Poly(l-lactide) and Poly(l-lactide-co-trimethylene carbonate) Melt-Spun

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Poly(L-lactide) and poly(L-lactide-co-trimethylene carbonate) melt-spun fibers: structure-processing-properties relationship Tiziana Fuoco, Torbjörn Mathisen, and Anna Finne-Wistrand Biomacromolecules, Just Accepted Manuscript • DOI: 10.1021/acs.biomac.8b01739 • Publication Date (Web): 21 Jan 2019 Downloaded from http://pubs.acs.org on January 22, 2019

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Poly(L-lactide) and poly(L-lactide-cotrimethylene carbonate) melt-spun fibers: structure-processing-properties relationship Tiziana Fuoco,1 Torbjörn Mathisen2 and Anna Finne-Wistrand1*

1) Department of Fibre and Polymer Technology, KTH Royal Institute of Technology, 100 44 Stockholm, Sweden 2) Novus Scientific AB, Virdings allé 2, 754 50 Uppsala, Sweden E-mail: [email protected]

ABSTRACT L-lactide/trimethylene carbonate copolymers have been produced as multifilament fibers by high-speed melt-spinning. The relationship existing between the composition, processing parameters and physical properties of the fibers has been disclosed by analyzing how the industrial process induced changes at the macromolecular level, i.e., the chain microstructure and crystallinity development. A poly(L-lactide) and three copolymers having trimethylene carbonate contents of 5, 10 and 18 mol% were synthesized with high molecular weight (Mn) up to 377 kDa and narrow dispersity. Their microstructure, crystallinity and thermal properties were dictated by the composition. The spinnability was then assessed for all the as-polymerized materials: four melt-spun multifilament fibers with increasing linear density were collected for each (co)polymer at a fixed take-up speed of 1800 m min-1 varying the mass throughput during the extrusion. A linear

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correlation resulted between the as-spun fiber properties and the linear density. The as-spun fibers could be further oriented, developing more crystallinity and improving their tensile properties by a second stage of hot-drawing. This ability was dependent on the composition and crystallinity achieved during the melt-spinning and the parameters selected for the hot-drawing, such as temperature, draw ratio and input speed. The crystalline structure evolved to a more stable form, and the degree of crystallinity increased from 0–52 % to 25–66 %. Values of tensile strength and Young’s modulus up to 0.32–0.61 GPa and 4.9–8.4 GPa were respectively achieved. Keywords Poly(L-lactide-co-trimethylene carbonate), melt-spun fibers, high-speed spinning, crystallinity development, tensile properties

INTRODUCTION Poly(lactic acid) or poly(lactide), PLA, is a bio-based, degradable and thermoplastic polymer belonging to the family of the aliphatic polyesters. PLA has a leading position because of its mechanical properties and thermoplastic processability that make it suitable for a broad range of commodity applications, in some cases representing an alternative to conventional plastic materials.1 More importantly, PLA and its copolymers have been the most-used polymers in biomedical research and applications for the past five decades since they are biocompatible and eventually degrade in vivo into non-toxic products.2 The poly(L-lactide), PLLA, usually prepared by ring-opening polymerization of the L-lactide monomer (LA), is a semicrystalline material having good mechanical properties with a melting point, Tm, of approximately 180 °C and a glass transition temperature, Tg, in the range 5560 °C.3 Its features make it an ideal candidate for load-bearing applications.4 PLA can indeed form high-strength monofilament5-7 and multifilament fibers,8-9 as well as electrospun 2 ACS Paragon Plus Environment

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fibers10-12 which also have investigated as scaffolding materials13-16 and injectable micelles carrier.17 The quality of PLA fibers makes them suitable for use in textile applications as well.18 The high crystallinity and the slow degradation time, which can be up to 3–5 years,3 could represent a limitation to the use of PLLA in biomedical applications. Copolymerization has been exploited as an effective strategy to modulate the physical and chemical properties of PLLA,19 as well as to tune the degradation profile.20 Trimethylene carbonate (TMC) represents a suitable co-monomer to modify PLLA for biomedical uses, since it can impart weaker acidity and higher flexibility to the polymeric chains. PTMC is a hydrophobic, amorphous and elastic polymer with a Tg of around -20 °C, usually used in applications for soft tissue regeneration.21 Moreover, unlike PLLA and other polyesters or their copolymers, PTMC mainly degrades by enzymatic surface erosion, and it does not produce acidic moieties upon degradation, the hydrolytic degradation being extremely slow.22 Bulk degradation has been found to be the degradation mechanism for LA/TMC copolymers.23 Several methods to synthesize copolymers consisting of LA and TMC, with various microstructures and exploiting different catalysts, have been developed for various applications and reported in patents24-25 and publications.26-30 More recently, LA/TMC copolymers have been used for the manufacturing of porous scaffolds by phase separation and a particulate leaching technique.31 In addition, cross-linked poly(LA-co-TMC) scaffolds for vascular engineering applications have been prepared.32 Electrospinning has also been used as a processing technique for the shaping of this class of materials into fibrous scaffolds.33-35 Herein, we have evaluated high-molecular-weight PLLA and related copolymers with TMC, having different compositions, to produce multifilament fibers through a high-speed meltspinning process. The production of PLA fibers by melt-spinning is indeed a well-established technology. The process consists of the melting of the polymer above its Tm and extruding it through a spinneret,



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a die with a small hole or several holes to produce a filament or a multifilament, respectively. The fiber or yarn is then cooled down, drawn and collected on a winder. After the spinning, fibers can also undergo a second stage of hot-drawing.36 It has been remarked that when compared with other methods for fibers production, melt-spinning has the advantages of higher spinning speed, solvent-free processes and the potential for mass production.37 The features of PLA melt-spun fibers have been researched by several authors.5-9,38-45 The first publications evaluated the spinning and/or drawing of monofilament at speed from a few centimeters to a few meters thus far from industrial processes,5-7 whose requirements are the production of a large quantity of material in a relatively short time. Spruiell et al. described for the first time the production of PLA multifilament fibers by a high-speed melt-spinning process. The effects of the spinning speed on the development of the fiber crystallinity, chain orientation and tensile properties were studied, concluding that the optimum take up speed of their singlestage process was between 2000 and 3000 m min-1.8 The effect of different draw ratio applied during the spinning of PLA fibers on their final properties was also lately described.38 Highspeed melt-spinning followed by hot-drawing was then reported by Schmack et al. for the production of multifilament yarn consisting of PLA containing a low percentage of mesolactide.9 It emerges that the end properties of PLA fibers can be modulated by selecting the molecular weight and composition of the polymer as well as the processing conditions. It has been disclosed how the crystallinity and mechanical performance develop after extrusion as the fibers are cooled down and drawn. If a second stage of further hot-drawing is performed, the chain orientation can be improved, thus, enhancing the properties of the fibers.9,45 The effect of fillers46 and nucleating agents47 on the final fiber properties has been also reported. However, an exhaustive correlation between all the above mentioned factors is prevented by the fact that the authors often analyzed a few parameters at time and did not give a complete overview of


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the process investigated. Comparison between the properties of the PLA fibers reported in different publications is also not possible because of the different material features, spinning devices and parameters, the diameter and number of filaments, and the characterization methods utilized. Furthermore, research dealing with fiber formation for PLLA and related copolymers through a two-stage industrial process, i.e., high speed melt-spinning followed by hot-drawing, is lacking in the literature. We reasoned that the comprehension of an industrial process from the macromolecular point of view, i.e., chain composition, microstructure and crystalline structure, could allow us to properly design the polymer features and select the processing conditions to match the mechanical properties of the demanded application. Indeed, all the practical applications of a material are dictated by its properties, and this is of key importance for the development of medical devices. Therefore, we aimed at understanding the relationship existing between the macromolecular structure, processing parameters and final properties of PLLA and poly(LA-co-TMC) multifilament fibers produced by a high-speed melt-spinning process followed by hot-drawing on an industrial scale. For this purpose, we have synthesized four (co)polymers, having different compositions. We have evaluated how the melt-extrusion and processing parameters, and the draw ratio, drawing temperature and speed modify the molecular weight and microstructure and induce crystallinity. These in turn determine the thermal properties and tensile behavior of the end fibers.

EXPERIMENTAL SECTION Materials Moisture and air-sensitive materials were manipulated under nitrogen or by using an MBraun glovebox. L-Lactide (LA) was purchased from Purac, trimethylene carbonate (TMC) was purchased from Richman chemical, and they were used as received. Stannous octanoate 5 ACS Paragon Plus Environment


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(SnOct2), propanediol and dry toluene were purchased from Sigma Aldrich and used as received. Polymers Synthesis Poly(L-lactide) 100LA Before the polymerization, the polymerization apparatus, consisting of a premixing chamber and a reactor, was heated at 130 °C under nitrogen flow for 24 hours. The L-lactide monomer (LA) (3002.06 g; 20.83 mol) was charged into the premixing chamber and was allowed to melt with stirring at 130 °C. 1,3-Propane diol (0.80 mL; 11.1 mmol) and stannous octanoate (1.60 mL of a stock solution in toluene C = 1.09 M; 1.74 mmol) were used as the initiator and catalyst, respectively, and were sequentially added to the premixing chamber. The polymerization mixture was stirred for 1 minute and then transferred to the reactor under nitrogen flow where it was allowed to polymerize at 130 °C for 48 hours. Poly(LA) = 1H NMR (400 MHz; CDCl3) δ 5.16 (q, J = 7.1 Hz, 1H, CHCH3), 1.58 (d, J = 7.1 Hz, 3H, CHCH3). 13C

NMR (100 MHz; CDCl3) δ 169.7 (OCH(CH3)C(O)), 69.2 (OCH(CH3)C(O)), 16.8

(OCH(CH3)C(O)). Poly(L-lactide-co-trimethylene carbonate) 95LA The polymerization was performed as above, but LA (3855.00 g; 26.75 mol) and TMC (142.88 g; 1.20 mol) were used as the monomers. 1,3-Propane diol (1.09 mL; 15.2 mmol) and stannous octanoate (1.60 mL of a stock solution in toluene C = 1.09 M; 1.74 mmol) were used as the initiator and catalyst, respectively.


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Poly(L-lactide-co-trimethylene carbonate) 90LA The polymerization was performed as above, but LA (3702.04 g; 25.71 mol) and TMC (298.65 g; 2.92 mol) were used as monomers. 1,3-Propane diol (1.09 mL; 15.2 mmol) and stannous octanoate (2.18 mL of a stock solution in toluene C = 1.09 M; 2.37 mmol) were used as the initiator and catalyst, respectively. Poly(L-lactide-co-trimethylene carbonate) 80LA The polymerization was performed as above, but LA (2774.49 g; 19.24 mol) and TMC (491.05 g; 4.81 mol) were used as monomers. 1,3-Propane diol (0.80 mL; 11.1 mmol) and stannous octanoate (1.60 mL of a stock solution in toluene C = 1.09 M; 1.74 mmol) were used as the initiator and catalyst, respectively. Poly(LA-co-TMC) = 1H NMR (400 MHz; CDCl3) δ 5.16 (q, J = 7.1 Hz, 1H, CHCH3), 4.304.16 (m, 4H, OCH2CH2CH2OC(O), OCH2CH2CH2OC(O)), 2.101.98 (m, 2H, OCH2CH2 CH2OC(O)), 1.58 (d, J = 7.1 Hz, 3H, CHCH3). 13C

NMR (100 MHz; CDCl3) δ 170.0 and 169.7 (OCH(CH3)C(O)), 155.0 and 154.4

(OCH2CH2CH2OC(O)), 71.5, 69.3 and 69.1 (OCH(CH3)C(O)), 64.9, 64.4 and 61.9 (OCH2CH2CH2OC(O), OCH2CH2CH2OC(O)), 28.0 (OCH2CH2CH2OC(O)) 16.79 (OCH(CH3)C(O)). Processing Poly(LA) and poly(LA-co-TMC) fibers were produced using a two-stage process, meltspinning followed by hot-drawing. Melt-spinning Melt-spun multifilament fibers, for each polymer, were prepared by using a spinline consisting of a single-screw extruder, a melt pump to control the mass throughput (Q), a spin-pack 7 ACS Paragon Plus Environment


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consisting of a filter, a breaker-plate and a multifilament spinneret. The multifilament is transported to the winder by means of several godets (Figure 1a). Specifically, the polymers were melted and homogenized in a single-screw extruder with three heating zones, set at specific temperatures for the different materials (Table 1), having a pressure-controlled rotational speed on the extruder screw to create an output pressure of 60 bar. The molten polymer was fed into a spin pump operated at constant speed to generate a constant flow of polymer to the spinneret, which contained several openings, each having an exit diameter of 0.45 mm. Table 1. Extruder zone temperatures (°C) for melt-spinning. Polymer

Zone 1

Zone 2

Zone 3

Spinneret

100LA

180

235

245/

250

95LA

165

235

240

240

90LA

165

235

240

250

80LA

160

220

230

250

The multifilament yarns were then coated with spin finishing oil and drawn through the draw line consisting of four godets (G) heated at 80 °C and a winder (W) (Figure 1a). The fibers were then collected at a final speed of W = 1800 m min-1. The speeds of the godets (G) were slightly slower than the winder (W) to keep tension on the yarn: G1 = 1760 m min-1 and G2, G3 and G4 = 1770 m min-1.


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Figure 1. Schematic representation (not to scale) of the fiber spinning line (a) and hotdrawing line (b). The take-up speed of the winder (W) was maintained constant at 1800 m min-1, while the mass throughput (Q) was increased in order to have, for each polymer batch, four fiber bobbins, named T1, T2, T3, T4, having linear density (LD1) of approximately 125, 135, 145 and 155 deniers, respectively (Table 2). After each change in the spin pump speed, the system was equilibrated until a stable pressure over the spin-pack was achieved. The fiber was then collected for 5 minutes for each bobbin T1 through T4. Table 2. Parameters used for melt-spinning. Bobbin

Q (g min-1)

G1 (m min-1)

G2, G3, G4 (m

W (m min-1)

min-1)

Draw

Theoretical

down

linear density

ratioa

(denier)

T1

25.06

1760

1770

1800

569.62

125.3

T2

27.00

1760

1770

1800

528.51

135.0

T3

28.95

1760

1770

1800

492.94

144.8

T4

31.04

1760

1770

1800

459.78

155.2

aCalculated

as previously reported.48



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Hot-drawing Before hot-drawing of the as-spun fibers, bobbins were stored under a vacuum for at least 48 hours. A schematic representation of the post-draw line is shown in Figure 1b. It consists of three godets (G1’, G2’, G3’), two ovens placed between the godets, and a winder (W’). For each of the four bobbins obtained from batches 100LA, 95LA and 90LA, hot-drawing of the fibers was carried out at two different temperatures, 80 and 110 °C, and for each temperature, two input speeds were evaluated, 50 and 300 m min-1, which correspond to the speed of G1’. Therefore, for each as-spun fiber sample, four bobbins were collected after hot-drawing, TXa (T = 80 °C, G1’ = 50 m min-1), TXb (T = 80 °C, G1’ = 300 m min-1), TXc (T = 110 °C, G1’ = 50 m min-1), and TXd (T = 110 °C, G1’ = 300 m min-1), where X is intended to be 1, 2, 3 or 4. Each of the four fiber bobbins melt-spun from batch 80LA was hot-drawn at 80 °C at both input speeds of 50 and 300 m min-1. Moreover, T1 (LD1 = 125 denier) was hot-drawn at different temperatures of 70, 80, 90, 100 and 110 °C, while the input speed was maintained constant at 300 m min-1. To hot-draw the fibers at a final linear density of 85 denier (LD2), the speed of the winder (W’) and of G2’ and G3’ was adjusted according to the following equations: 𝑊′ = 0.98

𝐺1′ 𝐿𝐷1 𝐿𝐷2

𝐺2′ = 𝐺3′ = 1.1 𝑊′ Characterization methods NMR spectra of the polymer and as-spun fiber samples were obtained in CDCl3 at room temperature on a Bruker Avance 400 spectrometer (1H: 400.13; 13C: 100.62 MHz; respectively). The resonances and coupling constants are reported in ppm (δ) and Hz (J), respectively. 1H NMR spectra were referenced to the residual solvent proton at δ 7.26 ppm; 13C NMR spectra were referenced to the 13C signal of CDCl3 at δ 77.16 ppm. Spectra were recorded on Bruker TopSpin v2.1 software. Data processing was performed using MestReNova v9.0.0 software. 10 ACS Paragon Plus Environment

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The peaks in the 1H and 13C NMR spectra were assigned according to the literature.29-30, 49 The average block lengths of lactidyl (LLL) and trimethylene carbonate blocks (LT) were calculated from the methine and methylene region of the

13C

NMR spectra according to the following

equations:

Molecular weights (Mn and Mw) and dispersities (Ð) were measured by size exclusion chromatography (SEC). The measurements were performed at 30 ºC on a Verotech PL-GPC 50 Plus system equipped with two PLgel 5 µm MIXED-D (300 × 7.5 mm) columns, a PL-RI detector (Varian, Germany) and a PL-GPC 50 Plus autosampler using CHCl3 as the eluent (1.0 mL min-1). Narrow polystyrene standards were used as references, and the flow rate fluctuations were corrected using toluene as an internal standard. Fiber linear density was determined by measuring the mass in grams, m (g), of 60 m of yarn. Three measurements were performed for each sample, and the mean value was converted to denier according to the following equation:

Glass transition temperatures (Tg), melting points (Tm) and enthalpies of fusion (ΔHm) of the polymeric samples and fibers were measured by differential scanning calorimetry (DSC) using aluminum pans and a Mettler Toledo DSC 1 calibrated with indium. Measurements were performed under nitrogen flow with a heating rate of 10 °C min-1 from  to  220 °C for samples from batches 100LA and 95LA and from 25 to  190 °C for samples from batches 11 ACS Paragon Plus Environment


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90LA and 80LA. The DSC data are reported for the first heating cycle, the glass transition temperature is taken as the midpoint ISO and the melting temperatures are taken as the maximum value of the endothermic peaks. The degree of crystallinity, Xc,DSC, was calculated from DSC considering an enthalpy of fusion for an infinitely large PLA crystal of 93.0 J g-1.50 Wide-angle X-ray diffraction (WAXD) scans of the raw materials (granulate) and as-spun and hot-drawn fibers were obtained using a Thermo Scientific ARL X'TRA X-ray diffractometer. A nickel-filtered CuK radiation source was used in the reflection mode and the goniometer was scanned in the 2θ range from 5 to 35°. Before WAXD, the fiber samples were cut into pieces 23 mm long and set on the sample holder. To calculate the degree of crystallinity, Xc,WAXD, the WAXD profile of the as-polymerized materials was fitted around the peaks at 2θ 14.9, 16.7, 19.1, and 22.3° as the sum of the amorphous component and the crystalline components using a Gaussian function. In contrast, the WAXD profile of the fiber samples was fitted around the main peak, 2θ = 16.2–16.8°, as the sum of the amorphous component and the crystalline components using a Gaussian function. The crystallite size was estimated by the Scherrer equation using the width of the main crystalline peak, 2θ = 16.2–16.8°, as previous reported for PLA fibers;38, 45 the lattice spacing d and the unit cell parameter a were estimated from the diffraction angle of the main crystalline peak, 2θ = 16.2–16.8°, relative to the (200)/(110) planes, assuming an orthorhombic crystal system. The tensile mechanical properties of fibers were measured at room temperature using a Zwick Z2.5 tensile testing machine equipped with pneumatic fiber grips. The free clamped length was 250 mm. A 100 N load cell was used, and the cross-head speed was 250 mm min-1. All the reported tensile properties reported are the average values of seven tests. The fiber morphologies and filament diameters were analyzed in a Hitachi S-4800 field emission scanning electron microscope (FE-SEM) at an acceleration voltage of 1 kV. Samples 12 ACS Paragon Plus Environment

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were sputter-coated with a layer of platinum. The filament diameters reported are the mean values of at least six measurements.

RESULTS AND DISCUSSION Synthesis and characterization of poly(L-lactide) and poly(L-lactide-co-trimethylene carbonate) copolymers To evaluate the effect of copolymerization and processing parameters on the fiber formation and their final properties, four different polymers were synthetized: a poly(L-lactide) (100LA) and three poly(L-lactide-co-trimethylene carbonate) copolymers having different content of TMC, 5 mol% (95LA), 10 mol% (90LA) and 18 mol% (80LA). The polymers were prepared by ring-opening (co)polymerization of L-lactide (LA) and trimethylene carbonate (TMC) using SnOct2 as the catalyst and propanediol as the initiator (Scheme 1). The polymerizations were performed at 130 °C for 48 hours. The as-prepared materials were fully characterized by NMR, SEC, DSC and WAXD, and the results are summarized in Table 3. O O

O O O+

SnOct2 OH O HO

O

bulk, 130 °C, 48 h

O

O O

O

O

0.82-1

O

O 0-0.18

n

Scheme 1. Ring-opening copolymerization of L-lactide and trimethylene carbonate. Table 3. Characterization of the as-prepared PLLA and P(TMC-co-LA) copolymers. Polymer

LAa

TMCa

LLLb

LTb

Mnc

(mol%)

(mol%)

100LA

100

-

-

-

361.8

95LA

95

5

30.7

1.0

90LA

90

10

13.8

80LA

82

18

8.1

Ðc

Tgd

Tmd

Xc,DSCd

Xc,WAXDe

(°C)

(°C)

(%)

(%)

1.2

68.6

196.2

82.5

76.8

376.8

1.2

51.9

185.75

69.3

74.5

1.3

316.6

1.3

46.0

179.3

54.3

42.1

1.3

320.2

1.3

46.9

167.9

22.0

18.8

(kDa)



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aCopolymer

composition calculated from 1H NMR. bCalculated from the methine and methylene region of the 13C NMR spectra. cDetermined by SEC in CHCl3. dDetermined by DSC, date reported for the first heating run. eDetermined from WAXD.

The composition, determined by 1H NMR, reflected the feed for 95LA and 90LA copolymers; however, for batch 80LA, it was slightly richer in LA than the feed, which contained a LA/TMC ratio of 80 to 20 mol%. The TMC monomer conversion after 48 hours was lower than the LA conversion, and the raw materials contained 4 mol% of unreacted TMC. This could be due to a slightly higher reactivity ratio of LA than TMC.32 Increasing the TMC content in the feed probably requires a longer reaction time. The copolymer microstructure was analyzed by inspection of the methine and methylene regions of the 13C NMR spectra. It has been reported that for LA/TMC copolymers the carbonyl signals are less sensitive to the chemical environment than the methine and methylene resonances, δ 60-75 ppm.30, 49 In Figure 2, the methine and methylene region of the 13C NMR spectra for all the copolymer samples is shown, and for comparison, the 13C NMR spectrum of PLLA (100LA) is also reported. A complete assignment of the peaks was carried out according to the literature (Table S2).29-30, 49


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Figure 2. Details of the methine and methylene region of the 13C NMR spectra (100 MHz, CDCl3) of the as-prepared PLLA and P(TMC-co-LA) copolymers. (*) LA and (#) TMC unreacted monomers. In the methine region of the spectra of the copolymers, δ 72–68 ppm, in addition to the signal of the long lactyl homosequence (LLLL) at 69.1 ppm, two more signals were observed at δ 71.5 and 69.3 ppm. These are due to block(s) junction(s), such as TLL and LLT sequences. An increment of their intensity was observed by increasing the amount of TMC. The methylene region of the 13C NMR spectra of the copolymers 80LA and 90LA displayed three resonances centered on the methylene carbons (-OCH2- and -CH2O-) of the TMC units, two of them attributable to the heterosequences (LT’ + T’’L at δ 64.9 ppm and T’L + LT’’ at 61.9) and one attributable to the homosequence (TT at δ 64.4 ppm). The resonance due to the TT homosequence was not observed for the copolymer 95LA. A quantitative analysis of the microstructure was performed by calculating the average block length of lactidyl (LLL) and trimethylene carbonate (LT) blocks (Table 3) from the methine and methylene regions for all the copolymers. The LLL decreased from 30.7 to 8.1 by increasing the amount of TMC in the copolymer, the LT remained nearly constant, and they are equal to 1.0 for 95LA and 1.3 for the 90LA and 80LA batches. Such results display a copolymer microstructure consisting of long lactidyl blocks split by isolated trimethylene carbonate units. Unimodal distribution was observed for all the samples and high molecular weights, Mn, up to 376.8 kDa were achieved. The Mn decreased slightly by increasing the amount of TMC. Ð was narrow, with values of 1.21.3 in all the cases, indicating a good control on the polymerization (Table 3). High molecular weight and narrower Ð are important for the spinnability of the material to obtain fibers with good physical properties.43 The composition and molecular weight have explicit effects on thermal and structural features, such as the glass transition temperature, Tg, melting temperature, Tm, and degree of crystallinity, 15 ACS Paragon Plus Environment


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Xc. Figure 3 shows the first DSC heating run and the WAXD scan of the as-polymerized materials. The experimental data are also summarized in Table 3.

Figure 3. DSC thermograms reported for the first heating run (a) and WAXD scans (b) of the as-prepared PLLA and P(TMC-co-LA) copolymers. The first heating runs for all the as-prepared (co)polymers are reported in Figure 3a. The PLLA sample (100LA) showed a glass transition with Tg of 68.6 °C and a melting endotherm peak with Tm of 196.2 °C. The relative degree of crystallinity was calculated to be 82.5 %. The high melting peak value and crystallinity are probably due to the high molecular weight of the polymer and to the favorable conditions for crystallization achieved during the polymerization process. Indeed, as-polymerized PLLA having viscosimetric molecular weight up to 106 kDa has been shown to be highly crystalline with melting points up to 207 °C.51 The higher-thanexpected Tg value observed in the first heating run is also a consequence of the degree of crystallinity.52 A decreasing Tg value with increasing incorporation of TMC was observed, with copolymers 95LA and 90 LA displaying the transition at 51.9 and 46.0 °C, respectively. However, for the 80LA sample, the glass transition overlapped with a small endothermic peak that caused inaccuracy in the Tg value determination at 46.9 °C (Figure 3a). Such a peak may be due either to the melting of the unreacted TMC,53 or related to a relaxation phenomenon.50 16 ACS Paragon Plus Environment

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A melting peak was also observed for all the copolymers in the first heating run, indicating that semicrystalline materials were obtained after the copolymerization (Table 3, Figure 3a). By plotting the Tm values versus the TMC content, a perfect linear correlation is obtained (Figure 4a). The calculated degree of crystallinity (%), determined by DSC on the basis of the melting heat, follows the same trend: it linearly decreases with the amount of TMC incorporated in the copolymer (Figure 4b). A reduction of the melting point and heat of fusion was already observed for LA/TMC copolymers by increasing the TMC content.26 It is interesting that the 80LA sample, containing 18 mol% of TMC, also displayed a certain degree of crystallinity (Table 3, Figure 4).

Figure 4. Plots of the melting temperature, Tm, as determined by the DSC first heating run (a) and of the degree crystallinity (%) as determined by WAXD and DSC versus the content of TMC (mol%) in the copolymers (b). In the cooling run no crystallization peaks were detected for any of the as-prepared (co)polymers. It is known from the literature that crystallization of PLLA does not occur on cooling the melt at a high rate54 because the low mobility of the high-molecular-weight PLLA 17 ACS Paragon Plus Environment


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hampered the process. The glass transition for all the (co)polymers was detected in the cooling run and second heating runs. DSC data obtained from the second heating run are summarized in Table S1. The glass transition for pure PLLA (100LA) in the second heating run occurred at a lower temperature, 57.4 °C, than during the first heating run. Moreover, the thermograms of the second heating run of the 100LA and 95LA samples showed exothermic transitions due to the cold crystallization at 137.1 and 133.2 °C, followed by the relative melting endotherm peaks, which revert to lower temperatures than in the first run, 175.5 and 166.6 °C, respectively. Negligible melting peaks were instead observed for 90LA and 80LA copolymers during the second heating run, indicating that by increasing the TMC content, the crystallization kinetics became slower. PLLA crystallizes preferentially in the most stable -form at a temperature above 120 °C.53 The single melting peak observed for the pure PLLA, 100LA, and for all the as-copolymerized materials could indicate the presence of only -crystals. The formation of the -form due to the favorable crystallization conditions achieved during the polymerization process was confirmed by wide-angle X-ray diffraction (WAXD). The WAXD scans of all the as-prepared (co)polymers are shown in Figure 3b. The two strong reflection peaks observed for 100LA at 2θ of 16.8 and 19.2°, corresponding to the (200)/(110) and (203) planes, respectively, are very close to the main peaks of the -form of PLLA reported by Zhang et al.55 at 2θ of 16.7 and 19.1°. Moreover, the peak at 2θ 22.4°, attributable to the (015) plane, is close to the typical peak at 2θ of 22.3° of the -form.55 All the reflection peaks observed for 100LA were also detected for the copolymers 95LA and 90LA. However, they were slightly shifted to lower 2θ values, which is an indication of greater lattice spacing, d, thus suggesting that the presence of the TMC unit along the polymeric chains hampers their packing in the crystal lattice. For the 80LA sample, the WAXD scan showed only low reflection peaks corresponding to the (200)/(110) and (203) planes. 18 ACS Paragon Plus Environment

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Therefore, the WAXD results proved that all the as-prepared materials were semicrystalline, and the degree of crystallinity determined by WAXD (Xc,WAXD, Table 3) shows the same dependence on the polymer composition (Figure 4b), although it differs slightly from the values determined by DSC. Moreover, on the basis of the plot in Figure 4b, it may be speculated that copolymers prepared under the same conditions having a content of TMC higher than 23-25 mol% would be amorphous. In other words, considering that the unit cell of the -form of PLLA contains two antiparallel chains and each chain has 5 monomeric LA residues in the cell,56 an increase in the amount of TMC will lead to a decrease in the LLL below 5, so that the copolymer could not crystallize. Melt-spinning The spinnability of the as-prepared (co)polymers, which had different compositions in LA and TMC, and how the processing parameters affect the macromolecular features and determine the end properties of the as-spun fibers were evaluated by a high-speed melt-spinning process. The granulate polymers were fed into the extruder, where their melting temperature was exceeded by approximately 50-60 °C; then, within the spin-pack, the molten polymers were channeled through a multi-capillary spinneret held at 240-250 °C (Table 1). Each hole had a diameter of 0.45 mm and was responsible for the formation of a single filament that was then combined in a multifilament yarn and transported to the winder via the godets (Figure 1a). All the synthesized (co)polymers showed a good spinnability, and for each of them, at least four fiber bobbins having different linear densities, named T1-4, were collected (Table 2). Interestingly, it was also possible to spin fibers from the copolymers 90LA and 80LA that had a high content of TMC. It is well known that the thermal processing of PLA causes an undesired reduction in the molecular weight due to its limited stability at high temperature.57 The rigorous conditions the material is subjected to could trigger chemical reactions such as hydrolysis, depolymerization, 19 ACS Paragon Plus Environment


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and oxidative degradation, as well as intramolecular and intermolecular trans-esterification reactions, which are responsible for the degradation of the polymeric chains. Because of the conditions used for the melt-spinning of the (co)polymers, especially the high temperature reached during the extrusion, we expected undesired reactions that would have affected the macromolecular structure of the polymer chains.58 Indeed, the reduction in molecular weight as well as microstructural changes have direct consequences on the final properties of the fibers. Thus, the molecular weight and the microstructure of the as-spun fiber bobbins (T1-4) collected for each (co)polymer were respectively analyzed by SEC and NMR and compared with those of the raw materials (Table 4). Table 4. Molecular weight, molecular weight loss and microstructural analysis of as-spun fibers. Polymer

Mna (kDa)

Ða

Mn lossb

LLLc

LTMCc

(%) 100LA

95LA

90LA

Raw material

361.8

1.2

-

-

-

T1

263.0

1.3

27.3

-

-

T2

250.0

1.3

30.9

-

-

T3

257.3

1.3

28.9

-

-

T4

241.1

1.3

33.3

-

-

Raw material

376.8

1.3

-

30.7

1.0

T1

221.3

1.4

41.3

23.8

1.0

T2

200.0

1.5

46.9

29.0

1.0

T3

219.8

1.4

41.7

29.4

1.1

T4

229.0

1.3

39.2

27.5

1.2

Raw material

316.6

1.3

-

13.8

1.3

T1

205.2

1.4

35.2

13.0

1.3

T2

222.7

1.3

29.6

11.5

1.3

T3

234.9

1.3

25.8

12.5

1.3


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Biomacromolecules

80LA

T4

212.6

1.4

32.9

11.6

1.3

Raw material

320.3

1.3

-

8.1

1.3

T1

177.8

1.4

44.5

7.6

1.3

T2

157.6

1.5

50.8

7.3

1.3

T3

167.2

1.5

47.8

6.9

1.3

T4

188.3

1.3

41.2

7.1

1.3

aDetermined cCalculated

by SEC in CHCl3. bMn loss (%) = [(Mn(raw material)-Mn(as-spun fiber))/Mn(raw material)]×100. from the methine and methylene region of the 13C NMR spectra.

For all the polymer batches, a drop in the molecular weight after extrusion was observed: Mn of the as-spun fibers was 2651 % lower than the Mn of the raw materials as shown in Figure 5.

Figure 5. Plot of the molecular weight, Mn, of raw materials and as-spun fibers prepared from (co)polymers 100LA, 95LA, 90LA and 80LA. As-spun fibers obtained from the polymer batch 100LA had on average the highest molecular weight. Mn was in the range 241–263 kDa, and the Mn loss with respect to the raw material was approximately 30 %. As-spun fibers collected from batches 95LA and 90LA have similar molecular weight but lower than that of the 100LA fibers, with Mn in the range of 200–235 kDa. However, for copolymer 95LA, which was the raw material with the highest molecular weight, the Mn loss after extrusion was on average 42 %, greater than the Mn loss of 31 % found for batch 90LA (Table 4, Figure 5). The melt-spinning of copolymer 80LA produced the asspun fibers with the lowest molecular weight. Mn was below 200 kDa and varied in the range 21 ACS Paragon Plus Environment


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of 157–188 kDa. Indeed, the major extent of molecular weight decrease was observed for this batch, with the Mn loss of 46 % on average (Table 4, Figure 5). As a consequence of the undesired reactions occurring during the extrusion, the molecular weight distribution of the asspun fibers displayed a slightly broader dispersity, Ð, than did the native (co)polymers (Table 4). The LLL of the as-spun fiber samples calculated from 13C NMR spectra, within the accuracy of the NMR measurements, was slightly lower than the LLL of the raw material (Table 4), indicating that the occurrence of side reactions is also responsible for small microstructural changes. Moreover, in the NMR spectra of the as-spun fibers, an increase in the peaks relative to the LA monomer of 24 % was observed with respect to the raw material. This suggests that depolymerization can be one of the side reactions responsible for molecular weight reduction. While degradation of the polymeric chain occurs in the extruder, the drawing of fibers at highspeed through the spinning line outlined in Figure 1a is responsible for their morphology and structure development. During the melt-spinning, the take-up speed was maintained constant at 1800 m min-1, while the mass throughput, Q, was changed in order to collect fibers with different and predictable linear density. The four fiber bobbins named T1, T2, T3 and T4, having linear densities of approximately 125, 135, 145 and 155 deniers, respectively, could be spun for each polymer batch using the spinning parameters reported in Table 2. The take-up speed was much higher than the extrusion velocity. The ratio between them is expressed in Table 2 as the draw down ratio. A lower linear density of the yarns, i.e., less mass throughput from the extruder, corresponds to a higher draw down ratio. The experimental linear density was determined for all the fiber bobbins (Table S3), and they were very close to the theoretical values, thus confirming that the spinning conditions were maintained during the process for all the polymer batches. The diameter of each as-spun filament should also increase according to the linear density of the yarn. The SEM images obtained for batch 95LA (Figure 22 ACS Paragon Plus Environment

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6) clearly show an increase in the diameter within the linear density, and it was estimated to vary in the range of 17-21 m.

Figure 6. SEM images of as-spun fibers prepared from copolymer 95LA having different linear densities: 125 (T1), 135 (T2), 145 (T3) and 155 denier (T4). Crystallinity, thermal and mechanical behavior of as-spun fibers The spinning line comprised four godets (G1-G4) heated at 80 °C, whose speed was maintained slightly lower than the winder (W), Figure 1a and Table 2. The drawing of the yarn through this system allowed the structure evolution of the fibers: the development of stress due to the high speed and the temperature above the glass transition promote structural changes, such as chain orientation and crystallization, that in turn affect the physical and mechanical properties of the as-spun fibers. The crystallinity and crystal structure of the as-spun fibers were analyzed by WAXD. In Figure 7a, the diffraction scans of as-spun fibers T1 (LD = 125 denier) obtained from the melt-spinning of the four different polymer batches are compared. The diffraction scans for all the as-spun fibers (T1-4) obtained for each polymer batch are reported in the Supporting Information (Figure S1).



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Figure 7. WAXD scans (a), DSC thermograms reported for the first heating run (b) and specific stress-strain curves with inset of the initial portion of the curves (up to 5 % strain) (c) of the as-spun rolls T1 for all the polymer batches. The obtained fiber diffraction scans are similar to those reported in the literature for PLLA fibers.38 A strong peak corresponding to the (200)/(110) plane was observed at 2θ approximately 16.3° (Figure 7a). Its position and width gave us information about the crystalline structure and the crystallite size. Data for all the as-spun fibers are summarized in Tables S4-S7. From the main peak position, the lattice spacing d was calculated to be in the 24 ACS Paragon Plus Environment

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Biomacromolecules

range of 0.547–0.541 nm, and the cell parameter a was determined to be 1.08–1.09 nm. It seems evident that the crystalline structure of the as-spun fibers is similar to the ’-form of PLLA, for which the strongest diffraction peak is reported at 2θ 16.4°, d is 0.541 nm and a is 1.08 nm.59 The formation of the ’-form is reasonable since the as-spun fibers could crystallize at a temperature of 80 °C, which is in the temperature range where the ’-form is obtained.60 From the width of the peak, the crystallite size of the (200)/(110) planes was calculated by the Scherrer equation. It decreased in the range of 13–4 nm from batch 100LA to batch 80LA, indicating that the presence of TMC units along the polymer chains prevented the long-range order of the lactidyl blocks in the crystalline lattice. The degree of crystallinity of the as-spun fibers also decreased with the amount TMC in the copolymers. For each polymer batch, a linear reduction of the degree of crystallinity with the linear density, i.e., by decreasing the draw down ratio or by increasing the mass throughput Q (Table 2), was found (Figure 8).

Figure 8. Plot of the degree of crystallinity (%) as determined by WAXD versus the linear density (denier) of the as-spun fibers prepared from (co)polymers 100LA, 95LA, 90LA and 80LA. The as-spun 100LA fibers have the highest degree of crystallinity, 5239 %, while the degree of crystallinity decreased in the range of 4025 % and 2515 % for the as-spun fibers collected



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from polymer batches 95LA and 90LA, respectively (Tables S4-S6). Interestingly, the as-spun fibers collected by extruding the 80LA batch also developed a certain degree of crystallinity, 13-7 %, apart from the yarn having the highest linear density, which resulted to be amorphous (Figure 8, Table S7). The linear dependence of the crystallinity with the draw down ratio is a clear indication that during the spinning, stress-induced crystallization was promoted, and the degree of crystallinity could be modulated by changing the spinning parameters. However, the temperature at which the godets are heated is also an important factor. It should be above the Tg in order to facilitate the macromolecular mobility. This was proved by spinning a yarn roll for batch 90LA (T5, Table S5) at a godet temperature of 50 °C, which is close to the Tg of the raw material. The resulting degree of crystallinity was 4 %, a value much lower than that determined for roll T3 (degree of crystallinity of 19 %) obtained under the same spinning conditions but at a godet temperature of 80 °C (Table S6, Figure S2). Representative DSC thermograms for the first heating run of the T1 fiber rolls spun from all the (co)polymers are shown in Figure 7b, while all the thermal analysis data are reported in the Supporting Information (Tables S8-S11). The degree of crystallinity and the crystal structure influence the thermal properties. By increasing the amount of TMC, the melting peak of the asspun fibers shifted to a lower temperature. Tm was observed in the range of 167–177 °C for the 100LA as-spun fibers, 164–167 °C for 95LA, 156–159 °C for 90LA and 145–146 °C for 80LA. The ΔHm decreased accordingly, confirming that the degree of crystallinity was reduced. However, a double-melting behavior was generally observed for the as-spun fibers: a shoulder at the low-temperature side was clearly visible in the melting region of the thermograms for rolls T1 in Figure 7b. The double-melting behavior has been generally observed for PLLA asspun fibers.8 It could be attributed to the melting of different crystalline forms that could coexist


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Biomacromolecules

in the sample or to the melting of the original crystals occurring at a lower temperature, followed by the melting of the crystals formed or perfected during the DSC scan at a higher temperature.61 From the thermograms reported in Figure 7b, it can be observed that the Tg, occurring in the range 55–51 °C for the 100LA, 95LA and 90LA as-spun fibers and 46–41 °C for the 80LA asspun fibers, overlaps with an endothermic peak, which is in turn followed by an exothermic peak. The endothermic peak can be related to the relaxation of the tie-macromolecules in the amorphous phase of the fibers.38 These are blocked in an unstable conformation between the crystallites, and just after the Tg, they have enough mobility to move and settle back to a more stable conformation. The related heat of relaxation, ΔHrel, is slightly higher for the as-spun fibers having more TMC, i.e., for batches 90LA and 80LA (Tables S8-S11). While the exothermic peak observed immediately above the glass transition, approximately 69 °C for the 100LA, 95LA and 90LA batches and 77–80 °C for the 80LA batch (Tables S8-S11), is presumably due to the crystallization of strained and/or oriented chain segments developed during the drawing that had not crystallized during the processing. Because of the high degree of chain orientation, the crystallization rate could be increased, thus explaining the low crystallization temperature in the DSC run.8, 43, 62 The related peak areas (ΔHc) decreased as the TMC content decreased and slightly diminished with the draw down ratio, indicating that the observed cold crystallization during DSC is less important for fiber samples that had achieved a higher crystallinity during the spinning. Because of the transitions overlapping, the Tg values and the degree of crystallinity calculated from the DSC for the as-spun fiber samples (Table S8-S11) should be taken with special care. Indeed, only modulated DSC could allow a more accurate separation of the glass transition from the relaxation and to distinguish between recrystallization and melting phenomenon.61 Figure 7c shows representative stress-strain curves for T1 as-spun fibers as a function of the TMC content, and an enlargement of the first portion of the curves is also shown. The



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mechanical properties of the as-spun fibers are reported in the Supporting Information (Tables S12-S15). Young’s modulus, E, decreases with the TMC content and therefore with the crystallinity trend as can be observed from the initial region of the stress-strain curves in the inset of Figure 7c. For the 100LA T1 as-spun fibers, E is 6.2 GPa, and it decreases for as-spun fibers collected from the same polymer but having a higher linear density. E is respectively 4.9 and 4.5 GPa for the T1 as-spun fibers from batches 95LA and 90LA. The lowest value of E, 4.1 GPa, was calculated for the T1 as-spun fibers from the 80LA batch. After the initial linear region, the stress-strain curves show a distinctive yield point, which occurs at a lower stress if the TMC content in the fiber sample is higher. After the yield point, the specimens undergo a high strain without much stress. This region of the curve corresponds to the natural draw of the fiber (Figure 7c). Increasing the TMC content, end therefore reducing the crystallinity, the as-spun fibers showed higher extensibility, meaning that at a given stress, the strain (%) increases and the region of the natural draw becomes gradually broader from 100LA to 80LA. The strain at break, ε, for the T1 roll from the 80LA batch is 84 %, twice as high as the T1 from the 100LA batch, for which it is 43 %. Moreover, at a given composition, ε increased by increasing the linear density, i.e., the fiber diameter, or in other word by increasing the mass throughput Q during the extrusion fibers with higher  were produced.48 Such results represent a trend opposite that previously found.62 While, the tensile strength, σ, is nearly constant, approximately 0.240.25 GPa, and it seems not to depend on composition and linear density (Tables S12-S15). The physical properties of the as-spun fibers are a consequence of the degree of crystallinity achieved during the spinning, which is in turn dictated by the draw down ratio, the draw temperature and the composition. However, the tensile behavior of the as-spun fibers indicates that the chain orientation is not complete and that the fiber structure is not optimal. Further 28 ACS Paragon Plus Environment

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Biomacromolecules

drawing is necessary to remove the region of natural draw and give to the fibers a higher tensile strength and initial modulus. Hot-drawing of 100LA, 95LA and 90LA as-spun fibers Crystallinity and mechanical properties of the as-spun PLLA fibers could be improved by hotdrawing since the polymer chains are unfolded and oriented along the drawing direction by applying an external tensile stress in the direction of the fiber axis. We wanted to evaluate the effect of composition, drawing temperature and drawing speed on the development of crystallinity, thermal and mechanical properties of the as-spun fibers obtained by melt-spinning of the different (co)polymers, whose properties have been discussed above. The as-spun fibers were hot-drawn using the drawing line depicted in Figure 1b. The drawing parameters, i.e., the godets and winder speed, were calculated in order to have a final linear density of 85 deniers for all the fibers samples. The SEM images confirmed that filaments with a similar diameter, approximately 14 m, were obtained after hot-drawing, regardless of the initial linear density of the as-spun yarns (Figure S3). Therefore, the draw ratio achieved during the hot-drawing was dependent on the initial linear density of the as-spun fibers. Specifically, each as-spun fiber bobbin collected from batches 100LA, 95LA and 90LA was hot-drawn at 80 and 110 °C. Two input speeds of 50 and 300 m min-1 were tested at each temperature. The speed of godets G2’ and G3’ was usually 1.1 times the speed of the winder W’ to allow the polymer chains to relax in the last part of the drawing line (Figure 1b). For batch 100LA, when the fibers were hot-drawn at the higher temperature and lower speed, i.e., 110 °C and 50 m min-1, the speed of godets G2’ and G3’ was kept at 1.05 times the speed of the winder W’, since fiber samples for 100LA relaxed quickly at a higher temperature, folding and leading to the rupture of the yarn between G3’ and W’.



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Crystallinity and crystal structure development during hot-drawing To evaluate the development of the crystalline structure, WAXD analysis was performed on all the hot-drawn fiber samples. The data regarding the degree of crystallinity, the crystallite size, the lattice spacing d and the cell parameter a are reported in the Supporting Information (Tables S4-S6). The WAXD patterns were similar to the ones detected for the as-spun fiber samples. A strong peak corresponding to the (200)/(110) plane was observed. After hot-drawing, the position of the main peak shifted to a higher value of 2θ, indicating that a reduction in the dimensions of the crystalline lattice had occurred. For the 100LA hot-drawn fibers, the main peak was observed to vary in the 2θ range of 16.416.7°. As a consequence, d was in the range of 0.540– 0.532 nm, and the cell parameter a was in the range of 1.08–1.06 nm. Thus, the crystalline ’form of the 100LA as-spun fibers tends to evolve into the more stable -form under the hotdrawing conditions, especially at a higher drawing temperature (Table S4). A similar trend was observed for the 95LA and 90LA fibers. After hot-drawing, the main peak position shifted to a higher value of 2θ, 16.216.5°, and the lattice dimension was found to be smaller than the corresponding as-spun fibers but larger than the 100LA hot-drawn fibers. The value of d was in the range 0.5460.536 nm and a between 1.091.07 nm (Tables S5-S6), values compatible with a crystalline form intermediate between the ’ and the . Therefore, during the hot-drawing process, as the fibers were subjected to drawing stress at temperature above the Tg, thermal- and stress-induced crystallization were promoted. As a consequence, changes in crystallinity occurred: the initial crystal structure was transformed, and the crystalline phase developed in a more stable form. Hence, the degree of crystallinity generally increased (Figure 9a,d,g). However, when the 100LA as-spun fibers were hot-drawn at 80 °C, a reduction in the degree of crystallinity with respect to the native as-spun samples was observed for T1 and T2, while 30 ACS Paragon Plus Environment

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Biomacromolecules

only a slight increase was calculated for T3 and T4. In contrast, the 100LA fibers with a higher degree of crystallinity of up to 66 % (T1c, Table S4) were obtained after hot-drawing at 110 °C. The enhancement of the degree of crystallinity increased by increasing the draw ratio, i.e., T1 through T4, while the input speed had a negligible effect on the crystallinity development of the 100LA fibers. Thus, the thermal-induced crystallization has a great effect on the structure development of the 100LA fibers, and 80 °C was found to be too low of a drawing temperature, since the applied stress led to a decrease in crystallinity (Figure 9a, Table S4). In contrast, despite an overall lower degree of crystallinity than the 100LA fibers, the hotdrawing of 95LA and 90LA fibers always led to an increase in the degree of crystallinity, and this increase was generally higher than that achieved for the 100LA fibers (Figure 9a,d,g and Tables S3-S6).



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Figure 9. Plots of the development of the degree of crystallinity calculated from WAXD (a, d, g) of tensile strength (b, e, h) and of Young’s modulus (c, f, i) of each as-spun yarn after hotdrawing under different conditions (Ta: T = 80 °C, input speed = 50 m min-1; Tb: T = 80 °C, input speed = 300 m min-1; Tc: T = 110 °C, input speed = 50 m min-1; and Td: T = 110 °C, input speed = 300 m min-1) for 100LA (top), 95LA (middle) and 90LA (bottom) (co)polymers. Degrees of crystallinity up to 58 % and 51 % were respectively determined for the 95 and 90LA hot-drawn fibers, and as observed for the 100LA batch, a higher drawing temperature and draw ratio led to fibers with higher crystallinity. A temperature of 110 °C was also found to be a good 32 ACS Paragon Plus Environment

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Biomacromolecules

drawing temperature for fibers containing up to 10 mol % of TMC. In contrast, with the 100LA fibers, at a given drawing temperature, the hot-drawing of 95 and 90LA fibers at slower input speed, 50 m min-1, led to fibers with a higher crystalline content than from hot-drawing at a faster input speed of 300 m min-1 (Figure 9d,g, Tables S5-S6). The TMC units along the polymeric chains probably slow down the crystallization rate and at a high temperature and low speed, the polymer chains have higher energy and more time to organize in the crystalline phase. The degree of crystallinity of T1 fibers hot-drawn under different conditions was overall higher than the others, suggesting that the initial degree of the crystallinity has an effect on the attainable crystallinity and for all the copolymers, fibers with the highest degree of crystallinity were obtained when T1 as-spun yarn was hot-drawn at 110 °C and the lower input speed (T1c, Figure 9d,g, Table S4-S6). However, the T4 as-spun fibers, having the least initial degree of crystallinity and the highest linear density, showed a substantial increase in the degree of crystallinity after hot-drawing at 110 °C (Figure 9a,d,g and Tables S4-S7). The higher draw ratio undergone by the T4 fibers could probably induce a greater development of the crystallinity during the hot-drawing. Therefore, the linear density, i.e., filament morphology, and crystallization of the as-spun fibers could be properly controlled by modulating the draw down ratio during the melt-spinning, and then by a second stage of hot-drawing, it was possible to further draw and orient the fibers, reducing the diameter and promoting additional crystallization. The attainable crystallinity after the second stage was the result of the material properties, such as composition, the initial degree of crystallinity, and the processing parameters that affected the structure development in this order: drawing temperature > draw ratio > input speed. The hot-drawing also induced changes in the crystallite size (Tables S4-S6). Specifically, the 100LA and 95LA hot-drawn fibers had a crystallite size in the range of 610 nm, thus smaller than the size calculated for the native as-spun samples. However, the crystallite size of the



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90LA fibers increased after hot-drawing, achieving similar dimensions to those calculated for the 100LA and 95LA fibers. A crystallite size in the range 7–12 nm, depending on the drawing temperature and degree of crystallinity, has been previously reported for PLLA fibers.38, 45 The observed alteration of the crystallite size with respect to the native as-spun sample could indicate that during hot-drawing, the applied stress led to a deformation and breaking of the crystallites and/or induced changes in the crystalline domains, probably into a more fibrillar morphology. As a consequence, the length of the long-range order can be reduced after hotdrawing. However, by increasing the drawing temperature from 80 to 110 °C, larger crystallites were generally obtained, indicating that at this temperature, the long-range order as well as the degree of crystallinity improved as the fibers were drawn. Thermal properties after hot-drawing The thermal behavior of the fibers after the hot-drawing was also investigated. In Figure 10, DSC thermograms of T1 fibers hot-drawn at 300 m min-1 at low and high speed are compared for the three polymers. All the data relative to the first DSC heating run are summarized in the Supporting Information (Tables S8-S11). However, data regarding the degree of crystallinity, as previously speculated, should be taken with special care since a double-melting behavior was generally observed.


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Figure 10. DSC thermograms reported for the first heating run of 100LA, 95LA and 90LA T1 hot-drawn fibers at 300 m min-1 at both drawing temperatures, 80 °C (a) and 110 °C (b). The thermograms of the first heating run of the 100LA hot-drawn fibers displayed a melting point of approximately 180 °C, a higher temperature than that of the corresponding as-spun samples. ΔHm also increases after hot-drawing, confirming a higher degree of crystallinity, with values approximately 60 % greater than those of the as-spun samples (Table S8). Tg was also observed in a temperature range between 4855 °C. Neither relaxation nor a cold crystallization peak were ever detected for the 100LA hot-drawn fibers, indicating that during the hot-drawing, the polymer chains had a high relaxation rate and could set and arrange in a stable conformation (Figure 10), regardless of the temperature and input speed. The absence of a cold crystallization peak in the DSC run of the hot-drawn PLA fibers has been previously observed.56 In the cooling run beside the glass transition, a crystallization peak on cooling the melt at approximately 100 °C (ΔHc of approximately 20 J g-1) was observed for all the 100LA hot-drawn fiber samples. The thermograms of the 95LA hot-drawn fibers showed the presence of a melting peak at a temperature approximately 166 °C (Figure 10). The ΔHm increased slightly with respect to the as-spun fiber samples. The overall degree of crystallinity after hot-drawing was approximately 40 % (Table S9). Interestingly, the thermograms of the 95LA hot-drawn fibers at 110 °C beside the melting peak only displayed the glass transition (Figure 10b). In contrast, when 95LA fibers were hot-drawn at 80 °C, small relaxation and cold crystallization peaks were observed in some cases right after the Tg. However, ΔHrel and ΔHc were decreased with respect to the as-spun fibers (Table S9). In the cooling run, only the glass transition was observed for the 95LA hotdrawn fibers. 90LA hot-drawn fibers had Tm in the range 156–160 °C, with a degree of crystallinity of approximately 30 % thus, increased with respect to the as-spun fibers (Figure 10, Table S10). Relaxation and cold crystallization peaks were generally observed in the thermograms of the 35 ACS Paragon Plus Environment


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90LA hot-drawn fibers at 80 °C, and a small relaxation peak could also be observed after hotdrawing at 110 °C (Figure 10b, Table S10). For the 90LA hot-drawn fibers, only the glass transition was observed in the cooling run. Therefore, the DSC data also confirmed that the degree of crystallinity was a function of the (co)polymer compositions and it increased during the hot-drawing process. The crystallization kinetics was however slower for fibers having a high TMC content, especially at a drawing temperature of 80 °C. The presence of small relaxation peaks after hot-drawing indicated that for the 95LA and 90LA fibers, the polymer chains were still blocked in an unstable conformation in the amorphous region. Thus, contrary to 100LA fibers, the processing conditions were in some cases not sufficient to allow a complete chain relaxation or a full structure development for fibers containing TMC up to 10 mol%. Tensile properties after hot-drawing As a consequence of the further orientation and crystallization, the tensile properties of the hotdrawn fibers were increased in all cases with respect to the native as-spun fibers; the tensile strength, σ, (Figure 9b,e,h) and Young’s modulus, E, (Figure 9c,f,i) were all higher after hotdrawing, while the strain at break, ε, decreased. Moreover, the stress-strain curves showed that the region of the natural draw ratio was removed. Representative stress-strain curves of T1 fibers hot-drawn at 300 m min-1 at both low and high speeds are shown for the three polymers in Figure 11. Tensile test data relative to all the hot-drawn fibers are summarized in the Supporting Information (Tables S12-S15).


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Figure 11. Specific stress-strain curves of 100LA, 95LA and 90LA T1 hot-drawn fibers at 300 m min-1 and at both drawing temperatures of 80 °C (a) and 110 °C (b). The 100LA hot-drawn fibers had the highest overall σ and E (Figure 9b,c), thus confirming that the mechanical properties strongly depended on composition and crystallinity. The higher molecular weight of the fibers (Table S2, Figure 5) also probably contributed to the better mechanical performance. The tensile properties of the 100LA fibers after hot-drawing were always improved with respect to the as-spun samples. This was also true for the 100LA T1 and T2 fibers, whose crystallinity after hot-drawing at 80 °C was apparently lower than the respective as-spun samples (Figure 9a). This result could be a further indication that the change in crystallinity and reduction in crystallite size should be due to a morphology modification, probably toward a more load-bearing fibrillar structure.63 σ was in the range 0.40–0.61 GPa, which is about twice that of the as-spun fibers, and its increase became more pronounced by increasing the draw ratio, i.e., from T1 through T4. Temperature also has a similar effect; σ values achieved after hot-drawing at 110 °C were higher than the values obtained at 80 °C (Figure 11). The values obtained for fibers hot-drawn at 110 °C and the higher input speed were slightly higher than the values obtained at the lower speed (Figure 9b, Table S12). E was in the range 6.9-8.4 GPa, and an analogous trend with temperature was found: E of the 100LA fibers hot-drawn at 110 °C was generally higher than for fibers hot-drawn at 80 °C, and 37 ACS Paragon Plus Environment


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contrary to σ, a higher E was calculated for fibers hot-drawn at the slower input speed, i.e., 50 m min-1 (Figure 9c). ε decreased significantly after hot-drawing to 17–25 % (Figure 11), and the higher the draw ratio was, the more it decreased (Table S12). In particular, the 100LA T3 and T4 fibers achieved the best mechanical properties after hot-drawing at 110 °C (Figure 9b,c). This confirms that a higher draw ratio promoted greater orientation and stress-induced crystallization, which, combined with thermally induced crystallization, led to a better mechanical response. Fibers with 5 mol% of TMC content achieved values of σ varying in the range of 0.36–0.61 GPa (Figure 9e) when hot-drawn and, thus, were similar to the value obtained for the 100LA fibers. A higher value of σ was reached when 95LA fibers were hot-drawn at 110 °C and at a higher draw ratio (Figure 11). Interestingly, the 95LA T1 fibers had similar values of σ, and E was the same for all of them, with a value of 6.5 GPa (Figure 9e,f), even though they were hotdrawn at a different input speeds and temperatures. E was determined to be in the range 6.1– 7.4 GPa for 95LA T2-T4 fibers with slightly higher values for fibers hot-drawn at 110 °C (Figure 9f). ε also decreased from a range of 5662 % to a range of 1527 % as a consequence of the hot-drawing stage (Table S13). The 90LA fibers, because of their higher TMC content and lower degree of crystallinity had σ in the range of 0.37–0.49 GPa, E in the range of 5.8–6.7 GPa and ε decreased from a range of 69–74 % to a range of 17–27 %(Figure 9h,i, Table S14). However, the drawing temperature and speed, which determine the fiber crystallinity, did not influence the end tensile properties to the same extent as observed for the 100LA and 95LA batches (Figure 11). The lower improvement in σ and E achieved by hot-drawing 90LA could be mainly due to a higher degree of chain orientation the fiber undergo during further drawing rather than to a significant development of the crystallinity.


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Hot-drawing of the 80LA as-spun fibers The hot-drawing of the 80LA as-spun fibers at high temperature was not possible since the fibers stuck to the godets resulting in the breaking of the yarn both at high and low speeds. However, hot-drawing at a temperature of up to 110 °C was possible for the T1 as-spun fibers, having the lowest linear density and the highest degree of crystallinity. The 80LA T1 as-spun fibers were hot-drawn at a fixed input speed of 300 m min-1 at varying temperatures of 70, 80, 90, 100 and 110 °C to evaluate the effect of the hot-drawing temperature on the development of crystallinity and how this determined the thermal and mechanical properties. WAXD analysis, DSC and mechanical tests were carried out. The data are summarized in the Supporting Information (Tables S7, S11, S15). In Figure 12, the degree of crystallinity and the crystallite size are plotted against the drawing temperature.

Figure 12. Plot of the degree of crystallinity (%) and crystallite size (nm) as determined by WAXD analysis for 80LA T1 hot-drawn fibers at different temperatures and at an input speed of 300 m min-1 (Table S7). The degree of crystallinity increased when the drawing temperature was raised from 70 to 80 °C, reaching the highest value of 42 %, while when the temperature was increased above 80 °C, the degree of crystallization slightly decreased and remained roughly constant at 39 ACS Paragon Plus Environment


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approximately 30 % (Figure 12). Therefore, the degree of crystallization of the 80LA T1 asspun fibers under hot-drawing at an input speed of 300 m min-1 had the highest value at 80 °C. Larger crystallites are formed by increasing the temperature, with the size growing from 5.8 to 8.9 nm (Figure 12). An optimum drawing temperature of 70 °C was observed for fibers made of L-lactide and -caprolactone (CL) copolymers containing 20 mol% of CL.6 The data obtained from the stress-strain curve regarding the tensile properties confirmed the crystallinity trend (Table S15). The highest tensile properties were indeed calculated for fibers hot-drawn at 80 °C and 90 °C, for which the value of σ was 0.33 GPa and that of E was 4.9 GPa (T1b and T1f, Table S15). However, good values of σ and E, 0.29 GPa and 4.3 GPa, respectively, were also determined for fibers hot-drawn at 100 °C (T1e, Table S15). Lower mechanical properties were instead achieved after hot-drawing at either a lower or higher temperature (T1g, T1d, Table S15). It was found that 80 °C was a good drawing temperature. The hot-drawing of 80LA as-spun fibers that had a higher linear density than T1, i.e., T2-T4, was also carried out at 80 °C and an input speed of either 50 or 300 m min-1. The results regarding the crystallinity and thermal and mechanical properties are also compared in the Supporting Information (Tables S7, S11, S15). As previously observed for the other (co)polymers, the hot-drawing of 80LA as-spun fibers induced changes in the crystalline structure with a consequent increase in the degree of crystallinity. A slight shift of the main reflection peak relative to the (200)/(110) plane was observed in the WAXD scans. The lattice parameters were found to be smaller with values compatible with the ’-form. The overall degree of crystallinity of the 80LA fibers is lower than for the other fibers due to the higher TMC content. However, regardless of the low drawing temperature, the increase in the degree of crystallinity attained by the 80LA as-spun fibers during hot-drawing is more significant than for the other (co)polymers. A degree of crystallinity of up to 42 % was reached (T2a, Table S7) and the amorphous 80LA T4 as-spun fibers achieved 40 ACS Paragon Plus Environment

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a degree of crystallinity up to 26 % after hot-drawing (T4a, Table S7). The crystallite size also increased from 4-6 nm up to 9 nm (Table S7). Although the Tg and the Tm did not change significantly after the hot-drawing, the DSC results confirmed that a higher degree of crystallinity was achieved after hot-drawing of the 80LA fibers (Table S11). The heating scans always showed the presence of a relaxation peak, confirming that by increasing the amount of TMC, the extent of the amorphous phase and the relaxation rate decreased. As a consequence of having the lowest molecular weight and degree of crystallinity, the mechanical properties of the 80LA fibers were lower than those of the 95LA and 90LA fibers. However, the hot-drawing at a lower temperature led to fibers with quite good tensile properties, with values of σ and E in the range 0.320.36 GPa and 4.9–5.5 GPa, respectively. E increased by increasing the draw ratio, while the value of ε decreased being however higher than the values obtained for the 100-90LA fibers (Table S15). The data obtained for the 80LA fibers showed that the as-spun fibers with quite a large amount of the soft TMC monomeric units and very low crystallinity after melt-spinning could still be oriented and develop a certain degree of crystallinity under hot-drawing, with a consequent improvement in their mechanical features.

CONCLUSION We have assessed the spinnability of four LA/TMC copolymers with a TMC content up to 18 mol% by high-speed melt-spinning, and we have elucidated how the structural and morphological features of the as-spun fibers could be precisely modulated by governing the processing conditions. The spinning of the fibers above their Tg allowed their crystallinity to be developed as a function of their composition and the draw down ratio into a crystalline form resembling the ’-form of PLLA. The good spinnability and the capability of tailoring the fiber properties were also due to the good control achieved during the polymerization. Indeed, 41 ACS Paragon Plus Environment


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(co)polymers with high molecular weights and narrow dispersity were synthesized, and their microstructure, crystallinity and thermal properties were functions of the TMC content. A second stage of hot-drawing was carried out to further orient the as-spun fibers, reducing the diameter, promoting additional crystallization and improving the mechanical properties. The unit cell parameters decreased, revealing that the crystalline structure tended to evolve to the more stable -form, and the crystallite size changed. This probably indicated that a more fibrillary morphology was developed during the second stage. The overall degree of crystallinity, the thermal features and the tensile properties all decreased by increasing the amount of TMC. Meanwhile, at a fixed composition, both thermal- and stressinduced crystallization contributed to the structure development during hot-drawing, although the former had a more important role; i.e. fibers hot-drawn at a higher temperature generally had a higher degree of crystallinity and better mechanical behavior. From our results, the fact also emerged that the presence of TMC units along the polymeric chains slow down the crystallization kinetics and the relaxation rate during the hot-drawing, causing the drawing speed to have an effect on the properties of the fibers containing TMC. Moreover, at a given composition and higher drawing temperature, the attainable final crystallinity was the highest for fibers that achieved a higher degree of crystallinity during melt-spinning, while the best tensile properties were achieved for fibers hot-drawn at a higher draw ratio. The good mechanical properties obtained for the 90LA and 80LA fibers suggested that the further drawing undergone by the fibers also probably led to chain orientation in the amorphous phase, greatly contributing to the mechanical performance improvement. The final fiber properties were the result of composition, initial linear density and degree of crystallinity, as well as the parameters selected during the hot-drawing process. By adjusting the melt-spinning variables and therefore the maximum achievable crystallization from the melt, followed by governing the hot-drawing parameters, it was possible to fabricate fibers with


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high crystallinity and good mechanical properties, even for copolymers having a consistent amount of TMC.

SUPPORTING INFORMATION DSC data of the as-polymerized materials, as-spun and hot drawn fibers (Table S1, S8-S11), homo- and heterosequence signals in the 13C NMR spectra (Table S2), linear density values of the as-spun fibers (Table S3), WAXD scans images and data of the as-spun and hot-drawn fibers (Figure S1, S2 and Table S4-S7), mechanical properties data of the as-spun and hotdrawn fibers (Table S12-S15) and SEM pictures of hot-drawn fibers are included in the Supporting Information.

ACKNOWLEDGMENTS The authors acknowledge the financial support from the Swedish Foundation for Strategic Research (RMA15-0010 and SM14-0011). T.F. acknowledges the Foundation Blanceflor Boncompagni Ludovisi.

REFERENCES 1. 2. 3. 4. 5. 6. 7. 8.

Hakkarainen, M.; Finne-Wistrand, A., Polylactide. In Handbook of Engineering and Speciality Thermoplastics, John Wiley & Sons, Inc.: 2011; pp 349-376. Ramot, Y.; Haim-Zada, M.; Domb, A. J.; Nyska, A., Biocompatibility and safety of PLA and its copolymers. Adv. Drug Delivery Rev. 2016, 107 (Supplement C), 153-162. Nair, L. S.; Laurencin, C. T., Biodegradable polymers as biomaterials. Prog. Polym. Sci. 2007, 32 (8), 762-798. Middleton, J. C.; Tipton, A. J., Synthetic biodegradable polymers as orthopedic devices. Biomaterials 2000, 21 (23), 2335-2346. Eling, B.; Gogolewski, S.; Pennings, A. J., Biodegradable materials of poly(l-lactic acid): 1. Melt-spun and solution-spun fibres. Polymer 1982, 23 (11), 1587-1593. Penning, J. P.; Dijkstra, H.; Pennings, A. J., Preparation and properties of absorbable fibres from l-lactide copolymers. Polymer 1993, 34 (5), 942-951. Fambri, L.; Pegoretti, A.; Fenner, R.; Incardona, S. D.; Migliaresi, C., Biodegradable fibres of poly(l-lactic acid) produced by melt spinning. Polymer 1997, 38 (1), 79-85. Mezghani, K.; Spruiell, J. E., High speed melt spinning of poly(L-lactic acid) filaments. J. Polym. Sci., Part B: Polym. Phys. 1998, 36 (6), 1005-1012. 43 ACS Paragon Plus Environment


9.

10.

11.

12.

13.

14.

15.

16.

17.

18. 19.

20.

21. 22.

23. 24.

Page 44 of 48

Schmack, G.; Tändler, B.; Vogel, R.; Beyreuther, R.; Jacobsen, S.; Fritz, H. G., Biodegradable fibers of poly(L-lactide) produced by high-speed melt spinning and spin drawing. J. Appl. Polym. Sci. 1999, 73 (14), 2785-2797. Wang, C.; Hsiue, T.-T., Core–shell fibers electrospun from phase-separated blend solutions: fiber formation mechanism and unique energy dissipation for synergistic fiber toughness. Biomacromolecules 2017 18 (9), 2906-2917. Mader, M.; Jérôme, V.; Freitag, R.; Agarwal, S.; Greiner, A., Ultraporous, compressible, wettable polylactide/polycaprolactone sponges for tissue engineering. Biomacromolecules 2018 19 (5), 1663-1673. Huan, S.; Liu, G.; Cheng, W.; Han, G.; Bai, L., Electrospun poly(lactic acid)-based fibrous nanocomposite reinforced by cellulose nanocrystals: impact of fiber uniaxial alignment on microstructure and mechanical properties. Biomacromolecules 2018 19 (3), 1037-1046. Lu, H. H.; Cooper, J. A.; Manuel, S.; Freeman, J. W.; Attawia, M. A.; Ko, F. K.; Laurencin, C. T., Anterior cruciate ligament regeneration using braided biodegradable scaffolds: in vitro optimization studies. Biomaterials 2005, 26 (23), 4805-4816. Zilberman, M.; Nelson, K. D.; Eberhart, R. C., Mechanical properties and in vitro degradation of bioresorbable fibers and expandable fiber-based stents. J. Biomed. Mater. Res., Part B 2005, 74B (2), 792-799. Laitinen, O.; Törmälä, P.; Taurio, R.; Skutnabb, K.; Saarelainen, K.; Iivonen, T.; Vainionpää, S., Mechanical properties of biodegradable ligament augmentation device of poly(l-lactide) in vitro and in vivo. Biomaterials 1992, 13 (14), 1012-1016. Feng, J.; Zhang, D.; Zhu, M.; Gao, C., Poly(l-lactide) melt spun fiber-aligned scaffolds coated with collagen or chitosan for guiding the directional migration of osteoblasts in vitro. J. Mater. Chem. B 2017, 5, 5176-5188. Chen, Z.; Liu, W.; Zhao, L.; Xie, S.; Chen, M.; Wang, T.; Li, X., Acid-labile degradation of injectable fiber fragments to release bioreducible micelles for targeted cancer therapy. Biomacromolecules 2018 19 (4), 1100-1110. Perepelkin, K. E., Polylactide fibres: fabrication, properties, use, prospects. A review. Fibre Chem. 2002, 34 (2), 85-100. Hiljanen-Vainio, M.; Karjalainen, T.; Seppälä, J., Biodegradable lactone copolymers. I. Characterization and mechanical behavior of ε-caprolactone and lactide copolymers. J. Appl. Polym. Sci. 1996, 59, 1281-1288. Malin, M.; Hiljanen-Vainio, M.; Karjalainen, T.; Seppälä, J., Biodegradable lactone copolymers. II. Hydrolytic study of ε-caprolactone and lactide copolymers. J. Appl. Polym. Sci. 1996, 59, 1289-1298. Fukushima, K., Poly(trimethylene carbonate)-based polymers engineered for biodegradable functional biomaterials. Biomater. Sci. 2016, 4 (1), 9-24. Zhang, Z.; Kuijer, R.; Bulstra, S. K.; Grijpma, D. W.; Feijen, J., The in vivo and in vitro degradation behavior of poly(trimethylene carbonate). Biomaterials 2006, 27 (9), 17411748. Cork, J.; Whittaker, A. K.; Cooper-White, J. J.; Grondahl, L., Tensile properties and in vitro degradation of P(TMC-co-LLA) elastomers. J. Mater. Chem. B 2015, 3, 4406-4416. Grijpma, D. W.; Van, W. P. G. T.; Nijenhuis, A. J.; Pennings, A. J., Copolymers based on cyclic esters and their use as surgical implants., WO Pat., WO9222599A1, 1992. 44 ACS Paragon Plus Environment

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25. Jarrett, P. K.; Casey, D. J., Deformable, absorbable surgical device., US Pat., US5376102A, 1994. 26. Grijpma, D. W.; Pennings, A. J., (Co)polymers of L-lactide, 1. Synthesis, thermal properties and hydrolytic degradation. Macromol. Chem. Phys. 1994, 195, 1633-1647. 27. Ruckenstein, E.; Yuan, Y., Molten ring-open copolymerization of L-lactide and cyclic trimethylene carbonate. J. Appl. Polym. Sci. 1998, 69 (7), 1429-1434. 28. Pospiech, D.; Komber, H.; Jehnichen, D.; Häussler, L.; Eckstein, K.; Scheibner, H.; Janke, A.; Kricheldorf, H. R.; Petermann, O., Multiblock Copolymers of l-Lactide and Trimethylene Carbonate. Biomacromolecules 2005, 6 (1), 439-446. 29. Dobrzynski, P.; Kasperczyk, J., Synthesis of biodegradable copolymers with low-toxicity zirconium compounds. V. Multiblock and random copolymers of L-lactide with trimethylene carbonate obtained in copolymerizations initiated with zirconium(IV) acetylacetonate. J. Polym. Sci., Part A: Polym. Chem., 2006, 44 (10), 3184-3201. 30. Guerin, W.; Helou, M.; Slawinski, M.; Brusson, J.-M.; Guillaume, S. M.; Carpentier, J.-F., Macromolecular engineering via ring-opening polymerization (2): l-lactide/trimethylene carbonate copolymerization - kinetic and microstructural control via catalytic tuning. Polym. Chem. 2013, 4 (13), 3686-3693. 31. Tyson, T.; Finne-Wistrand, A.; Albertsson, A.-C., Degradable porous scaffolds from various L-lactide and trimethylene carbonate copolymers obtained by a simple and effective method. Biomacromolecules 2009, 10 (1), 149-154. 32. Dargaville, B. L.; Vaquette, C.; Peng, H.; Rasoul, F.; Chau, Y. Q.; Cooper-White, J. J.; Campbell, J. H.; Whittaker, A. K., Cross-linked poly(trimethylene carbonate-co-l-lactide) as a biodegradable, elastomeric scaffold for vascular engineering applications. Biomacromolecules 2011, 12 (11), 3856-3869 33. Dargaville, B. L.; Vaquette, C.; Rasoul, F.; Cooper-White, J. J.; Campbell, J. H.; Whittaker, A. K., Electrospinning and crosslinking of low-molecular-weight poly(trimethylene carbonate-co-l-lactide) as an elastomeric scaffold for vascular engineering. Acta Biomater. 2013, 9 (6), 6885-6897. 34. Cork, J.; Whittaker, A. K.; Cooper-White, J. J.; Grondahl, L., Electrospinning and mechanical properties of P(TMC-co-LLA) elastomers. J. Mater. Chem. B 2017, 5 (12), 2263-2272. 35. Li-Jie, J.; Kui-Lin, L.; Bin, H.; Gang, W.; Li-Qing, S.; Yao, W.; Zhong-Wei, G., Study on poly(l-lactide -cotrimethylene carbonate): synthesis and cell compatibility of electrospun film. Biomed. Mater. 2010, 5 (4), 045009. 36. Lim, L. T.; Auras, R.; Rubino, M., Processing technologies for poly(lactic acid). Prog. Polym. Sci. 2008, 33 (8), 820-852. 37. Gupta, B.; Revagade, N.; Hilborn, J., Poly(lactic acid) fiber: An overview. Prog. Polym. Sci. 2007, 32 (4), 455-482. 38. Solarski, S.; Ferreira, M.; Devaux, E., Thermal and mechanical characteristics of polylactide filaments drawn at different temperatures. J. Text. I. 2007, 98 (3), 227-236. 39. Cicero, J. A.; Dorgan, J. R., Physical properties and fiber morphology of poly(lactic acid) obtained from continuous two-step melt spinning. J.Polym. Environ. 2001, 9 (1), 1-10. 40. Kim, I.-H.; Lee, S. C.; Jeong, Y. G., Tensile behavior and structural evolution of poly(lactic acid) monofilaments in glass transition region. Fibers Polym. 2009, 10 (5), 687-693. 45 ACS Paragon Plus Environment


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41. Paakinaho, K.; Ellä, V.; Syrjälä, S.; Kellomäki, M., Melt spinning of poly(l/d)lactide 96/4: Effects of molecular weight and melt processing on hydrolytic degradation. Polym. Degrad. Stab. 2009, 94 (3), 438-442. 42. Schmack, G.; Jehnichen, D.; Vogel, R.; Tändler, B.; Beyreuther, R.; Jacobsen, S.; Fritz, H. G., Biodegradable fibres spun from poly(lactide) generated by reactive extrusion. J. Biotechnol. 2001, 86 (2), 151-160. 43. Schmack, G.; Tändler, B.; Optiz, G.; Vogel, R.; Komber, H.; Häußler, L.; Voigt, D.; Weinmann, S.; Heinemann, M.; Fritz, H. G., High-speed melt spinning of various grades of polylactides. J. Appl. Polym. Sci. 2004, 91 (2), 800-806. 44. Cicero, J. A.; Dorgan, J. R.; Garrett, J.; Runt, J.; Lin, J. S., Effects of molecular architecture on two-step, melt-spun poly(lactic acid) fibers. J. Appl. Polym. Sci. 2002, 86 (11), 28392846. 45. Ghosh, S.; Vasanthan, N., Structure development of poly(L-lactic acid) fibers processed at various spinning conditions. J. Appl. Polym. Sci. 2006, 101 (2), 1210-1216. 46. Persson, M.; Lorite, G. S.; Cho, S.-W.; Tuukkanen, J.; Skrifvars, M., Melt Spinning of Poly(lactic acid) and Hydroxyapatite Composite Fibers: Influence of the Filler Content on the Fiber Properties. ACS Appl. Mater. Interfaces 2013, 5 (15), 6864-6872. 47. Zhang, H.; Bai, H.; Liu, Z.; Zhang, Q.; Fu, Q., Toward High-Performance Poly(l-lactide) Fibers via Tailoring Crystallization with the Aid of Fibrillar Nucleating Agent. ACS Sustainable Chem. . Eng. 2016, 4 (7), 3939-3947. 48. R. Beyreuther, H. Brünig, Modelling of steady state fibre formation process in melt spinning. In Dynamics of Fibre Formation and Processing: Modelling and Application in Fibre and Textile Industry, Springer Berlin Heidelberg: Berlin, Heidelberg, 2007; pp 43143. 49. Hua, J.; Gebarowska, K.; Dobrzynski, P.; Kasperczyk, J.; Wei, J.; Li, S., Influence of chain microstructure on the hydrolytic degradation of copolymers from 1,3-trimethylene carbonate and L-lactide. J. Polym. Sci., Part A: Polym. Chem. 2009, 47 (15), 3869-3879. 50. Kalb, B.; Pennings, A. J., General crystallization behaviour of poly(l-lactic acid). Polymer 1980, 21 (6), 607-612. 51. Nijenhuis, A. J.; Grijpma, D. W.; Pennings, A. J., Highly crystalline as-polymerized poly(llactide). Polym. Bull. 1991, 26 (1), 71-77. 52. Fambri, L.; Migliaresi, C., Crystallization and thermal properties. In Poly(Lactic Acid), John Wiley & Sons, Inc.: 2010; pp 113-124. 53. Carothers, W. H.; Natta, F. J. V., Studies on polymerization and ring formation. iii. glycol esters of carbonic acid. J. Am. Chem. Soc. 1930, 52 (1), 314-326. 54. Androsch, R.; Iqbal, H. M. N.; Schick, C., Non-isothermal crystal nucleation of poly (llactic acid). Polymer 2015, 81 (Supplement C), 151-158. 55. Zhang, J.; Duan, Y.; Sato, H.; Tsuji, H.; Noda, I.; Yan, S.; Ozaki, Y., Crystal modifications and thermal behavior of poly(l-lactic acid) revealed by infrared spectroscopy. Macromolecules 2005, 38 (19), 8012-8021. 56. Sasaki, S.; Asakura, T., Helix distortion and crystal structure of the α-form of poly(llactide). Macromolecules 2003, 36 (22), 8385-8390. 57. Jamshidi, K.; Hyon, S. H.; Ikada, Y., Thermal characterization of polylactides. Polymer 1988, 29 (12), 2229-2234. 46 ACS Paragon Plus Environment

Page 47 of 48 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Biomacromolecules

58. Yuan, X.; Mak, A. F. T.; Kwok, K. W.; Yung, B. K. O.; Yao, K., Characterization of poly(L-lactic acid) fibers produced by melt spinning. J. Appl. Polym. Sci. 2001, 81 (1), 251-260. 59. Wasanasuk, K.; Tashiro, K., Crystal structure and disorder in Poly(l-lactic acid) δ form (α′ form) and the phase transition mechanism to the ordered α form. Polymer 2011, 52 (26), 6097-6109. 60. Wasanasuk, K.; Tashiro, K., Structural regularization in the crystallization process from the glass or melt of poly(l-lactic acid) viewed from the temperature-dependent and timeresolved measurements of FTIR and wide-angle/small-angle X-ray scatterings. Macromolecules 2011, 44 (24), 9650-9660. 61. Solarski, S.; Ferreira, M.; Devaux, E., Characterization of the thermal properties of PLA fibers by modulated differential scanning calorimetry. Polymer 2005, 46 (25), 1118711192. 62. Hossain, K. M. Z.; Parsons, A. J.; Rudd, C. D.; Ahmed, I.; Thielemans, W., Mechanical, crystallisation and moisture absorption properties of melt drawn polylactic acid fibres. Eur. Polym. J. 2014, 53, 270-281. 63. Hoogsteen, W.; Postema, A. R.; Pennings, A. J.; Ten Brinke, G.; Zugenmaier, P., Crystal structure, conformation and morphology of solution-spun poly(L-lactide) fibers. Macromolecules 1990, 23, 634-642.



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For Table of Contents use only

Poly(L-lactide) and poly(L-lactide-cotrimethylene carbonate) melt-spun fibers: structure-processing-properties relationship Tiziana Fuoco,1 Torbjörn Mathisen2 and Anna Finne-Wistrand1*

1) Department of Fibre and Polymer Technology, KTH Royal Institute of Technology, 100 44 Stockholm, Sweden 2) Novus Scientific AB, Virdings allé 2, 754 50 Uppsala, Sweden E-mail: [email protected]


Poly(l-lactide) and Poly(l-lactide-co-trimethylene carbonate) Melt-Spun

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