c298-45f9-4df3-8b83-e9a2292f26e6","added_by":"auto","created_at":"2024-06-05 16:39:29","extension":"png","order_by":7,"title":"Figure 7","display":"","copyAsset":false,"role":"figure","size":13406,"visible":true,"origin":"","legend":"\u003cp\u003eDSC heating curves measured at 10 °C min\u003csup\u003e–1\u003c/sup\u003e for the original and stretched films.\u003c/p\u003e","description":"","filename":"7.png","url":"https://assets-eu.researchsquare.com/files/rs-4463452/v1/d78e123aa1f7742367687172.png"},{"id":57788157,"identity":"3cd85adf-ec4c-4415-becb-a9bf71b0f24d","added_by":"auto","created_at":"2024-06-05 16:47:28","extension":"png","order_by":8,"title":"Figure 8","display":"","copyAsset":false,"role":"figure","size":19239,"visible":true,"origin":"","legend":"\u003cp\u003eTemperature dependence of the tensile storage modulus \u003cem\u003eE’\u003c/em\u003e (filled symbols) and loss modulus \u003cem\u003eE”\u003c/em\u003e (open symbols) at 10 Hz for the original (circles) and stretched (diamonds) films.\u003c/p\u003e","description":"","filename":"8.png","url":"https://assets-eu.researchsquare.com/files/rs-4463452/v1/ffdf047aa2b07aad0584c0c4.png"},{"id":57788159,"identity":"440c8504-d6cd-40f1-9b52-459520f5eeb8","added_by":"auto","created_at":"2024-06-05 16:47:29","extension":"png","order_by":9,"title":"Figure 9","display":"","copyAsset":false,"role":"figure","size":129439,"visible":true,"origin":"","legend":"\u003cp\u003e(Left) Schematic diagram of the experimental set-up for birefringence measurement and (Right) photographs under crossed polars with/without a quarter-wave (1/4l) plate for the original film and stretched films under various strains.\u003c/p\u003e","description":"","filename":"9.png","url":"https://assets-eu.researchsquare.com/files/rs-4463452/v1/16a0cdd2029b31944c6e6742.png"},{"id":57788522,"identity":"ccce8eb3-31c6-4f54-abe8-5403b273de54","added_by":"auto","created_at":"2024-06-05 16:55:28","extension":"png","order_by":10,"title":"Figure 10","display":"","copyAsset":false,"role":"figure","size":22363,"visible":true,"origin":"","legend":"\u003cp\u003e(a) 2\u003cem\u003eθ\u003c/em\u003e profiles and (b) azimuthal angle distribution of the (020) plane for stretched films under various strains \u003cem\u003eε\u003c/em\u003e.\u003c/p\u003e","description":"","filename":"10.png","url":"https://assets-eu.researchsquare.com/files/rs-4463452/v1/a26cc28543c35e89c4d9c71c.png"},{"id":57787742,"identity":"c87c4f54-be4b-43db-886a-67d22cfc7389","added_by":"auto","created_at":"2024-06-05 16:39:28","extension":"png","order_by":11,"title":"Figure 11","display":"","copyAsset":false,"role":"figure","size":72230,"visible":true,"origin":"","legend":"\u003cp\u003eSchematic illustration of amorphous tie chains and crystallites in the stretched PHBHHx film during the reloading process.\u003c/p\u003e","description":"","filename":"11.png","url":"https://assets-eu.researchsquare.com/files/rs-4463452/v1/35b0abbbfc6557492a2dbf9f.png"},{"id":62298907,"identity":"54cac99e-38a1-422f-84e5-9073d94387ee","added_by":"auto","created_at":"2024-08-12 16:17:13","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":936655,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-4463452/v1/8ae0ebc3-c049-423d-81a0-2dd95484bc18.pdf"}],"financialInterests":"No competing interests reported.","formattedTitle":"Anomalous Mechanical Response of Stretched Film of Poly(3-Hydroxybutyrate-co-3-Hydroxyhexanoate)","fulltext":[{"header":"Introduction","content":"\u003cp\u003ePoly(3-hydroxybutyrate) (PHB) and its copolymers show good biodegradability even in marine environments [\u003cspan additionalcitationids=\"CR2\" citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e], which makes them attractive to replace conventional plastics; e.g., isotactic polypropylene (PP). To date, various basic properties of PHB and its copolymers have been clarified, such as rheological properties in the molten state [\u003cspan additionalcitationids=\"CR5\" citationid=\"CR4\" class=\"CitationRef\"\u003e4\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR6\" class=\"CitationRef\"\u003e6\u003c/span\u003e], thermal properties including crystallization behavior [\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e, \u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e, \u003cspan additionalcitationids=\"CR8 CR9\" citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e], processability [\u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e, \u003cspan citationid=\"CR11\" class=\"CitationRef\"\u003e11\u003c/span\u003e, \u003cspan citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e], mechanical properties in the solid state [\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e, \u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e, \u003cspan additionalcitationids=\"CR14 CR15 CR16\" citationid=\"CR13\" class=\"CitationRef\"\u003e13\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e], structure and properties of blends and composites with other materials [\u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e, \u003cspan additionalcitationids=\"CR19 CR20 CR21 CR22\" citationid=\"CR18\" class=\"CitationRef\"\u003e18\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR23\" class=\"CitationRef\"\u003e23\u003c/span\u003e], and degradation behavior [\u003cspan additionalcitationids=\"CR2\" citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e, \u003cspan citationid=\"CR24\" class=\"CitationRef\"\u003e24\u003c/span\u003e, \u003cspan citationid=\"CR25\" class=\"CitationRef\"\u003e25\u003c/span\u003e]. The melting point (\u003cem\u003eT\u003c/em\u003e\u003csub\u003e\u003cem\u003em\u003c/em\u003e\u003c/sub\u003e) of PHB is approximately 180\u0026deg;C and decreases with increasing comonomer content [\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e, \u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e]. The glass transition temperature (\u003cem\u003eT\u003c/em\u003e\u003csub\u003e\u003cem\u003eg\u003c/em\u003e\u003c/sub\u003e) of PHB is slightly lower than room temperature [\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e, \u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e]. Therefore, the mechanical toughness of PHB is usually poor especially at high strain rates, where craze formation initiates failure [\u003cspan additionalcitationids=\"CR14 CR15 CR16\" citationid=\"CR13\" class=\"CitationRef\"\u003e13\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e]. Various methods have been proposed to improve the mechanical toughness of PHB, such as adding a high-molecular-weight fraction [\u003cspan citationid=\"CR14\" class=\"CitationRef\"\u003e14\u003c/span\u003e] and mixing with rubbery materials [\u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e, \u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e]. Copolymers of PHB show ductile behavior to some degree because their crystallinity decreases with increasing comonomer content. In fact, copolymers with no or low crystallinity behave like a rubber at room temperature [\u003cspan citationid=\"CR19\" class=\"CitationRef\"\u003e19\u003c/span\u003e]. These properties of PHB including \u003cem\u003eT\u003c/em\u003e\u003csub\u003e\u003cem\u003eg\u003c/em\u003e\u003c/sub\u003e and \u003cem\u003eT\u003c/em\u003e\u003csub\u003e\u003cem\u003em\u003c/em\u003e\u003c/sub\u003e are similar to those of PP. Therefore, the mechanical properties of PHB and its copolymers are often compared with those of PP [\u003cspan citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e]. The marked difference between PHB and PP is their crystallization rate; PHB shows very slow crystallization compared with PP [\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e]. For the industrial application of PHB and its copolymers, their crystallization rate needs to be increased by doping with nucleating agents and/or plasticizers [\u003cspan citationid=\"CR8\" class=\"CitationRef\"\u003e8\u003c/span\u003e]. Flow-induced crystallization [\u003cspan citationid=\"CR9\" class=\"CitationRef\"\u003e9\u003c/span\u003e, \u003cspan citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e] and the melt memory effect [\u003cspan citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e] can also be used to enhance the crystallization rate of polymers. Slow crystallization of PHB has also been detected during post-processing annealing [\u003cspan citationid=\"CR19\" class=\"CitationRef\"\u003e19\u003c/span\u003e, \u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e]. Consequently, the structure and properties of PHB and its copolymers depend on their storage conditions after processing, which is also the case for PP. Another important but unfavorable property of PHB is its severe thermal degradation [\u003cspan citationid=\"CR25\" class=\"CitationRef\"\u003e25\u003c/span\u003e]. To avoid molecular scission during processing, the resin temperature must be low. Despite these limitations, some copolymers of PHB have been successfully commercialized. One example is poly(3-hydroxybutyrate-\u003cem\u003eco\u003c/em\u003e-3-hydroxyhexanoate) (PHBHHx), which has been already used for various applications [\u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e, \u003cspan citationid=\"CR27\" class=\"CitationRef\"\u003e27\u003c/span\u003e, \u003cspan citationid=\"CR28\" class=\"CitationRef\"\u003e28\u003c/span\u003e]. Here, we investigated the mechanical properties of stretched films of PHBHHx, which differ considerably from those of stretched PP films.\u003c/p\u003e \u003cp\u003eIn general, polymer solids show complicated mechanical responses including viscoelastic properties and plasticity. In particular, plastic deformation strongly affects mechanical toughness. Therefore, various approaches have been developed to clarify the plastic deformation behavior of polymers. One of the phenomenological methods to characterize plastic deformation is loading, unloading, and reloading cyclic tests [\u003cspan additionalcitationids=\"CR30 CR31 CR32 CR33 CR34 CR35 CR36 CR37 CR38 CR39 CR40\" citationid=\"CR29\" class=\"CitationRef\"\u003e29\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR41\" class=\"CitationRef\"\u003e41\u003c/span\u003e], which are also used to investigate the Bauschinger effect [\u003cspan citationid=\"CR30\" class=\"CitationRef\"\u003e30\u003c/span\u003e]. In glassy plastics, the residual strain after an unloading process, i.e., plastic strain, is proportional to the applied strain beyond the yield point [\u003cspan citationid=\"CR29\" class=\"CitationRef\"\u003e29\u003c/span\u003e, \u003cspan citationid=\"CR34\" class=\"CitationRef\"\u003e34\u003c/span\u003e, \u003cspan citationid=\"CR37\" class=\"CitationRef\"\u003e37\u003c/span\u003e]. Similar behavior has been reported for crystalline polymers with \u003cem\u003eT\u003c/em\u003e\u003csub\u003e\u003cem\u003eg\u003c/em\u003e\u003c/sub\u003e below ambient temperature, such as PP [\u003cspan citationid=\"CR29\" class=\"CitationRef\"\u003e29\u003c/span\u003e, \u003cspan additionalcitationids=\"CR32\" citationid=\"CR31\" class=\"CitationRef\"\u003e31\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR33\" class=\"CitationRef\"\u003e33\u003c/span\u003e, \u003cspan citationid=\"CR35\" class=\"CitationRef\"\u003e35\u003c/span\u003e, \u003cspan citationid=\"CR36\" class=\"CitationRef\"\u003e36\u003c/span\u003e, \u003cspan additionalcitationids=\"CR40\" citationid=\"CR39\" class=\"CitationRef\"\u003e39\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR41\" class=\"CitationRef\"\u003e41\u003c/span\u003e]. Therefore, the residual strain increases monotonically with the pre-strain level. Beyond the yield point, the initial modulus upon reloading increases in general because of the molecular orientation present from the stretching history. The initial modulus upon reloading is not affected by the magnitude of pre-strain when the applied pre-strain is in the neck-formation region of the original stress\u0026ndash;strain curve. In this study, we investigated the mechanical responses of a tubular blown film of PHBHHx during cyclic tensile deformations and its structure before and after testing were characterized.\u003c/p\u003e"},{"header":"Experimental Procedure","content":"\u003cdiv id=\"Sec3\" class=\"Section2\"\u003e \u003ch2\u003eMaterial and sample preparation\u003c/h2\u003e \u003cp\u003eThe PHBHHx (Kaneka Corporation, Tokyo, Japan) was composed of two fractions; one contained 28 mol.% of 3-hydroxyhexanoate (HHx) and the other contained 2 mol.% of HHx. Because both fractions had the same weight ratio, i.e., 50 wt.%, the average HHx content of PHBHHx was 15 mol.%, which was confirmed by gas chromatography measurement after methanolytic degradation [\u003cspan citationid=\"CR42\" class=\"CitationRef\"\u003e42\u003c/span\u003e]. The molecular weight and distribution of PHBHHx were characterized by size-exclusion chromatography (SEC; TSK-GEL GMHXL 16141; Tosoh, Tokyo, Japan) at 40\u0026deg;C using chloroform as a solvent. The number- and weight-average molecular weights of PHBHHx were \u003cem\u003eM\u003c/em\u003e\u003csub\u003e\u003cem\u003en\u003c/em\u003e\u003c/sub\u003e = 380,000 and \u003cem\u003eM\u003c/em\u003e\u003csub\u003e\u003cem\u003ew\u003c/em\u003e\u003c/sub\u003e = 690,000 (Da), respectively, which were measured using a polystyrene standard. The SEC curve of PHBHHx is shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003e. Both fractions had a similar molecular weight and distribution, so the curve was narrow with \u003cem\u003eM\u003c/em\u003e\u003csub\u003e\u003cem\u003ew\u003c/em\u003e\u003c/sub\u003e/\u003cem\u003eM\u003c/em\u003e\u003csub\u003e\u003cem\u003en\u003c/em\u003e\u003c/sub\u003e of 1.8.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eAfter drying the PHBHHx pellets, a blown film was prepared by tubular extrusion using a single-screw extruder (Hokushin Sangyo, Furukawa, Japan) at a screw rotation speed of 30 rpm. The out-put rate was 25 kg h\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e. The diameter of die was 100 mm and the die gap was 1 mm. Therefore, the shear rate at the die wall \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\dot {\\gamma }\\)\u003c/span\u003e\u003c/span\u003e, calculated by Eq.\u0026nbsp;(1) [\u003cspan citationid=\"CR43\" class=\"CitationRef\"\u003e43\u003c/span\u003e], was 40 s\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e, assuming that the melt density was 900 kg m\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e [\u003cspan citationid=\"CR44\" class=\"CitationRef\"\u003e44\u003c/span\u003e]. The temperature at the screw top and the die was controlled at 165\u0026deg;C, and the air temperature was 35\u0026deg;C. The take-up speed was 7.5 m min\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e and the blow-up ratio was 2.55.\u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec4\" class=\"Section2\"\u003e \u003ch2\u003eMeasurements\u003c/h2\u003e \u003cp\u003eVarious mechanical tests were performed using a tensile testing machine (EZ-LX HS; Shimadzu, Kyoto, Japan) at 25\u0026deg;C. Rectangular samples with a width of 2 mm were cut from the blown film. The initial distance between the clamps was 8 mm. The whole stress\u0026ndash;strain curves were measured in both the machine direction (MD) and transverse direction (TD). One of the clamps was moved at a constant speed of 1.6 mm s\u003csup\u003e\u0026ndash;1\u003c/sup\u003e. Therefore, the initial strain rate was 0.2 s\u003csup\u003e\u0026ndash;1\u003c/sup\u003e. Measurements were carried out five times and the average values of initial modulus, yield stress, ultimate stress, and elongation at break were calculated along with their standard deviations.\u003c/p\u003e \u003cp\u003eFor the multiple-cycle deformation tests, the measurements were performed under MD stretching at a constant crosshead speed of 0.16 mm s\u003csup\u003e\u0026ndash;1\u003c/sup\u003e during loading, unloading, and reloading processes. The first loading was performed at a gauge length of 16 mm, i.e., engineering strain of 2, which was immediately followed by the unloading process until the load was 0 N. Then, reloading was immediately performed until the gauge length was 16 mm longer than the minimum gauge length of the previous unloading process.\u003c/p\u003e \u003cp\u003eFor the single-cycle deformation tests, the clamp was moved at 0.16, 0.8, and 1.6 mm s\u003csup\u003e\u0026ndash;1\u003c/sup\u003e during all processes; i.e., the initial strain rates were 0.02, 0.1, and 0.2 s\u003csup\u003e\u0026ndash;1\u003c/sup\u003e, respectively. First, the sample was stretched to a gauge length of 32 mm, i.e., engineering strain of 4, in the MD. Then, the crosshead was immediately moved in the opposite direction, i.e., unloading, until the load became 0 N, followed by the reloading process.\u003c/p\u003e \u003cp\u003eStructural characterization was performed using the original blown film and the necked region of the stretched film, which is denoted as the stretched film. To obtain the stretched film, the original film was stretched in the MD at a stretching speed of 1.6 mm s\u003csup\u003e\u0026ndash;1\u003c/sup\u003e (strain rate of 0.2 s\u003csup\u003e\u0026ndash;1\u003c/sup\u003e). The initial distance between clamps was 8 mm and the final distance was 32 mm. After removal of the sample from the tensile testing machine, the necked region cut out for subsequent measurements.\u003c/p\u003e \u003cp\u003eThe optical retardation of the original and stretched films was measured by an optical birefringence analyzer (KOBRA-WPR; Oji Scientific Instruments, Amagasaki, Japan) at a wavelength of 589 nm, i.e., the Fraunhofer D-line.\u003c/p\u003e \u003cp\u003eTwo-dimensional wide-angle X-ray diffraction (2D-WAXD) patterns of the original and stretched films were obtained using an X-ray diffractometer (SmartLab; Rigaku, Tokyo, Japan) with a camera length of 45 mm. The diffractometer was operated using CuKα radiation (45 kV and 200 mA) with a charge-coupled device detector (HyPix-400; Rigaku). The exposure time was 15 min. The data were converted into 2\u003cem\u003eθ\u003c/em\u003e profiles by integrating the intensity along 2\u003cem\u003eθ\u003c/em\u003e from 10\u0026deg; to 30\u0026deg;. The 2D-WAXD crystallinity \u003cem\u003eχ\u003c/em\u003e\u003csub\u003e\u003cem\u003eXRD\u003c/em\u003e\u003c/sub\u003e was calculated by the following equation:\u003cdiv id=\"Equ1\" class=\"Equation\"\u003e\u003cdiv format=\"TEX\" class=\"mathdisplay\" id=\"FileID_Equ1\" name=\"EquationSource\"\u003e\n$${\\chi _{XRD}}=\\frac{{\\sum\\limits_{i} {{A_{{c_i}}}} }}{{\\sum\\limits_{i} {{A_{{c_i}}}} +{A_a}}}$$\u003c/div\u003e\u003cdiv class=\"EquationNumber\"\u003e2\u003c/div\u003e\u003c/div\u003e,\u003c/p\u003e \u003cp\u003ewhere \u003cem\u003eA\u003c/em\u003e\u003csub\u003e\u003cem\u003eci\u003c/em\u003e\u003c/sub\u003e and \u003cem\u003eA\u003c/em\u003e\u003csub\u003e\u003cem\u003ea\u003c/em\u003e\u003c/sub\u003e are the integrated areas of the crystalline diffraction peaks and amorphous halo, respectively. The values of \u003cem\u003eA\u003c/em\u003e\u003csub\u003e\u003cem\u003ec\u003c/em\u003e\u003c/sub\u003e and \u003cem\u003eA\u003c/em\u003e\u003csub\u003e\u003cem\u003ea\u003c/em\u003e\u003c/sub\u003e were determined by fitting the sum of the Gaussian functions to the 2D-WAXD profile along 2\u003cem\u003eθ\u003c/em\u003e.\u003c/p\u003e \u003cp\u003eThermal properties of films were evaluated by differential scanning calorimetry (DSC) using a calorimeter (DSC8500; PerkinElmer, Waltham, MA, USA). The sample films encapsulated in an aluminum pan were heated from 20 to 165\u0026deg;C at a heating rate of 10\u0026deg;C min\u003csup\u003e\u0026ndash;1\u003c/sup\u003e under a nitrogen atmosphere. The DSC crystallinity \u003cem\u003ec\u003c/em\u003e\u003csub\u003e\u003cem\u003eDSC\u003c/em\u003e\u003c/sub\u003e was calculated by the following equation:\u003cdiv id=\"Equ2\" class=\"Equation\"\u003e\u003cdiv format=\"TEX\" class=\"mathdisplay\" id=\"FileID_Equ2\" name=\"EquationSource\"\u003e\n$${\\chi _{DSC}}=\\frac{{\\Delta {h_F}}}{{\\Delta {h_F}^{0}}}$$\u003c/div\u003e\u003cdiv class=\"EquationNumber\"\u003e3\u003c/div\u003e\u003c/div\u003e\u003c/p\u003e \u003cp\u003ewhere \u003cem\u003eΔh\u003c/em\u003e\u003csub\u003e\u003cem\u003eF\u003c/em\u003e\u003c/sub\u003e is the heat of fusion and \u003cem\u003eΔh\u003c/em\u003e\u003csub\u003e\u003cem\u003eF\u003c/em\u003e\u003c/sub\u003e\u003csup\u003e0\u003c/sup\u003e is the heat of fusion of a perfect crystal of PHB, which was reported to be 146 J g\u003csup\u003e\u0026ndash;1\u003c/sup\u003e [\u003cspan citationid=\"CR45\" class=\"CitationRef\"\u003e45\u003c/span\u003e].\u003c/p\u003e \u003cp\u003eThe temperature dependence of the dynamic tensile moduli of the original and stretched films was measured at 10 Hz using a dynamic mechanical analyzer (Rheogel-E4000; UBM, Muko, Japan) over the temperature range from \u0026minus;\u0026thinsp;100 to 180\u0026deg;C at a constant heating rate of 2\u0026deg;C min\u003csup\u003e\u0026ndash;1\u003c/sup\u003e. The gauge length was 10 mm and width was 5 mm. The oscillatory strain was applied in the MD.\u003c/p\u003e \u003c/div\u003e"},{"header":"Results and Discussion","content":"\u003cp\u003eThe whole stress\u0026ndash;strain curves measured during MD and TD stretching are shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003e. Both stress and strain were the engineering values. The initial distance between the clamps was 8 mm and the crosshead speed was 1.6 mm s\u003csup\u003e\u0026ndash;1\u003c/sup\u003e (strain rate of 0.2 s\u003csup\u003e\u0026ndash;1\u003c/sup\u003e). Figure\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003e reveals that both curves were composed of elastic deformation, yielding, neck-propagation, and strain-hardening regions. This is typical behavior for semi-crystalline polymers. The Young\u0026rsquo;s modulus and stress values of the film that underwent MD stretching were higher than those of the film subjected to TD stretching. The average values of Young\u0026rsquo;s modulus, yield stress, strain at break, and ultimate stress along with their standard deviations are summarized in Table\u0026nbsp;\u003cspan refid=\"Tab1\" class=\"InternalRef\"\u003e1\u003c/span\u003e.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003e \u003cdiv class=\"gridtable\"\u003e\u003ctable float=\"Yes\" id=\"Tab1\" border=\"1\"\u003e \u003ccaption language=\"En\"\u003e \u003cdiv class=\"CaptionNumber\"\u003eTable 1\u003c/div\u003e \u003cdiv class=\"CaptionContent\"\u003e \u003cp\u003eTensile properties of PHBHHx films stretched in machine and transverse directions\u003c/p\u003e \u003c/div\u003e \u003c/caption\u003e \u003ccolgroup cols=\"5\"\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c1\" colnum=\"1\"\u003e\u003c/div\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c2\" colnum=\"2\"\u003e\u003c/div\u003e \u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c3\" colnum=\"3\"\u003e\u003c/div\u003e \u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c4\" colnum=\"4\"\u003e\u003c/div\u003e \u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c5\" colnum=\"5\"\u003e\u003c/div\u003e \u003cthead\u003e \u003ctr\u003e \u003cth align=\"left\" colname=\"c1\"\u003e\u0026nbsp;\u003c/th\u003e \u003cth align=\"left\" colname=\"c2\"\u003e \u003cp\u003eYoung\u0026rsquo;s modulus\u003c/p\u003e \u003cp\u003e(MPa)\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c3\"\u003e \u003cp\u003eYield stress (MPa)\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c4\"\u003e \u003cp\u003eStrain at break (-)\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c5\"\u003e \u003cp\u003eUltimate stress (MPa)\u003c/p\u003e \u003c/th\u003e \u003c/tr\u003e \u003c/thead\u003e \u003ctbody\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eMD\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e211 (13)\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e \u003cp\u003e12.0 (0.8)\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e \u003cp\u003e8.9 (0.5)\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c5\"\u003e \u003cp\u003e21.5 (1.2)\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eTD\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e206 (22)\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e \u003cp\u003e10.4 (0.1)\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e \u003cp\u003e9.5 (0.8)\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c5\"\u003e \u003cp\u003e18.7 (1.5)\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003c/tbody\u003e \u003c/colgroup\u003e \u003c/table\u003e\u003c/div\u003e \u003c/p\u003e \u003cp\u003e* Standard deviations are shown in parentheses.\u003c/p\u003e \u003cp\u003eFigure \u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e shows the stress\u0026ndash;strain curves of the films during multiple-cycle deformation. The crosshead speed during all processes was 0.16 mm s\u003csup\u003e\u0026ndash;1\u003c/sup\u003e and the initial distance between the clamps was 8 mm. Each reloading process was carried out until the gauge length became 16 mm longer than the final one (where the load became zero) of the previous unloading process. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e, the residual strain after unloading processes and initial modulus of the reloading processes decreased as the number of cycles increased. As a result, almost the same curves were obtained after ten cycles. The stress\u0026ndash;strain curve eventually became an inverted \u0026ldquo;S\u0026rdquo; shape, which was resembling to the stress\u0026ndash;strain curve expressed by the non-Gaussian statistical model of rubber elasticity [\u003cspan citationid=\"CR46\" class=\"CitationRef\"\u003e46\u003c/span\u003e, \u003cspan citationid=\"CR47\" class=\"CitationRef\"\u003e47\u003c/span\u003e]. This result differed markedly from those reported previously for various plastics [\u003cspan citationid=\"CR29\" class=\"CitationRef\"\u003e29\u003c/span\u003e, \u003cspan additionalcitationids=\"CR32 CR33 CR34 CR35 CR36\" citationid=\"CR31\" class=\"CitationRef\"\u003e31\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR37\" class=\"CitationRef\"\u003e37\u003c/span\u003e, \u003cspan additionalcitationids=\"CR40\" citationid=\"CR39\" class=\"CitationRef\"\u003e39\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR41\" class=\"CitationRef\"\u003e41\u003c/span\u003e]. The present result suggested that rubber-like deformation occurred during the reloading process. It was also demonstrated that stress/strain-induced crystallization barely occurred for PHBHHx at 25\u0026deg;C, which is quite different from the case for conventional crystalline polymers such as polyethylene and PP [\u003cspan additionalcitationids=\"CR49 CR50\" citationid=\"CR48\" class=\"CitationRef\"\u003e48\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR51\" class=\"CitationRef\"\u003e51\u003c/span\u003e].\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eFigure \u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003e shows the results of single-cycle deformation tests at strain rates of 0.02, 0.1, and 0.2 s\u003csup\u003e\u0026ndash;1\u003c/sup\u003e, which correspond to crosshead speeds of 0.16, 0.8, and 0.16 mm s\u003csup\u003e\u0026ndash;1\u003c/sup\u003e, respectively. The unloading process started at an engineering strain of 4. The yield stress during the first loading process decreased at a low strain rate; i.e., 0.02 s\u003csup\u003e\u0026ndash;1\u003c/sup\u003e. The reloading curves were almost independent of the applied strain rate; i.e., the stretched film showed rubber-like behavior.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eBased on the experimental results, we then prepared stretched films to examine their mechanical responses during reloading. The original blown film was stretched at 0.2 s\u003csup\u003e\u0026ndash;1\u003c/sup\u003e in the MD with an engineering strain of 4 and then removed from the tensile machine to give the stretched film. Figure\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003e shows the 2D-WAXD profiles of the original and stretched films.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eThe 2\u003cem\u003eθ\u003c/em\u003e profiles of the original and stretched films are presented in Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003e(a). Some diffraction peaks attributed to α-crystals were observed. In this figure, the Millar indices are denoted in parentheses [\u003cspan citationid=\"CR52\" class=\"CitationRef\"\u003e52\u003c/span\u003e]. The \u003cem\u003eχ\u003c/em\u003e\u003csub\u003e\u003cem\u003eXRD\u003c/em\u003e\u003c/sub\u003e values calculated using Eq.\u0026nbsp;(\u003cspan refid=\"Equ1\" class=\"InternalRef\"\u003e2\u003c/span\u003e) were 0.42 for the original film and 0.33 for the stretched film. These results indicated that the crystallinity of the film was decreased by the stretching process. The azimuthal angle distributions of the (020) plane of crystallites in the original and stretched films are shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003e(b). Intense diffraction peaks were detected on the equator, suggesting that the α-crystals were orientated along the MD. This orientation must be induced by the flow history during film processing. It should be noted that a similar distribution was detected for the stretched film, demonstrating that some oriented crystals were not destroyed by the stretching treatment.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eThe birefringence of the original and stretched films was evaluated from the optical retardation and film thickness. The birefringence at 589 nm was \u0026minus;\u0026thinsp;2.5 \u0026times; 10\u003csup\u003e\u0026ndash;4\u003c/sup\u003e for the original film and \u0026minus;\u0026thinsp;2.6 \u0026times; 10\u003csup\u003e\u0026ndash;4\u003c/sup\u003e for the stretched film. Because α-crystals of PHB show negative birefringence [\u003cspan citationid=\"CR53\" class=\"CitationRef\"\u003e53\u003c/span\u003e], this was a reasonable result and corresponded to the azimuthal angle distribution in Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003e(b).\u003c/p\u003e \u003cp\u003eThe DSC heating curves of the original and stretched films are depicted in Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003e. Although both films showed almost the same \u003cem\u003eT\u003c/em\u003e\u003csub\u003e\u003cem\u003em\u003c/em\u003e\u003c/sub\u003e of 154\u0026deg;C, their heat of fusion, i.e., crystallinity, was different. The \u003cem\u003eχ\u003c/em\u003e\u003csub\u003e\u003cem\u003eDSC\u003c/em\u003e\u003c/sub\u003e values calculated using Eq.\u0026nbsp;(\u003cspan refid=\"Equ2\" class=\"InternalRef\"\u003e3\u003c/span\u003e) were 0.19 for the original film and 0.17 for the stretched film.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eFigure \u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e8\u003c/span\u003e presents the temperature dependence of the tensile storage modulus \u003cem\u003eE\u0026rsquo;\u003c/em\u003e and loss modulus \u003cem\u003eE\u0026rdquo;\u003c/em\u003e of the original and stretched films when oscillatory strain was applied in the MD. Both films exhibited typical viscoelastic properties of crystalline polymers. Although the \u003cem\u003eE\u0026rsquo;\u003c/em\u003e values of both films were almost the same below \u003cem\u003eT\u003c/em\u003e\u003csub\u003e\u003cem\u003eg\u003c/em\u003e\u003c/sub\u003e, those of the stretched film decreased greatly beyond \u003cem\u003eT\u003c/em\u003e\u003csub\u003e\u003cem\u003eg\u003c/em\u003e\u003c/sub\u003e. The \u003cem\u003eE\u0026rsquo;\u0026rsquo;\u003c/em\u003e peak ascribed to \u003cem\u003eT\u003c/em\u003e\u003csub\u003e\u003cem\u003eg\u003c/em\u003e\u003c/sub\u003e was located at approximately \u0026minus;\u0026thinsp;10\u0026deg;C. The \u003cem\u003eE\u0026rsquo;\u0026rsquo;\u003c/em\u003e peak of the stretched film was sharper than that of the original film, which was attributed to the decrease in crystallinity after stretching. These results demonstrated that stretching decreased the crystallinity of the polymer film, as revealed by the 2D-WAXD profiles. Therefore, the stretched film exhibited a lower modulus than that of the original film, as shown in Figs.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e and \u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003e. Considering that the modulus decrease was detected only in a relatively low temperature region, i.e., from 0 to 80\u0026deg;C, only small crystals with low \u003cem\u003eT\u003c/em\u003e\u003csub\u003e\u003cem\u003em\u003c/em\u003e\u003c/sub\u003e must be destroyed by the stretching process; these small crystals do not reorganize into well-oriented crystals with high \u003cem\u003eT\u003c/em\u003e\u003csub\u003e\u003cem\u003em\u003c/em\u003e\u003c/sub\u003e after stretching. This result should be noted because conventional crystalline plastics including PP show higher modulus after stretching because of the stress/strain-induced crystallization in the solid state. The DSC results (Fig.\u0026nbsp;\u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e8\u003c/span\u003e) also indicated that well-developed crystals with relatively high \u003cem\u003eT\u003c/em\u003e\u003csub\u003e\u003cem\u003em\u003c/em\u003e\u003c/sub\u003e were not destroyed by the stretching process.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eThe optical retardation during the reloading process of the stretched film was measured. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig9\" class=\"InternalRef\"\u003e9\u003c/span\u003e, some strains were applied to the stretched films by the tensile machine, and then images under crossed polars with/without a quarter-wave plate were captured. The slow axis of the quarter-wave plate was oriented along the MD and the polarizer and analyzer were tilted at 45\u0026deg; with respect to the MD, as illustrated in Fig.\u0026nbsp;\u003cspan refid=\"Fig9\" class=\"InternalRef\"\u003e9\u003c/span\u003e. As mentioned, the stretched film showed negative birefringence at a strain (\u003cem\u003eε\u003c/em\u003e) of 0. Then, the birefringence became almost 0 at approximately \u003cem\u003eε\u003c/em\u003e\u0026thinsp;=\u0026thinsp;0.3. Finally, positive birefringence was detected under further stretching, e.g., \u003cem\u003eε\u003c/em\u003e\u0026thinsp;=\u0026thinsp;0.7, demonstrating that the sign of the birefringence changed from negative to positive during the reloading process.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eWe also evaluated the 2D-WAXD profiles at the same strain levels as those used in Fig.\u0026nbsp;\u003cspan refid=\"Fig9\" class=\"InternalRef\"\u003e9\u003c/span\u003e to reveal the deformation behavior of the stretched film. Figure\u0026nbsp;\u003cspan refid=\"Fig10\" class=\"InternalRef\"\u003e10\u003c/span\u003e shows the 2\u003cem\u003eθ\u003c/em\u003e profiles and azimuthal angle distribution of the (020) plane for the stretched film at different \u003cem\u003eε\u003c/em\u003e. The crystallinity and crystallite orientation were not affected during the reloading process, i.e., the α-crystals oriented to the MD even after stretching. Considering that a simple addition rule is applicable for birefringence [\u003cspan citationid=\"CR46\" class=\"CitationRef\"\u003e46\u003c/span\u003e, \u003cspan citationid=\"CR54\" class=\"CitationRef\"\u003e54\u003c/span\u003e, \u003cspan citationid=\"CR55\" class=\"CitationRef\"\u003e55\u003c/span\u003e], the growth of positive birefringence must be attributed to the orientation of amorphous tie chains under the applied \u003cem\u003eε\u003c/em\u003e.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eFigure \u003cspan refid=\"Fig11\" class=\"InternalRef\"\u003e11\u003c/span\u003e shows a schematic illustration of crystals and amorphous tie chains in the stretched film during tensile deformation. As revealed in the azimuthal angle distribution of the 2D-WAXD profile, the crystals oriented to the MD with no/low orientation of amorphous chains before reloading. Therefore, the birefringence of the stretched film without applied stress was negative. During deformation, i.e., the reloading process, the amorphous tie chains were orientated along MD while keeping the orientation of α-crystals. Considering that the stretched film showed rubber-like behavior with a small residual strain, amorphous tie chains between crystals must be responsible for the mechanical response during reloading.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e"},{"header":"Conclusion","content":"\u003cp\u003eThe mechanical properties of a tubular blown PHBHHx film were studied by cyclic tensile tests. During the first stretching process, the film showed a typical stress\u0026ndash;strain curve for a crystalline polymer. During stretching, the film lost small crystals with low \u003cem\u003eT\u003c/em\u003e\u003csub\u003e\u003cem\u003em\u003c/em\u003e\u003c/sub\u003e and did not show stress/strain-induced crystallization in the solid state. As a result, the crystallinity and thus the modulus of the PHBHHx film were decreased by the stretching history. Moreover, well-developed crystals oriented in the MD were not destroyed during stretching. These crystals acted as crosslink points during the reloading process. Therefore, the stress\u0026ndash;strain curve during reloading was an inverted \u0026ldquo;S\u0026rdquo; shape, and the residual strain after unloading was minimal. During the reloading process, the amorphous chains connecting crystallites oriented, which was supported by the results of birefringence measurements.\u003c/p\u003e"},{"header":"Declarations","content":"\u003cp\u003e\u003cstrong\u003eAcknowledgment\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eWe thank Natasha Lundin, PhD, from Edanz (https://jp.edanz.com/ac) for editing a draft of this manuscript.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003e\u0026nbsp;\u003c/strong\u003e\u003cstrong\u003eFunding\u0026nbsp;\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThis work was supported by JST SPRING, Japan\u0026nbsp;Grant Number JPMJSP2102.\u003c/p\u003e\n\u003cp\u003e\u0026nbsp;\u003cstrong\u003eAuthor Contributions\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eData curation and Investigation,\u0026nbsp;Y.F.; Methodology,\u0026nbsp;Y.F. and\u0026nbsp;K.J.; Writing Original Draft,\u0026nbsp;Y.F.; Samples, T.S.; Writing Review and Edit, M.Y. All authors read and approved the final manuscript.\u003c/p\u003e\n\u003cp\u003e\u0026nbsp;Conflict of interest: The authors declare no conflict of interest.\u0026nbsp;\u003c/p\u003e\n"},{"header":"References","content":"\u003col\u003e\n\u003cli\u003eDoi Y (1990) Microbial Polyesters. VCH Publishers, New York\u003c/li\u003e\n\u003cli\u003eDoi Y, Steinb\u0026uuml;chel A Editors (2001) Biopolymers: Polyesters I and II Vol 3. Wiley-VCH, Weinheim\u003c/li\u003e\n\u003cli\u003eOmura T, Tsujimoto S, Kimura S, Maehara A, Kabe K, Iwata T (2023) Marine biodegradation of poly[(R)-3-hydroxybutyrate-co-4-hydroxybutyrate] elastic fibers in seawater: dependence of decomposition rate on highly ordered structure. Front Bioeng Biotechnol 11:1303830\u003c/li\u003e\n\u003cli\u003eArakawa K, Yokohara T, Yamaguchi M (2007) Enhancement of melt elasticity for poly(3-hydroxybutyrate-co-3-hydroxyvalerate) by addition of weak gel. 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Cellulose 20:83-96\u003c/li\u003e\n\u003c/ol\u003e"}],"fulltextSource":"","fullText":"","funders":[],"hasAdminPriorityOnWorkflow":false,"hasManuscriptDocX":true,"hasOptedInToPreprint":true,"hasPassedJournalQc":"","hasAnyPriority":false,"hideJournal":false,"highlight":"","institution":"","isAcceptedByJournal":true,"isAuthorSuppliedPdf":false,"isDeskRejected":"","isHiddenFromSearch":false,"isInQc":false,"isInWorkflow":false,"isPdf":false,"isPdfUpToDate":true,"isWithdrawnOrRetracted":false,"journal":{"display":true,"email":"[email protected]","identity":"journal-of-polymers-and-the-environment","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":false,"externalIdentity":"jooe","sideBox":"Learn more about [Journal of Polymers and the Environment](https://www.springer.com/journal/10924)","snPcode":"10924","submissionUrl":"https://submission.nature.com/new-submission/10924/3","title":"Journal of Polymers and the Environment","twitterHandle":"","acdcEnabled":true,"dfaEnabled":true,"editorialSystem":"em","reportingPortfolio":"Springer Hybrid","inReviewEnabled":true,"inReviewRevisionsEnabled":false},"keywords":"Poly(3-hydroxybutyrate-co-3-hydroxyhexanoate), Mechanical property, Tensile test, Residual strain","lastPublishedDoi":"10.21203/rs.3.rs-4463452/v1","lastPublishedDoiUrl":"https://doi.org/10.21203/rs.3.rs-4463452/v1","license":{"name":"CC BY 4.0","url":"https://creativecommons.org/licenses/by/4.0/"},"manuscriptAbstract":"\u003cp\u003eThe mechanical responses during loading, unloading, and reloading cyclic tensile tests of a tubular blown film of poly(3-hydroxybutyrate-\u003cem\u003eco\u003c/em\u003e-3-hydroxyhexanoate) are studied. Although the stress\u0026ndash;strain curve recorded during the initial stretching process is typical for a crystalline polymer, the stretched film behaves like a rubber during the reloading process; that is, low modulus with a small residual strain after unloading. Furthermore, the stress\u0026ndash;strain curves during the reloading process are an inverted \u0026ldquo;S\u0026rdquo; shape. During the first stretching process of the polymer film, small crystals are destroyed without reorganization into a crystalline structure, leading to the observed decrease of crystallinity. In contrast, well-developed crystals that orient to the machine direction of the film do not disappear during the first stretching and act as crosslink points during reloading. As a result, a rubber-like response is detected. This mechanical response during reloading is considerably different from those of conventional crystalline plastics such as polyethylene and polypropylene.\u003c/p\u003e","manuscriptTitle":"Anomalous Mechanical Response of Stretched Film of Poly(3-Hydroxybutyrate-co-3-Hydroxyhexanoate)","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2024-06-05 16:39:23","doi":"10.21203/rs.3.rs-4463452/v1","editorialEvents":[{"type":"communityComments","content":0},{"type":"decision","content":"Revision requested","date":"2024-07-06T14:46:52+00:00","index":"","fulltext":""},{"type":"editorInvitedReview","content":"","date":"2024-06-20T03:04:11+00:00","index":"hide","fulltext":""},{"type":"reviewerAgreed","content":"265304245147713596218256600631755279148","date":"2024-06-14T08:54:47+00:00","index":"hide","fulltext":""},{"type":"reviewerAgreed","content":"100436976714518996421808103855330295307","date":"2024-06-13T11:37:42+00:00","index":"hide","fulltext":""},{"type":"editorInvitedReview","content":"","date":"2024-06-05T11:02:32+00:00","index":"hide","fulltext":""},{"type":"reviewerAgreed","content":"90595239447159525882870416330672256072","date":"2024-05-28T12:28:22+00:00","index":"hide","fulltext":""},{"type":"reviewerAgreed","content":"955486721069032813438512795986068341","date":"2024-05-26T11:37:10+00:00","index":"hide","fulltext":""},{"type":"reviewersInvited","content":"","date":"2024-05-26T11:25:34+00:00","index":"","fulltext":""},{"type":"checksComplete","content":"","date":"2024-05-24T07:38:45+00:00","index":"","fulltext":""},{"type":"editorAssigned","content":"","date":"2024-05-24T07:38:45+00:00","index":"","fulltext":""},{"type":"submitted","content":"Journal of Polymers and the Environment","date":"2024-05-23T01:11:58+00:00","index":"","fulltext":""}],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"journal-of-polymers-and-the-environment","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":false,"externalIdentity":"jooe","sideBox":"Learn more about [Journal of Polymers and the Environment](https://www.springer.com/journal/10924)","snPcode":"10924","submissionUrl":"https://submission.nature.com/new-submission/10924/3","title":"Journal of Polymers and the Environment","twitterHandle":"","acdcEnabled":true,"dfaEnabled":true,"editorialSystem":"em","reportingPortfolio":"Springer Hybrid","inReviewEnabled":true,"inReviewRevisionsEnabled":false}}],"origin":"","ownerIdentity":"efc2c2e6-8f0e-48d6-bb00-4e01f71fb87b","owner":[],"postedDate":"June 5th, 2024","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"published-in-journal","subjectAreas":[],"tags":[],"updatedAt":"2024-08-12T16:10:31+00:00","versionOfRecord":{"articleIdentity":"rs-4463452","link":"https://doi.org/10.1007/s10924-024-03370-x","journal":{"identity":"journal-of-polymers-and-the-environment","isVorOnly":false,"title":"Journal of Polymers and the Environment"},"publishedOn":"2024-08-05 15:58:13","publishedOnDateReadable":"August 5th, 2024"},"versionCreatedAt":"2024-06-05 16:39:23","video":"","vorDoi":"10.1007/s10924-024-03370-x","vorDoiUrl":"https://doi.org/10.1007/s10924-024-03370-x","workflowStages":[]},"version":"v1","identity":"rs-4463452","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-4463452","identity":"rs-4463452","version":["v1"]},"buildId":"8U1c8b4HqxoKbykW_rLl7","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}