Biodegradable, Amorphous Copolyester‐Urethane Networks Having Shape‐Memory Properties

In great shape: Star-shaped telechelic oligomers can be coupled with a junction unit to yield amorphous copolyester-urethane networks, which are not only biodegradable but can also be fixed in a new, temporary shape after they have been processed into a permanent shape. The permanent shape can be recovered by heating (see picture). Such polymer networks may find application in ophthalmology. Degradable polymers having a thermally induced shape memory can be fixed in a new, temporary shape after they have been processed into a permanent shape. They have great potential for biomedical applications, especially in the area of minimally invasive surgery.1 One example is the insertion of a bulky medical device in a compressed temporary shape through a small surgical incision. When the implant is heated above a switching temperature (Ttrans), it returns to its application-relevant permanent shape. After a given time the device degrades, and a second surgery for its removal is not necessary.2, 3 Shape-memory polymers generally consist of two components: cross-links determining the permanent shape and switching segments fixing the temporary shape at temperatures below Ttrans. Cross-linkage can be achieved either by physical interaction (e.g., in thermoplastic polymers) or by chemical bonds (e.g., in thermosets or photosets). In covalently cross-linked shape-memory polymer networks a maximum weight content of switching segments is possible. In constrast, thermoplastic materials must contain a sufficient amount of hard-segment-determining blocks so that a sufficient number of physical cross-links exist at temperatures above Ttrans.4 The blocks that determine the switching segment may display Ttrans as either a melting temperature or a glass transition temperature. In biodegradable shape-memory polymers previously described as thermoplastic multiblockcopolymers3 or photoset AB polymer networks2, 5 Ttrans is the melting point of crystallizable oligo(ε-caprolactone) segments. Hydrogels with hydrophobic and crystallizable side chains as molecular switches also can show a thermo-responsive one-way shape-memory effect.6 Based on noncrystallizable switching segments, completely amorphous shape-memory polymer networks having a glass transition temperature as Ttrans can be designed. These networks are transparent, and they should show a more homogenous degradation than semicrystalline polymers since hydrolytic degradation occurs faster in amorphous regions than in crystalline ones.7–11 Both features are most relevant for certain medical applications and therapeutic methods. Transparency is required for applications like replacement of ocular tissue.12–15 The homogenous degradation behavior of amorphous aliphatic copolyesters qualifies them as preferred matrices in the area of controlled drug delivery.10 However, the completely amorphous polymer networks with thermally induced shape-memory effects described in literature, such as poly(styrene)16 and poly(acrylate) resins,17 were not originally developed for medical applications and are not hydrolytically degradable. In this paper we report on, to our knowledge, the first amorphous, biodegradable, shape-memory polymer networks. Our structural concept is based on polymer networks that are synthesized by coupling well-defined star-shaped telechelic oligomers with a low-molecular-weight junction unit. In this polymer network architecture the functionality of the cross-links is defined by the branch points of the telechelic precursor. The switching segment chains are formed by coupling the end groups of two different arms of the precursors. When the structure–property relationships are known, we will be able to adjust the macroscopic properties of the polymer networks to the specific requirements of each medical application by varying these structural parameters. The amorphous shape-memory thermosets introduced in this communication were prepared from star-shaped hydroxytelechelic co-oligoesters. The cross-linking reaction was performed by adding an isomeric mixture of 1,6-diisocyanato-2,2,4-trimethylhexane and 1,6-diisocyanato-2,4,4-trimethylhexane (7) as junction units (Scheme 1). The co-oligoester segments consist of oligo[(rac-lactide)-co-glycolide] (5, 6) synthesized by copolymerization of rac-dilactide (1, mixture of 93.8 mol % D,D- and L,L-dilactide and 6.2 mol % meso-dilactide) and diglycolide (2). The cross-link points were introduced by ring-opening polymerization of 1 and 2 with the initiators 1,1,1-tris(hydroxymethyl)ethane (3) and pentaerythrite (4), resulting in trifunctional and tetrafunctional netpoints, respectively. Poly(rac-lactide)s and poly[(rac-lactide)-co-glycolide]s containing up to 75 mol % glycolide are amorphous.18 The Ttrans values of the resulting shape-memory polymers are the glass transition temperatures Tg, which are between ca. 59 °C for poly(rac-lactide)19 and ca. 36 °C for poly(glycolide).20 The homo- and copolymers of 1 and 2 are established in biomaterial science and known to be biocompatible and biodegradable.21, 22 Synthetic route to amorphous, biodegradable copolyester-urethane networks. I: initiator, f: functionality of initiator, R: oligo[(rac-lactide)-co-glycolide] segment. DBTO=dibutyltin oxide. The obtained copolyester-urethane networks exhibited a gel content higher than 90 %. The degree of swelling of the copolyester-urethane networks increased with increasing chain length of the copolyester segment due to the lower cross-link density. The networks prepared from the three-armed oligomers 5 showed a higher degree of swelling than those from four-armed species 6 when the molecular weight of the oligomers was similar, indicating a lower cross-link density. The thermal properties of polymer networks were characterized with differential scanning calorimetry (DSC) by cooling samples from the melt at 150 to −50 °C at 10 K min−1 and reheating them to 150 °C at 10 K min−1. As expected, the optically clear materials did not show any crystallization or melting peak in DSC experiments. The Tg values (second heating run) of the copolyester-urethane networks were in the range between 48 and 66 °C and about 10–50 K above the Tg values of the respective telechelic cooligomers 5 and 6 (Figure 1). For both types of polymer networks the Tg values were found to be around 55 °C when the number average molecular weight (Mn) of the precursors was higher than 4000 g mol−1. For lower Mn values, the Tg values of the networks prepared from four-armed precursors 6 are higher than those synthesized from three-armed precursors 5 and increase significantly with decreasing Mn. This effect is in agreement with the theory of Fox and Flory,23 which predicts an increase in Tg with decreasing molecular weight of the chain segments and increasing functionality of the netpoints. However, it must be noted that the chemical composition of the polymer networks (e.g. the diurethane content) depends drastically on the molecular weight of the precursors. Intermolecular interactions like H-bonding through the diurethane units and nonideal cross-linking because of possible dangling chain ends or loop formation may also influence Tg. For biomedical applications the Tg value of the networks could potentially be adjusted around body temperature by increasing the glycolide content. Glass transition temperature of copolyester-urethane networks (DSC) from the second heating cycle as a function of the molecular weight of the telechelic precursors Mn (determined by 1H NMR spectrscopy). Copolyester-urethane networks: ○: three-armed oligomers 5, •: four-armed oligomers 6. At 70 °C, approximately 20 °C above Tg, the polymer networks are in a rubber-elastic state, and no necking occurred in the stress–strain test. Young's modulus E and elongation at break εb depend on the segment length of the covalent network: with increasing molecular weight of the precursors, in other words, increasing chain lengths of the networks, E decreases and εb increases, respectively. The results of mechanical tests at 70 °C showed that it was possible to realize deformations from the permanent to the temporary shape of up to 470 %. The value for E decreased by a factor between 60 and 530 compared to that of identical sample measured at 25 °C (Table 1). The value for εb decreased from 470 % at 70 °C to 250 % at 25 °C for the loosely cross-linked material N-P-LG(17)-10 000. The tensile stress at break σb increased by at least an order of magnitude when the measurement temperature was decreased from 70 to 25 °C. Network[b] Mn [g mol−1][c] E [MPa][d] σb [MPa][d] εb [%][d] E [MPa][e] σb [MPa][e] εb [%][e] εm [%] Rf(1) [%] Rr(1) [%] R̄f,2–5 [%] R̄r,2–5 [%] N-T-LG(17)-1000 980 330±30 25.8±1.4 90±10 2.84±0.53 1.80±0.01 115±30 n.d. n.d. n.d. n.d. n.d. N-T-LG(15)-2000 2270 520±150 20.2±4.2 170±30 1.84±0.14 2.85±0.43 250±30 n.d. n.d. n.d. n.d. n.d. N-T-LG(17)-5000 4500 490±35 29.3±4.7 165±30 0.92±0.23 3.78±1.12 470±75 50[f] 91.5 98.5 94.5±2.5 98.5±1.0 N-T-LG(17)-6000 6000 345±70 26.2±1.2 305±45 1.11±0.25 1.84±0.93 375±65 100 94.5 >99.0 94.5±0.5 >99.0±0 N-T-LG(16)-8000 7900 600±125 34.7±0.3 150±5 1.22±0.14 2.67±0.54 370±10 100 95.5 >99.0 91.0±0.0 99.0±0.5 N-T-LG(18)-12 000 11 700 390±35 24.2±4.2 150±5 0.77±0.12 0.71±0.33 875±90 100 91.8 97.3 91.7±0.1 96.9±0.4 N-P-LG(17)-1000 820 340±80 29.2±4.8 100±20 5.89±0.78 3.23±0.24 90±10 n.d. n.d. n.d. n.d. n.d. N-P-LG(18)-2000 2450 495±145 26.1±3.2 50±10 2.62±0.04 2.30±0.45 180±110 n.d. n.d. n.d. n.d. n.d. N-P-LG(15)-5000 4900 375±60 30.4±5.3 240±90 1.16±0.12 2.90±0.01 240±10 50[f] 90.5 >99.0 91.0±2.5 96.5±1.5 N-P-LG(15)-7000 7300 420±35 21.4±5.3 180±5 1.21±0.19 2.38±1.07 300±130 100 92.0 >99.0 92.5±0.0 >99±0.0 N-P-LG(16)-8000 8200 350±80 18.3±3.3 265±50 1.03±0.26 1.75±0.36 430±115 100 96.0 >99.0 97.0±2.0 98.5±1.5 N-P-LG(17)-10 000 10 000 340±60 36.2±5.9 250±210 1.86±0.35 3.51±0.21 470±5 100 96.5 92.5 95.0±0.0 90.0±1.0 The shape-memory effect of amorphous biodegradable copolyester-urethane networks is illustrated in Scheme 2. In the programming step, the amorphous polyurethane network was deformed from its permanent shape to its temporary shape at a high temperature (Thigh=70 °C), held in its deformed shape, and cooled below Ttrans to Tlow. After fixation of the deformed temporary shape was complete, the load was removed. In the recovery step the network was reheated above Ttrans to Thigh, and the original permanent shape was recovered. The macroscopic shape-memory effect of these materials is demonstrated in Figure 2. The complex transformation from the temporary shape (“SM”) to the permanent shape (“corkscrew”) took approximately 5 min at a temperature of 70 °C. The series of photographs demonstrates the macroscopic shape-memory effect for a biodegradable copolyester-urethane network. The transition from the temporary shape (“SM”) to the permanent shape (“corkscrew”) took 300 s at 70 °C. Illustration of shape-memory effect for amorphous, biodegradable copolyester-urethane networks prepared from oligo[(rac-lactide)-co-gylcolide]tetrols 6 and low-molecular-weight diisocyanate 7. —: oligo[(rac-lactide)-co-glycolide] segment, ⧫⋅⋅⋅⧫: diurethane link, •: netpoint. The shape-memory properties of the polymer networks were quantified by cyclic, thermomechanical tensile experiments.4 The strain fixity rate Rf and the strain recovery rate Rr were calculated to quantify the fixation of the temporary and the recovery of the permanent shape of the polymer networks. Rf and Rr of all prepared networks were higher than 90 % for all cycles (Table 1). The shape-memory properties varied slightly between the first and the second cycle due to segment-chain orientation and relaxation effects. After the second cycle, Rf and Rr reached almost constant values. The recovery stress was in the range of 0.8–1.0 MPa when polymer network N-P-LG(17)-10 000 was heated to Th=70 °C in its temporary shape (εm=100 %). The mechanical properties of the polymer networks in their temporary shape (εu) are important characteristics for potential applications. For this purpose, a series of samples were programmed by elongating to εm and cooling to Tlow, resulting in an elongation at the stress-free state εu. Figure 3 shows the effect of εu on the mechanical properties of the polymer network N-P-LG(17)-10 000. Applying a deformation in the same direction as that used for forming the temporary shape resulted in a significant increase of E, while εu was lower than 100 %. For higher values of εu no further increase could be observed. This behavior indicates that a certain preorientation of chain segments in the amorphous polymer networks leads to higher mechanical strength and a lower elongation at break εb.24 Young's modulus E (▪) and elongation at break εb (○) for the network N-P-LG(17)-10 000 in its temporary shape at 25 °C as a function of elongation in the stress-free state εu after fixation in cyclic thermomechanical tensile experiments. The hydrolytic degradation experiments showed three major phases in the degradation of the polymer networks: the mass was unchanged during an induction phase (1), which was followed by a phase characterized by accelerated mass loss (2), and finally by a phase showing retarded mass loss (3). The degradation curves shown in Figure 4 indicate bulk degradation. Linear poly[(rac-lactide)-co-glycolide] (25 % glycolide) microspheres with Mn≈80 000 g mol−1 degraded homogenously, while the scaffold made from a copolymer with Mn>100 000 g mol−1 showed an induction period of 48 d.10, 25 The induction period is characterized by the diffusion of water into the polymer networks and the beginning of the hydrolytic degradation of the copolyester segment chains. The diffusion of water is influenced by the hydrophilicity of polymer networks which increases with increasing diurethane unit content and the glycolide content of the copolyester (Table 1 in the Supporting Information). The beginning cleavage of ester bonds led to an increase in the swelling capability. During this period the swollen and partly degraded chain segments were still connected to multifunctional netpoints. Copolyester-urethane networks from four-armed cooligoesters 6 had an induction period of 60 to 125 d, while 45 to 65 d were characteristic for the polymer networks from the three-armed oligomers 5, which reflects their lower crosslink density (Figure 4). Hydrolytic degradation of copolyester-urethane networks in aqueous phosphate buffer solution (pH 7) at 37 °C. Relative mass loss: ratio of mass mtd at degradation time td (in days) to the initial mass mt0, for polymer networks with varying segment chain length and functionality of crosslinks. Copolyester-urethane networks: □: N-T-LG(17)-1000, ○: N-T-LG(17)-5000, ▵: N-T-LG(17)-6000, ▪: N-P-LG(17)-1000, •: N-P-LG(15)-5000, ▴: N-P-LG(17)-10 000. In summary, we have demonstrated that a series of copolyester-urethane networks are amorphous and biodegradable, and have shape-memory properties. The synthetic route to the polymer networks allows the adjustment of certain molecular parameters such as functionality of the crosslinks and the segment chain length. In this way, polymer networks with variable mechanical properties but nearly constant shape-memory properties (Ttrans, Rf and Rr) could be obtained. The copolyester-urethane networks are transparent and undergo bulk degradation. Biodegradable, transparent biomaterials with shape memory may be crucial in the development of minimally invasive surgery in ophthalmology. 5, 6: All polymerizations were carried out in bulk under nitrogen at 130 °C. A mixture of 85 wt % rac-dilactide 1, 15 wt % diglycolide 2, and the initiators (3 or 4) was heated and stirred. When the melt became optically clear, dibutyltin oxide (0.2 wt %, DBTO) was added. After 5 days the reaction mixture was cooled to room temperature. The cooligomers were dissolved in a six- to tenfold excess of CH2Cl2 and then precipitated in hexane. The precipitate was washed with hexane and dried to constant weight at 80 °C. Synthesis of the copolyester-urethane network: The telechelic cooligomer was dissolved in a tenfold excess of CH2Cl2 under nitrogen. Compound 7 (molar ratio of isocyanato to hydroxy functional groups of 1.05) was added to the solution with stirring. After 5 min the reaction mixture was poured into teflon dishes and kept under nitrogen flow for 24 h at room temperature to evaporate the solvent carefully during formation of the polymer network. To complete film crosslinking, the films were kept at 80 °C for further 4 days. The crude films were extracted with chloroform and dried to constant weight under vacuum (0.1 mbar) at 80 °C. Supporting information for this article is available on the WWW under http://www.wiley-vch.de/contents/jc_2002/2005/z461360_s.pdf or from the author. Please note: The publisher is not responsible for the content or functionality of any supporting information supplied by the authors. Any queries (other than missing content) should be directed to the corresponding author for the article.