Synthesis and properties of ω-pentadecalactone-co-δ-hexalactone copolymers: a biodegradable thermoplastic elastomer as an alternative to poly(ε-caprolactone)

Abstract

The copolymerization of ω-pentadecalactone with the little-known δ-hexalactone, a cyclic ester with an identical structure to δ-valerolactone but possessing a methyl pendant group, creates a new kind of low glass transition polyester that shows improved biodegradability in comparison to poly(ε-caprolactone) (PCL). The poly(ω-PDL-co-δ-HL), with ω-PDL molar contents ranging from 39 to 82%, were synthesized using triphenyl bismuth as catalyst and presented a chain microstructure that deviated slightly from a random distribution (R > 0.74) while WAXS measurements proved that only the ω-PDL blocks were able to crystallize. Therefore, the incorporation of δ-HL decreased the crystalline fraction of the poly(ω-pentadecalactone) to 21–44% and the T_ms of the copolymers shifted from 104 °C to temperatures between 53 and 88 °C. On the other hand, the low glass transition temperatures (<−36 °C) of these thermoplastic elastomers allow a rapid crystallization from melt, which prevents physical aging during their storage or application. The degradation rates obtained for the poly(ω-PDL-co-δ-HL), after carrying out in vitro degradation studies in phosphate buffer solution (at 37 °C for 182 days), ranged from 0.0013 to 0.0019 d⁻¹. Thermogravimetric analysis also demonstrated their great thermal stability, since they were completely degraded at temperatures close to 500 °C. These ω-PDL-co-δ-HL copolymers displayed tunable mechanical properties, adjusted according to their composition, with lower elastic modulus values than PCL and a mechanical performance that was quite steady at both room temperature (21 °C) and body temperature (37 °C).

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1. Introduction

Nowadays, one of the challenges of polymer chemistry is synthesizing biodegradable materials with the same properties as petroleum based polymers. Since the 1980s, an important research work has been focused on aliphatic polyesters while the ring-opening polymerization (ROP) of lactones and macrolactones^1–10 has been proven as a useful approach in obtaining copolyesters with tuned topology and microstructure. However, the polyesters most often employed in medical science either display a glassy mechanical behavior with high elastic modulus and rather low elongation at break values, owing to their high glass transition (T_g) and melting temperatures (T_m) (i.e. polylactide (PLA) or polyglycolide (PGA)¹¹), or show an optimal combination of properties but biodegrade at a very slow rate, as is the case of poly(ε-caprolactone) (PCL)^12,13 and the polyethylene-like polymacrolactones.^14,15 These materials are, therefore, clearly inappropriate for numerous clinical uses, where highly elastomeric biodegradable materials are required.

The copolymerization of lactide or glycolide with a complementing monomer that reduces the crystallinity and the melting point is a good strategy to enhance the flexibility and ductility of PLA and PGA, as long as the glass transition of these new materials is kept above room temperature. However, some of the ε-caprolactone, δ-valerolactone, ε-decalactone, trimethylene carbonate or p-dioxanone^16–33 based copolymers frequently suffer from supramolecular rearrangements in their chains, associated with the development of lactide/glycolide crystalline domains, which cause undesirable changes in their properties during storage or biodegradation. Moreover, these lactide or glycolide-rich copolymers usually do not exhibit a suitable mechanical performance at both room temperature (21 °C) and at human body temperature (37 °C) because their mechanical properties are not stable at temperatures above T_g.

On the contrary, linear copolymers with a relatively low concentration of ester groups show low glass transition (<−25 °C) and melting temperatures (<100 °C), thereby ensuring a very efficient processing with thermoplastic techniques. Likewise, these polyesters present high thermal stability and a rapid crystallization from melt, which prevent physical aging. Another important feature of these low T_g copolymers is their attractive mechanical behavior, with high elongation at break values and elastic modulus below 400 MPa. Hence, the synthesis of poly(ε-caprolactone)-like biomaterials with an improved biodegradability is extremely interesting and could meet the requirements for a large number of potential applications in the biomedical field. However, it should also be taken into consideration that the T_m of these polymers should be above human body temperature in order to guarantee their mechanical performance during use. In addition, if possible, this T_m value should preferably be higher than 60–65 °C, the T_m of PCL, so as to assure more stable mechanical properties. The mechanical behavior of the copolymers with low T_g is heavily dependent on the crystalline phase (on their T_m and melting enthalpies), especially at temperatures close to their melting points. Therefore, with the incorporation of another monomer, the ε-caprolactone-rich copolymers^34,35 present lower T_m and melting enthalpies and as a result exhibit lower stress related properties at 37 °C than at 21 °C.³⁴

ω-Pentadecalactone, whose homopolymer is highly crystalline and melts at around 100 °C,³⁶ has been copolymerized with numerous monomers. Thus, several ω-PDL-containing copolyesters including poly(ω-pentadecalactone-co-ε-caprolactone), poly(ω-pentadecalactone-co-p-dioxanone), poly(ω-pentadecalactone-co-δ-valerolactone) and poly(ω-pentadecalactone-co-trimethylene carbonate) have been proposed as alternative bioresorbable materials.^37–47 However, in most of these cases both comonomers cocrystalize. Only, in the case of poly(ω-pentadecalactone-co-ε-decalactone)^48–50 was the crystalline fraction reduced, owing to the racemic stereochemistry of the butyl side chain of this decalactone. Nevertheless, the latter materials were found to be very resistant to hydrolytic degradation because of their blocky distribution of sequences and the steric effect of the ε-DL units.⁵⁰ Therefore, in the this paper δ-hexalactone,^51–54 a six-membered lactone with identical structure to δ-valerolactone but possessing a methyl pendant group (also known as δ-caprolactone and found in heated milk fat), was employed in place of ε-DL with the aim of lowering the steric effect of the ε-DL and at the same time increasing the hydrophilicity of the copolymers.

The poly(ω-pentadecalactone-co-δ-hexalactone) of this study were synthesized using triphenyl bismuth (Ph₃Bi)⁵⁵ as catalyst and characterized using proton and carbon nuclear magnetic resonance spectroscopy (¹H and ¹³C NMR), gel permeation chromatography (GPC) measurements and thermogravimetric analysis (TGA). Their crystallization and melting behaviour was studied by means of differential scanning calorimetry (DSC) and Wide Angle X-Ray Diffraction (WAXRD). In addition, films of the ω-PDL-co-δ-HL copolymers with ω-PDL molar contents ranging from 39 to 82% were prepared for mechanical testing at room temperature (21 °C) and at 37 °C, the working temperatures of these biomaterials. Finally, an in vitro hydrolytic degradation study was also carried out at 37 °C for a period up to 26 weeks in phosphate buffer solution (PBS). Thus, the changes in water absorption, weight loss, macroscopic morphology, crystallinity, phase structure, molecular weight and mechanical properties of the copolymers were monitored.

2. Materials and methods

2.1. Materials

δ-Hexalactone monomer (assay > 98%) was provided by Sigma Aldrich (W316709). ω-Pentadecalactone monomer (assay > 98%) was also supplied by Sigma Aldrich (W284009). The triphenyl bismuth (Ph₃Bi) catalyst was obtained from Gelest. Phosphate buffer solution (PBS) (pH 7.2) was obtained from Fluka Analytical (Sigma Aldrich). For proper evaluation of the behavior of the ω-PDL-co-δ-HL copolymers, a poly(ω-pentadecalactone) (PPDL) homopolymer, synthesized in a previous work,⁵⁰ was also employed. This material had a M_w of 365 kDa with a dispersity (D) of 1.98.

2.2. Synthesis procedure

Statistical copolymers from ω-pentadecalactone and δ-hexalactone at different compositions were synthesized in bulk by one pot-one-step ring-opening polymerizations (ROP). The synthesis reactions were conducted in a flask immersed in a controlled temperature oil bath. In each polymerization, predetermined amounts of the different comonomers at the chosen mass feed ratio were simultaneously added and melted into the flask. The flask was purged for 30 minutes with a nitrogen stream under the surface of the melt. The catalyst (Ph₃Bi) was then added (at 100 : 1 comonomers/catalyst molar ratio) and the magnetic stirrer maintained at 100 rpm. No initiator compound was added to the reaction mix so the catalyst was activated by the ROH provided by the monomers (H₂O and impurities). The ω-pentadecalactone-co-δ-hexalactone polymerizations were carried out at 130 °C over 6 days. Long reaction times were necessary to obtain substantial yields due to the low reactivity of the mixture of monomers. Moreover, at lower reaction temperatures or when less PH₃Bi was added the efficiency of the process decreased and products with low molecular mass were obtained.

After the corresponding period of reaction time the product was dissolved in chloroform and precipitated, pouring the polymer solution into an excess of methanol in order to remove the catalyst impurities and those monomers that had not reacted. Finally the product was dried at room temperature and then subjected to a heat treatment at 140 °C for 1 hour to ensure the complete elimination of any remaining solvent. The polymer sample then weighed, obtaining the yield of the synthesis process which is shown in Table 1.

Table 1 NMR characterization data of the ω-PDL-co-δ-HL copolymers of different compositions synthesized at 130 °C with a monomers/catalyst molar ratio of 100

Sample	Feed molar composition		Composition^a		Yield (%)	Conversion (%)		Mw, kDa	D	Microstructural magnitudes^b
	%PDL	%HL	%PDL	%HL		PDL	HL			lPDL	lHL	R
a Calculated averaging the copolymer molar compositions obtained by ^1H and ^13C NMR.b lPDL and lHL are the PDL and HL number average sequence lengths of the ω-PDL-co-δ-HL copolymers obtained from the average dyad relative molar fraction (PDL-HL) in ^13C NMR. These values are compared to the Bernoullian random number-average sequence lengths obtaining the randomness character value (R) of the different copolymers.
PDL-HL 82	72.9	27.1	82.4	17.6	74.9	80.0	46.0	249.2	2.03	6.67	1.42	0.85
PDL-HL 76	65.5	34.5	75.8	24.2	65.4	71.0	43.0	180.1	1.91	4.91	1.57	0.84
PDL-HL 72	58.8	41.2	71.6	28.4	76.0	85.3	48.2	156.7	1.82	4.30	1.71	0.82
PDL-HL 63	50.2	49.8	62.6	37.4	67.4	77.2	46.5	103.9	1.79	3.58	2.14	0.75
PDL-HL 42	36.7	63.3	42.2	57.8	57.0	62.8	49.9	83.5	1.77	2.36	3.22	0.74
PDL-HL 39	28.0	72.0	38.7	61.3	48.9	62.0	38.2	62.0	1.81	2.22	3.51	0.74

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2.3. Methods

Proton and carbon nuclear magnetic resonance (¹H and ¹³C NMR) spectra were recorded in a Bruker Avance DPX 300 at 300.16 MHz and at 75.5 MHz of resonance frequency respectively, using 5 mm O.D. sample tubes. All spectra were obtained at room temperature from solutions of 0.7 mL of deuterated chloroform (CDCl₃). Experimental conditions were as follows: (a) for ¹H NMR: 10 mg of sample; 3 s acquisition time; 1 s delay time; 8.5 μs pulse; spectral width 5000 Hz and 32 scans; (b) for ¹³C NMR: 40 mg, inverse gated decoupled sequence; 3 s acquisition time; 4 s delay time; 5.5 μs pulse; spectral width 18 800 Hz and more than 10 000 scans. The assignment of the different signals was made employing the tables of structural determination devised by Prestch et al.⁵⁶ The copolymer composition data, average sequence lengths and randomness character in Table 1 were calculated by averaging the values of molar contents and the PDL-HL dyad relative molar fractions that were obtained by means of ¹H and ¹³C NMR spectroscopy. Eqn (1)–(3) (ref. 57) were employed to obtain the number-average sequence lengths (l_i), the Bernoullian random number-average sequence lengths (l_i) and the randomness character (R):where (HL) and (PDL) are the δ-hexalactone and ω-pentadecalactone molar fractions and (PDL-HL) is the PDL-HL average dyad relative molar fraction.

200–300 μm films were prepared by pressure melting at 200 °C followed by water quenching. These were then stored for 24 hours in a fridge (at 0–5 °C), the typical storage temperature for biopolymers. From these films repetitive square samples for the in vitro degradation study (1 × 1 cm²) and repetitive samples for mechanical characterization (10 × 1 cm²) were obtained. The specimens for mechanical testing at 37 °C were stored for another 24 hours at 37 °C before tests were conducted at this temperature. In addition, DSC scans were made at 20 °C min⁻¹ for each polymer sample before mechanical testing, in order to monitor the thermal properties of the specimens.

The mechanical properties were determined by tensile tests with an Instron 5565 testing machine at a crosshead displacement rate of 10 mm min⁻¹. These tests were performed at room temperature (21 ± 2 °C) and at human body temperature (37 °C) following ISO 527-3/1995. The specimens had the following dimensions: overall length = 100 mm, distance between marks = 50 mm, width = 10 mm; and were cut from 200–300 μm thick films. The mechanical properties reported (secant modulus at 2%, yield strength, ultimate tensile strength and elongation at break) correspond to average values of at least 5 determinations. Mechanical testing at 37 °C was conducted in an Instron controlled temperature chamber. The tests were stopped at 300% of strain due to the size limitations of the temperature chamber.

For the in vitro degradation study, square samples (25–35 mg (W₀) (n = 3)) of the different copolymers were placed in Falcon tubes containing phosphate buffer solution (PBS) (pH = 7.2) maintaining a surface area to volume ratio equal to 0.1 cm⁻¹. The samples were stored in an oven at 37 °C. Three samples of each polymer were removed at different times from the PBS and weighed wet (W_W) immediately after wiping the surface with filter paper to absorb the surface water. These samples were air-dried overnight at 37 °C. Then they were weighed again to obtain the dry weight (W_d). Water absorption (%WA) and remaining weight (%RW) were calculated according to eqn (4) and (5). At the end of the degradation study (182 days) the mechanical properties of the poly(ω-pentadecalactone-co-δ-hexalactone) were also measured at 37 °C.

In order to compare the degradation rate of the studied copolymers, the exponential relationship between molecular weight and degradation time for biodegradable polyesters degrading under bulk degradation was used:^58,59

ln M_w = ln M_w0 − K_Mwt (6)

t_1/2 = 1/K_Mw ln 2 (7)

where M_w is the weight-averaged molecular weight, M_w0 is the initial weight-averaged molecular weight, K_Mw is the apparent degradation rate and t_1/2 is the half degradation time (the amount of time required to fall to half the initial value of molecular weight).

The molecular weights of the polymers were determined by GPC using a Waters 1515 GPC device equipped with two Styragel columns (10² to 10⁴ Å). Chloroform was used as eluent at a flow rate of 1 mL min⁻¹ and polystyrene standards (Shodex Standards, SM-105) were used to obtain a primary calibration curve. The samples were prepared at a concentration of 10 mg in 1.5 mL. The reported values are likely to be higher than the actual molecular weights owing to the differences in hydrodynamic volume of the copolymers and polystyrene.

The thermal properties of the polymers were studied on a DSC 2920 (TA Instruments). Samples of 6–9 mg were melted at 140 °C and then quenched to 5 °C at different cooling rates (5, 10 and 20 °C min⁻¹) to study the crystallization process during the cooling. Finally, a scan was also made at 20 °C min⁻¹ from −85 to 140 °C to determine the glass transition temperature (T_g), the melting temperature (T_m) and the heat of fusion or melting enthalpy per gram (ΔH_m) of the samples. During the in vitro degradation study, the DSC samples were heated from 21 °C to 140 °C at 20 °C min⁻¹, immediately after the drying at 37 °C of the polymer samples removed from the degradation medium. After this first scan, the samples were quenched in the DSC and a second scan was made from −85 °C to 140 °C at 20 °C min⁻¹.

Wide angle X-ray diffraction (WAXRD) data were collected on a Bruker D8 Advance diffractometer operating at 30 kV and 20 mA. This device is equipped with a Cu tube (λ = 1.5418 Å), a Vantec-1 PSD detector and an Anton Parr HTK2000 high-temperature furnace. The powder patterns were recorded in 2θ steps of 0.033° in the 10 ≤ 2θ ≤ 38 range, counting for 0.2 s per step, from 30 to 120 °C every 2 °C using a heating rate of 0.16 °C s⁻¹. The degree of crystallinity (X_c) was determined as the ratio of the crystalline peak areas to the total area under the scattering curve, while the average crystal size was obtained employing the Scherrer equation⁶⁰ with a shape factor “k”of 0.90. The latter data were calculated from the two most intense peaks in the diffraction pattern: at 2θ = 21.7° and at 2θ = 24.1°. On the other hand, the melting temperatures (T_ms) estimated are the temperatures until the crystalline phase became stable.

Thermal degradation was studied under nitrogen by means of thermogravimetric analysis (TGA) into a TGA model Q50-0545 (TA Instruments). Samples of 10–15 mg were heated from room temperature to 500 °C at a heating rate (β) of 5 °C min⁻¹, with the heat flow, sample temperature, residual sample weight and its time derivative being continuously recorded. In this temperature range, the polymers degraded completely.

3. Results and discussion

3.1. Synthesis characterization

Table 1 summarizes the characterization data of several poly(ω-PDL-co-δ-HL) with different composition synthesized at 130 °C for 6 days with a monomers/catalyst molar ratio (M/C) of 100. As can be observed, the molecular weights of the copolymers increase proportionally with the ω-PDL content, ranging from 62 to 249 kDa, while their dispersity values (D) are in the range of 1.77 to 2.03. The ω-PDL molar content in this set of copolymers ranges from ∼39% (PDL-HL 39) to 82% (PDL-HL 82) with average sequence lengths of ω-PDL (l_PDL) from 2.22 to 6.67. The consequence of the different reactivity of the comonomers is that the synthesis reactions led to gradient copolymers with a slightly random distribution of sequences (R → 1). The randomness character (R) value is, in all cases, higher than 0.74 while the copolymers richer in ω-PDL showed a more random distribution of sequences (R > 0.82).

As can be seen in Table 1, the conversion of ω-PDL was higher than that of δ-HL in all polymerizations. Therefore, higher yield of reactions and molecular weights were obtained when the feed was richer in ω-PDL. For example, PDL-HL 39, the copolymer with the lowest content of ω-PDL, only 38% of δ-HL became a part of the copolymer after the synthesis reaction and the molar content of this unit failed to reach to 61% (when the feed composition had a 72% molar content of this monomer). The initiating species formed by means of the activation of the catalyst with ROH (H₂O or impurities provided by the monomers or the catalyst) may prefer the incorporation of ω-PDL rather than δ-HL units. Thus, the consumption of ω-PDL was faster than that of δ-HL owing to the low reactivity of the latter. This is different from what was observed in the study of the ω-pentadecalactone-co-ε-decalactone copolymers. In accordance with our own results⁵⁰ and the reports of Jasinska-Walc and Duchateau,^48,49 the reactivity of ε-DL was higher than that of ω-PDL leading to a blockier distribution of sequences (R < 0.49). Thermodynamic polymerizability has been shown to increase along with the ring size in the cyclic small-medium lactones.⁶¹ So, it may be possible that the smaller ring strain of δ-HL (six-membered ring) when compared to ε-DL (seven-membered ring) explains the contrasting reactivity of these two lactones in presence of ω-PDL, whose polymerization route is, in contrast, mainly entropically driven.⁶²

3.2. NMR characterization

Fig. 2 and 3 present the ¹H and ¹³C NMR spectra of the ω-pentadecalactone-co-δ-hexalactone copolymer with 63% of ω-PDL (PDL-HL 63). In the ¹H NMR spectrum, the calculus of the molar composition was made based on the signals of the ω-PDL methylene bonded to the ester group, which appears centered at 4.09 ppm, with regard to the δ-HL methine at around 4.95 ppm. The rest of the peaks (see scheme in Fig. 1) are assigned to the other hydrogens of the copolymer in the table of Fig. 2, although some signals appear overlapped.

Fig. 1 Scheme of the units of ω-pentadecalactone (above) and δ-hexalactone (below).

Fig. 2 ¹H NMR spectrum of PDL-HL 63 with the assignment of the different signals to the hydrogens of the ω-PDL-co-δ-HL.

Fig. 3 shows the ¹³C NMR spectrum of the ω-PDL-co-δ-HL copolymer. The different signals appearing between 19 and 175 ppm of chemical shift (δ) are assigned in Table 2 to the different carbons numbered in Fig. 1. Hence, the molar composition of these copolymers can be easily determined by comparing the areas under the peaks due to the ω-PDL and δ-HL carbons. On the other hand, the signals of the ¹³C NMR spectrum at 172.8 and at 173.5 ppm, which belong to the carbonyl carbons of the δ-hexalactone and ω-pentadecalactone units respectively, show sequence sensitivity. The peak at 172.83 split into two shoulders that can be assigned to the HL-HL and -PDL dyads (underlining is used to emphasize that the analysed nuclei belong to that unit). Furthermore, the PDL-PDL dyad appears at 173.93 ppm whereas the signals found at 173.36 ppm are related to the -HL dyad. However, the carbonyl peaks were not well-resolved and it is a better idea to calculate the dyad relative molar fractions from another region, in which the four PDL-HL dyads are clearly distinguishable. The signals from the region enlarged in Fig. 3 and centered at 64.40 and 70.25 ppm, can be assigned to the ω-PDL methylene bonded to the ester group (carbon 1) and the δ-HL methine (carbon 1′), respectively. The HL-HL and -PDL dyads, along with the -HL and PDL-PDL dyads, appear from left to right at 70.31 and 70.07 ppm, and at the peak at 64.40 ppm, at around 64.50 and 64.34 ppm.

Fig. 3 ¹³C NMR spectrum of PDL-HL 63. The region enlarged shows the signals of the ω-PDL methylene bonded to the ester group (carbon 1) and the δ-HL methine (carbon 1′), respectively. They show dyad sensitivity.

Table 2 Assignment of the different signals to the carbons of the ω-PDL-co-δ-HL

Number of C	C ω-PDL	C δ-HL
	δ (ppm)	δ (ppm)
1	64.34 PDL-PDL	70.07 -PDL
	64.50 -HL	70.31 HL-HL
2	28.6–29.6	35.19
3	25.90	20.79
4	28.6–29.6	33.90–34.62
5		172.82 HL-HL
		172.85 -PDL
6		19.86
7
8
9
10
11
12
13	24.98
14	34.36
15	173.32 -HL
	173.36 -HL
	173.42 -HL
	173.93 PDL-PDL

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3.3. Thermal properties

In this section the crystallization and melting behavior of the copolymers in Table 1 were studied. First of all, several cooling treatments at different rates (5, 10 and 20 °C min⁻¹) were conducted in the DSC. The crystallization temperatures (T_c) and the enthalpy of crystallization (ΔH_c) of each material are shown in Table 4.

As can be seen in Table 3, the differences between the respective ΔH_c from the same polymer were small (particularly for those with higher contents of δ-HL), although it was noted that at the slowest cooling rate, the ΔH_c was slightly larger than those corresponding to the other cooling treatments. Therefore, it can be stated that the ω-PDL chains crystallized from the melt very fast and the ΔH_c of the PPDL homopolymer and their copolymers were almost independent of the cooling rate. Likewise, the crystallization peaks of each polymer appeared at a higher temperature as the cooling rate decreased. This is due to the fact that at lower cooling rates the crystal nuclei have more time to develop.

Table 3 Crystallization enthalpies and temperatures obtained for the ω-PDL-co-HL copolymers at cooling rates in the 5–20 °C min⁻¹ range

Sample	5 °C min^−1		10 °C min^−1		20 °C min^−1
	ΔH (J g^−1)	Tc (°C)	ΔH (J g^−1)	Tc (°C)	ΔH (J g^−1)	Tc (°C)
PPDL	129.2	76.0	127.5	72.2	123.8	63.2
PDL-HL 82	100.3	64.6	95.5	61.1	86.3	55.7
PDL-HL 76	87.0	62.7	85.5	58.8	76.9	52.7
PDL-HL 72	81.0	56.4	77.9	53.6	71.4	49.1
PDL-HL 63	67.3	52.7	66.4	49.6	64.8	45.0
PDL-HL 42	55.0	47.0	53.6	43.3	51.1	38.3
PDL-HL 39	39.5	34.7	38.4	31.3	35.8	28.3

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The DSC heating curves obtained after the cooling treatments were virtually identical for each polymer, all of them having practically the same melting temperatures and enthalpies (T_m and ΔH_m) after the cooling process. Fig. 4 shows the DSC curves of scans made from −85 to 140 °C at 20 °C min⁻¹ while the data obtained from them (the glass transition temperature (T_g) with its associated heat capacity (ΔC_p) and the melting enthalpies and temperatures) are summarized in Table 4 along with the T_m, the crystal fraction (χ_c) and the average crystal size (c.s.) from Wide Angle X-ray Scattering (WAXS).

Fig. 4 DSC heating curves at 20 °C min⁻¹ from −85 to 140 °C of the different ω-PDL-co-δ-HL copolymers.

Table 4 Calorimetric and X-ray diffraction data of the ω-PDL-co-δ-HL copolymers

Sample	lPDL	Tm^a (°C)	Xc (%)	c.s (nm)	ΔHm (J g^−1)	Tm^b (°C)	Tg (°C)	ΔCp (J g^−1 °C^−1)
a Obtained by WAXS. This is the crystalline phase-stable temperature.b Obtained from a DSC at 20 °C min^−1 from −85 to 140 °C.
PPDL	—	104.0	53.8	27	136.1	104.3	−35.6	0.16
PDL-HL 82	6.67	78.0	41.1	24	107.1	87.5	−38.6	0.19
PDL-HL 76	4.91	76.0	44.1	19	101.2	85.0	−37.4	0.21
PDL-HL 72	4.30	70.0	38.4	17	94.8	76.3	−40.2	0.27
PDL-HL 63	3.58	65.0	35.7	17	84.5	74.1	−42.5	0.28
PDL-HL 42	2.36	59.0	27.6	12	68.9	63.2	−41.8	0.45
PDL-HL 39	2.22	50.0	21.1	17	51.2	53.1	−42.5	0.48

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As can be noted, the PPDL homopolymer and the ω-PDL rich copolymers showed broader melting peaks and higher crystallization capability in comparison to other copolymers such as PDL-HL 39, the poly(ω-PDL-co-δ-HL) that presented the lowest melting peak. The associated ΔH_ms values of these δ-HL copolymers, with ω-PDL contents from 39 to 82% and average sequence lengths of ω-PDL in the range of 2.22–6.67, were slightly larger than those of the crystallization peaks (from Table 3) and were between ∼51 to 107 J g⁻¹. Their calculated crystalline fraction ranged from 21.1 to 44.1%, values lower than those of PPDL (with a χ_c of 53.8%, a ΔH_m of 136.1 J g⁻¹ and a T_m at 104 °C). On the other hand, the melting temperature values obtained by WAXS for the poly (ω-PDL-co-δ-HL) (from 50 to 78 °C) were consistent with their corresponding T_ms obtained by DSC (from 53 to 88 °C), although they are somewhat lower. With regard to their glass transition behaviour, the T_gs shifted to lower values when the δ-HL content was raised and were more easily distinguished (with higher ΔC_p values) as the amorphous phase increased. The measured values were between −37 and −43 °C, temperatures higher than those of the ω-pentadecalactone-co-ε-decalactone copolymers.

To provide another framework for comparison, some of the melting results were contrasted with those obtained for the ω-pentadecalactone-co-ε-decalactone copolymers, in which ε-DL, with an identical structure to ε-caprolactone but with a butyl pendant group, also lowered the crystallization capability of ω-PDL. Fig. 5 shows the melting enthalpies and temperatures of the copolymers in this work (in bold squares) along with those of the poly(ω-PDL-co-ε-DL).⁵⁰ As can be seen the ΔH_m values and the crystallinity degree (not shown) of both kinds of copolymers are related to the ω-PDL content. This is quite unexpected because the poly(ω-PDL-co-ε-DL) have a blockier character (0.38 < R < 0.49), and consequently should exhibit an increased crystalline fraction with the same composition. Nevertheless, the T_ms of the δ-HL based copolymers were certainly lower than those measured for the ω-PDL-co-ε-DL copolymers. This makes more sense, because it has been reported that random copolymers possess lower T_ms than the blocky copolymers.²¹ One possible explanation for this fact may be that the butyl pendant chain of the ε-DL, could, to a certain extent, hinder the crystallization of ω-PDL. During the crystallization, few crystal nuclei would be formed facilitating the growth of the crystals. Thus, the ω-PDL crystals of the poly(ω-PDL-co-ε-DL) would be more stable, more perfect and consequently melt at higher temperatures. Indeed, the average crystal size measurements, determined by WAXS, support this theory. While the crystal size of the poly(ω-PDL-co-δ-HL) ranged from 12 to 19 nm (with the exception of PDL-HL 82), it was in the 20 to 23 nm range for the ε-DL based copolymers, so it can be stated that the latter are bigger.

Fig. 5 Melting enthalpies and temperatures of poly(ω-pentadecalactone-co-δ-hexalactone) and poly(ω-pentadecalactone-co-ε-decalactone) with different chain microstructures at different compositions.

Fig. 6 shows the diffraction profiles of some poly(ω-PDL-co-δ-HL) together with the diffractogram of the reference homopolymer (PPDL). As can be seen, at higher contents of δ-HL in the copolymers the reflection intensity decreases. However, the poly(ω-PDL-co-δ-HL) presented exactly the same signals as PPDL at 2θ of 21.7, 24.1, 30.1 and 36.3° (the peak at 27.3° may be due to the bismuth compounds because it also appeared at high temperatures when the polymers were fully amorphous). Hence, it was demonstrated that the ω-PDL-co-δ-HL copolymers and the PPDL present the same crystal structure, confirming the fact that only the ω-PDL units were able to crystallize.

Fig. 6 WAXS profiles of PPDL and the ω-PDL-co-δ-HL copolymers in the 10 ≤ 2θ ≤ 38 range.

3.4. Thermogravimetric analysis

Fig. 7 shows the first derivative of weight loss against temperature (DTG curves) belonging to the PPDL homopolymer and the ω-PDL-co-δ-HL copolymers. The curves of percent weight loss are also shown in the ESI.† As can be seen, four well-defined peaks (especially true in the case of PDL-HL 39) can be distinguished in the DTG curves, which were generated because of the high residual catalyst content of these polymers (a monomer/catalyst molar ratio of 100 was used in their synthesis). With increased metal catalyst content there are more units bonded to the metal atoms so metal-accelerated depolymerization is a more significant phenomenon in the pyrolysis of the polyesters.⁶³ As a result of this effect, the peaks of thermal degradation of δ-HL and ω-PDL split into two stages: the decomposition of the bismuth accelerated sequences and the degradation of the catalyst free chains.

Fig. 7 Differential thermogravimetric (DTG) curves of PPDL and the ω-PDL-co-δ-HL copolymers.

The first two peaks of thermal degradation at around 275 °C and 350 °C were attributed to the decomposition of δ-HL-rich sequences. These two peaks did not appear in the curve of the PPDL and were more pronounced in the case of PDL-HL 39, the copolymer with the highest δ-HL content. As the ω-PDL content rose, the intensity of these peaks decreased and they eventually vanished in PPDL. Conversely, two stages of degradation of ω-PDL were also found in the DTG curve of PPDL that also appeared in the poly(ω-PDL-co-δ-HL) curves. The first peak, the more intense, appeared at ∼400 °C whereas the thermal degradation of the ω-PDL sequences that are less exposed to the residual metal was delayed to around 440 °C. These two peaks of degradation of the ω-PDL blocks were more stable under thermal degradation than PCL (which degrades in the range of 350 to 400 °C), owing to the lower proportion of ester groups in the macrolactone.

On the other hand, it is also worth mentioning that these kind of copolymers show a more random distribution of sequences than the poly(ω-pentadecalactone–ε-decalactone) studied previously by our group (0.38 < R < 0.49).⁵⁰ This is reflected in the DTG curves of the ω-PDL-co-δ-HL copolymers with a higher content of ω-PDL (from 72 to 82%). Due to their random tendency (0.82 < R < 0.85) the peaks assigned to the thermal degradation of δ-HL blocks were difficult to distinguish and appear almost overlapped with the peak at 400 °C.

3.5. Hydrolytic degradation study

DSC scans were made from 21 to 140 °C: at room temperature (before submerging the samples in PBS), after storage for 24 hours at 37 °C and at different degradation times (after drying). The data relating to melting enthalpies and temperatures at the start, on day 98 and at the end of the study for all the polymer films are gathered in Table 5.

Table 5 Melting enthalpies and temperatures of all the polymer films at the start and at the end of the degradation study

Sample	21 °C		37 °C		On day 98 of degradation at 37 °C		On day 182 of degradation at 37 °C
	ΔH (J g^−1)	Tm (°C)	ΔH (J g^−1)	Tm (°C)	ΔH (J g^−1)	Tm (°C)	ΔH (J g^−1)	Tm (°C)
PPDL	133.3	102.2	133.8	102.8	149.2	102.9	147.4	106.1
PDL-HL 82	106.8	91.0	111.1	88.9	118.4	88.1	115.5	90.2
PDL-HL 76	102.1	88.2	105.9	84.9	109.6	87.5	111.9	87.9
PDL-HL 72	95.2	79.6	96.5	79.8	103.7	79.8	101.2	84.4
PDL-HL 63	77.1	74.4	80.7	76.2	84.3	79.1	90.4	80.2
PDL-HL 42	58.4	69.2	60.6	65.3	66.6	63.4	65.2	72.3
PDL-HL 39	48.3	55.4	44.0	59.0	49.7	58.7	52.7	61.4

---

The thermal properties of the poly(ω-PDL-co-δ-HL) films at 21 °C and at 37 °C can be considered as very similar (the small differences may be due to the precision of the measurements). However, it can be observed that on days 98 and 182 of degradation, the samples were slightly more crystalline with higher ΔH_m values. Moreover, the T_m values shifted to higher temperatures (2–7 °C higher), a tendency that is observed for all the materials owing to rearrangements of the crystalline phase. This is particularly true in the case of the copolymers with lower contents of ω-PDL. However, as a consequence of the highly crystalline domains of all the polymers (>44 J g⁻¹), water absorption was negligible and did not exceed a value of 2.5% in any case. Consequently, the polymers did not lose mass and the weight of the samples remained almost constant during the entire study (see Fig. 2 of ESI†).

Fig. 8 shows the progress of ln M_w against degradation time of the PPDL homopolymer and the δ-HL copolymers. As the degradation evolves, the weight average molecular weight (M_w) of the samples decreased while the dispersity rose, indicating a broader distribution of polymer chain length. As can be seen, the ω-PDL-co-δ-HL copolymers curves displayed similar slight drops, while the M_w of PPDL homopolymer remained practically constant during the 182 days.

Fig. 8 ln M_w against degradation time of PPDL and its δ-HL copolymers.

Table 6 presents the degradation rate values (K_Mw) and half degradation times (t_1/2) calculated from the slope of the fitting curve. The M_w experimental data adapts to the fitting curve (R² > 0.98) very well over the 182 days and the K_Mw estimated for the copolymers were in the 0.0013–0.0019 d⁻¹ range with corresponding t_1/2 of 365 to 533 days. All the values of degradation rate were above the obtained for PCL (0.0010 d⁻¹) in a previous study by this group³⁴ and also higher than those of the poly(ω-PDL-co-ε-DL),⁵⁰ which are almost non-biodegradable polymers.

Table 6 K_Mw, calculated on the final day of the study (day 182), and t_1/2 of the poly(ω-PDL-co-δ-HL)

Sample	KMw (d^−1)	Half-molecular weight degradation time t1/2
PPDL	<0.0002	>9.5 years
PDL-HL 82	0.0017	408 days
PDL-HL 76	0.0014	495 days
PDL-HL 72	0.0013	533 days
PDL-HL 63	0.0013	533 days
PDL-HL 42	0.0015	462 days
PDL-HL 39	0.0019	365 days

---

With the incorporation of δ-HL units, the crystal structure of ω-PDL (54% crystalline fraction) was disrupted and the copolymers became significantly more amorphous (21–44% crystalline fraction). Likewise, the proportion of ester bonds increased, which should help to promote the hydrolysis process. However, despite the fact that PDL-HL 39 (the copolymer with the highest δ-HL content) which showed the highest degradation rate (0.0019 d⁻¹), the differences between the respective K_Mw of the poly(ω-PDL-co-δ-HL) were small. This may be due to the more random chain microstructure of the copolymers of increased ω-PDL content. On the other hand, it is important to mention that no morphological changes were appreciated on the PPDL and poly(ω-PDL-co-δ-HL) samples, and the M_w of PDL-HL 82, PDL-HL 76, PDL-HL 72, PDL-HL 63, PDL-HL 42 and PDL-HL 39 decreased slightly to final values of 193, 145, 124, 82, 65 and 43 kDa. In the latter case, the film of this material became brittle after day 126 but it was studied on day 182 by polarized light optical microscopy and no sign of deterioration was found with respect to the initial sample.

3.6. Mechanical properties

Table 7 summarizes the mechanical testing results at 21° and at 37 °C (at the start and on day 182 of degradation) for the poly(ω-PDL-co-δ-HL) films, whose DSC data, before performing the mechanical tests, are already shown in Table 5. As can be seen in Fig. 9, which displays typical stress–strain curves of the polymers studied, the ω-PDL-co-δ-HL copolymers showed a marked improvement in flexibility (with lower secant modulus and increased strain recovery rates) compared to PPDL and poly(ε-caprolactone) homopolymers (secant modulus from 300 to 400 MPa, ultimate tensile strength of 28–32 MPa and deformation at break values > 750%).^34,50 Nevertheless, the tensile samples of the copolymers having a ω-PDL content below 63% (with a crystalline fraction under 36% and T_ms < 74 °C) broke early at strains lower than 100% and their performance was greatly influenced by their crystalline phase. Owing to their rather random distribution of sequences (0.74 < R < 0.85), the ω-PDL average sequence lengths of PDL-HL 63, PDL HL 42 and PDL-HL 39 were only in the range of 2.22–3.58 and that was not enough to induce a process of crystallization by orientation during the tensile test, that could increase the ductility. Conversely, in the case of the poly(ω-PDL-co-ε-DL) with ω-PDL content of 59 and 46%, the alignment of the polymer chains allowed elongation at breaks of 968 and 193% to be reached, respectively. These two copolymers, with a more perfect crystalline structure (with higher average crystal sizes and T_ms), were blockier (R ∼ 0.38) and had larger ω-PDL-blocks, l_PDL of 6.26 and 4.88, in their chain microstructure.⁵⁰

Table 7 Mechanical properties of the PPDL and the ω-PDL-co-δ-HL copolymers

Sample	Testing temperature^a	Secant modulus at 2% (MPa)	Yield strength (yield point) or offset yield strength at 10%^b (MPa)	Tensile strength at 300%^c (MPa)	Ultimate tensile strength^d (MPa)	Elongation at break (%)	Strain recovery after break^e (%)
a At 37 °C, the tests were stopped at 300% of strain due to the size limitations of the temperature chamber.b Offset yield strength at 10% was calculated for PDL-HL 42 and PDL-HL 39.c Some specimens broke before 300% of strain.d The tensile strength was determined as ultimate stress value (σu).e Measured 24 hours after break.
PDL-HL 82	21 °C	179.0 ± 13.5	7.9 ± 0.4 (18.2%)	8.8 ± 0.2	20.2 ± 2.2	1145 ± 100	43.3 ± 6.0
	37 °C	154.3 ± 2.8	6.9 ± 0.2 (28.4%)	7.6 ± 0.1	—	>300	—
	On day 182 (37 °C)	185.3 ± 9.2	7.5 ± 0.5 (14.6%)	7.4 ± 0.6	—	>300	—
PDL-HL 76	21 °C	155.8 ± 6.3	7.0 ± 0.5 (18.6%)	7.3 ± 0.2	12.6 ± 1.3	961 ± 80	37.7 ± 4.7
	37 °C	118.3 ± 4.4	5.4 ± 0.2 (22.1%)	6.2 ± 0.3	—	>300	—
	On day 182 (37 °C)	161.3 ± 1.6	6.4 ± 0.2 (13.1%)	6.8 ± 0.2	—	>300	—
PDL-HL 72	21 °C	115.8 ± 2.3	6.3 ± 0.4 (36.8%)	7.2 ± 0.2	9.9 ± 1.1	765 ± 87	41.1 ± 2.7
	37 °C	81.2 ± 3.1	5.1 ± 0.2 (46.5%)	5.6 ± 0.1	—	>300	—
	On day 182 (37 °C)	111.2 ± 4.1	5.4 ± 0.4 (25.7%)	—	4.8 ± 0.4	102 ± 13	—
PDL-HL 63	21 °C	82.9 ± 1.4	4.7 ± 0.2 (38.9%)	—	4.4 ± 0.3	82 ± 10	80.0 ± 5.9
	37 °C	58.9 ± 1.3	3.1 ± 0.1	—	3.3 ± 0.1	25 ± 7	—
	On day 182 (37 °C)	80.0 ± 4.8	—	—	3.7 ± 0.2	13 ± 2	—
PDL-HL 42	21 °C	50.9 ± 4.7	3.1 ± 0.1	—	4.0 ± 0.1	64 ± 8	92.2 ± 1.9
	37 °C	34.7 ± 2.3	2.1 ± 0.1	—	2.4 ± 0.2	25 ± 5	—
	On day 182 (37 °C)	38.4 ± 2.5	—	—	1.9 ± 0.1	8.3 ± 0.7	—
PDL-HL 39	21 °C	25.7 ± 1.1	1.8 ± 0.1	—	2.0 ± 0.2	28 ± 6	95.1 ± 1.3
	37 °C	15.8 ± 0.3	—	—	0.81 ± 0.06	8.7 ± 0.5	—
	On day 182 (37 °C)	14.4 ± 1.9	—	—	0.31 ± 0.08	2.3 ± 0.3	—

---

Fig. 9 Tensile stress–strain curves of PPDL and the ω-PDL-co-δ-HL copolymers at 21 °C.

Despite the low strain at break of the three poly(ω-PDL-co-δ-HL) mentioned above, which is even more remarkable at 37 °C, they all presented strain recoveries higher than 80% with secant modulus values between 25.7 and 82.9 MPa and final strength values in the range of 2.0 and 4.4 MPa. On the contrary, the rest of the copolymers with ω-PDL contents over 72% exhibited lower strain recoveries but elongation at break values well-above 765%. The secant modulus values obtained for PDL-HL 82 and PDL-HL 72 were 179 and 115.8 MPa and presented yield strengths of 7.9 and 6.3 MPa and tensile strengths at break of 20.2 and 9.9 MPa.

Fig. 10 show the tensile stress–strain curves up to 150% of strain of the tensile tests conducted at 37 °C with non-degraded (curves in bold type) and degraded specimens. As a result of the crystallization during the hydrolytic degradation process, the stress related properties were found slightly higher than at the start of the study. As can be noted, the elastic modulus of the samples submerged for 182 days in PBS (see also Table 7) increased with respect to that of the non-degraded films, reaching secant modulus values similar to those obtained at 21 °C. In addition, the stress also took greater values at the same strain and the yield points were narrower (with higher yield strengths), and shifted to lower values. However, owing to the changes in molecular weight and the mechanical deterioration experienced in an aqueous medium, the poly(ω-PDL-co-δ-HL) with ω-PDL contents lower than 76% lost their deformation capability. For example, PDL-HL 72 exhibited elongation at break values above 300% but at the end of the degradation study broke at only 102%.

Fig. 10 Tensile stress–strain curves up to 150% of strain of PPDL and the ω-PDL-co-δ-HL copolymers at 37 °C. The curves of the non-degraded samples are in bold type.

With the exception of PDL-HL 39, the ω-PDL-co-δ-HL copolymers displayed good mechanical stability between 21 and 37 °C, although the poly(ω-PDL-co-ε-DL) (with higher l_PDL and melting temperatures) showed steadier mechanical properties.⁵⁰ Thus, at 37 °C the loss in secant modulus and strength of PDL-HL 82 was about 14% compared to its performance at room temperature, when the PCL homopolymer suffered a drop of around 12%. These changes were even more relevant (>27%) on the PCL copolymers studied previously by this group.³⁴

4. Conclusions

In this work several statistical copolymers based on ω-pentadecalactone and δ-hexalactone were synthesized with triphenyl bismuth as catalyst. The poly(ω-PDL-co-δ-HL)s showed a slightly deviated from random distribution of sequences (0.74 < R < 0.85) and were carefully characterized by NMR, GPC, DSC, WAXS and TGA measurements. In addition, their mechanical properties were tested at room temperature (21 °C) and at body temperature (37 °C), whereas an in vitro degradation study was also carried out at 37 °C for 182 days in phosphate buffer solution. Owing to their lower stiffness and upgraded biodegradability in comparison to poly(ε-caprolactone) there are high hopes for these new thermoplastic elastomers, which also present high thermal stability and rapid crystallization from melt. However, significant improvements are necessary in the methodology of synthesis in order to shorten the reaction times and make the process better adapted to industry.

Acknowledgements

The authors are thankful for funds from the Basque Government, Department of Education, Universities and Research (GIC12/161-IT-632-13) and the Spanish Ministry of Innovation and Competitiveness MINECO (MAT2013-45559-P). J.F. thanks the University of the Basque Country (UPV-EHU) for a pre-doctoral grant.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c5ra23404b

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