Biodegradable elastomeric scaffolds for soft tissue engineering

Journal of Controlled Release 87 (2003) 69–79 www.elsevier.com / locate / jconrel

Biodegradable elastomeric scaffolds for soft tissue engineering ˆ Ana Paula Pego, Andre´ A. Poot, Dirk W. Grijpma, Jan Feijen* Institute for Biomedical Technology ( BMTI) and Department of Polymer Chemistry and Biomaterials, Faculty of Chemical Technology, University of Twente, P.O. Box 217, 7500 AE Enschede, The Netherlands

Abstract Elastomeric copolymers of 1,3-trimethylene carbonate (TMC) and ´-caprolactone (CL) and copolymers of TMC and (DLLA) have been evaluated as candidate materials for the preparation of biodegradable scaffolds for soft tissue engineering. TMC-DLLA copolymers are amorphous and degrade more rapidly in phosphate-buffered saline (PBS) of pH 7.4 at 37 8C than (semi-crystalline) TMC-CL copolymers. TMC-DLLA with 20 or 50 mol% TMC loose their tensile strength in less than 5 months and are totally resorbed in 11 months. In PBS, TMC-CL copolymers retain suitable mechanical properties for more than a year. Cell seeding studies show that rat cardiomyocytes and human Schwann cells attach and proliferate well on the TMC-based copolymers. TMC-DLLA copolymers with either 20 or 50 mol% of TMC are totally amorphous and very flexible, making them excellent polymers for the preparation of porous scaffolds for heart tissue engineering. Porous structures of TMC-DLLA copolymers were prepared by compression molding and particulate leaching techniques. TMC-CL (co)polymers were processed into porous two-ply tubes by means of salt leaching (inner layer) and fiber winding (outer layer) techniques. These grafts, seeded with Schwann cells, will be used as nerve guides for the bridging of large peripheral nerve defects.  2002 Elsevier Science B.V. All rights reserved. D,L-lactide

Keywords: Poly(ester carbonate)s; Soft tissue engineering; In vitro hydrolysis; Porous structures

1. Introduction An essential stage in tissue engineering is the design and fabrication of porous three-dimensional scaffolds. In general the scaffold should be fabricated from a biocompatible material that degrades and resorbs at a controlled rate to match cell and tissue growth in vitro and / or in vivo. The material’s surface should sustain cell adhesion and proliferation. Finally, the material should be processable into porous three-dimensional structures with intercon*Corresponding author. Tel.: 131-53-4892-968; fax: 131-534893-823. E-mail address: [email protected] (J. Feijen).

nected pores in a variety of shapes and sizes with mechanical properties after cell seeding and proliferation that match those of the tissues at the site of implantation. In the past few years much attention has been given to natural materials like collagen [1,2] and synthetic materials like poly(lactide) (poly(LA)) and poly(glycolide) (poly(GA)) and their copolymers [3– 6] as candidate materials for the preparation of porous scaffolds. However, these materials have limitations for use in the field of soft tissue engineering. Collagen presents some disadvantages associated with polymers of natural source such as large batch-to-batch variations upon isolation from biological tissues, as well as restricted versatility in

0168-3659 / 02 / $ – see front matter  2002 Elsevier Science B.V. All rights reserved. doi:10.1016/S0168-3659(02)00351-6

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designing devices with specific mechanical properties. Poly(LA) and poly(GA) are rather stiff and brittle materials that can have too high degradation rates for certain applications [7]. For soft tissue engineering applications, synthetic elastomeric materials with tunable degradation properties would be preferable. The aim of this work is to prepare biodegradable, elastomeric porous structures for scaffolds in soft tissue engineering. Poly(1,3-trimethylene carbonate) (poly(TMC))—a rubbery and amorphous polymer— was taken as a starting point in the design of alternative synthetic materials. In order to obtain materials with suitable mechanical properties and degradation rates TMC was copolymerized with either D,L-lactide (DLLA) or ´-caprolactone (CL). We have previously reported on the physical and mechanical properties of poly(ester carbonate)s based on TMC [8,9]. In the present study the effect of copolymer composition on the degradation in phosphate-buffered saline (PBS) (pH 7.4, 37 8C) has been evaluated. Furthermore, TMC (co)polymers have been used for culturing rat cardiomyocytes and human Schwann cells. Based on their mechanical properties, degradation rates and adhesion and proliferation of appropriate cells, (co)polymers have been selected and processed into porous scaffolds, which will be investigated later for their suitability for heart tissue engineering and peripheral nerve guides.

2. Materials and methods

2.1. Materials Polymer grade TMC was obtained from Boehringer Ingelheim, Germany. CL (Acros Organics, Belgium) was purified by drying over CaH 2 and distillation under reduced argon pressure. Polymer grade DLLA (Purac Biochem, The Netherlands) was used without further purification. Stannous octoate (SnOct 2 ) was used as received from Sigma, USA. Powdered sugar (Suiker Unie, The Netherlands) and sodium chloride (Merck, Germany) used as porosifying agents in the preparation of porous structures were fractionated by means of sieving, using standard test sieves (Endecotts, UK). Solvents were of analytical grade.

2.2. Polymer synthesis The synthesis and characterization of the (co)polymers have been described previously [8,9]. Briefly, polymerizations were conducted by ringopening polymerization in evacuated and sealed glass ampoules using stannous octoate as a catalyst. All homo- and copolymerizations were carried out for a period of 3 days at 13062 8C. The obtained polymers were purified by dissolution in chloroform and precipitation into a 10-fold volume of isopropanol. The precipitated polymers were collected, washed with fresh isopropanol and dried under reduced pressure at room temperature until constant weight.

2.3. Hydrolytic degradation For hydrolysis tests, rectangular films (10035 mm 2 ) of the different polymers were melt-pressed at 140 8C for 4 min to a thickness of 600 mm and rapidly cooled in the mold to approximately 15 8C with cooling water. The specimens were incubated at 37 8C in PBS (pH 7.4, NPBI, The Netherlands) containing 0.02 wt / vol% of NaN 3 (Sigma, USA) to prevent bacterial growth. PBS was changed once a month. At selected time points two samples were removed from the PBS and, after blotting, weighed and their dimensions monitored. The mechanical properties of the wet samples were evaluated at room temperature. After drying under vacuum at room temperature samples were weighed and subsequently processed for evaluation of composition, molecular weights and thermal properties. Water absorption was defined as follows: water absorption (wt%)51003(w w 2w r ) /w r , where w w represents the weight of the wet sample after blotting and w r represents the remaining weight of the sample after drying. Mass loss was defined as: mass loss (wt%)51003(w i 2w r ) /w i , where w i represents the initial dry sample weight.

2.4. Analysis and characterization Polymers were characterized with respect to residual monomer content and chemical composition by nuclear magnetic resonance (NMR). Three hundred MHz 1 H NMR (Varian Inova 300 MHz) spectra were recorded using polymer solutions in CDCl 3 (Sigma, USA).

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Molecular weights and molecular weight distributions were determined by gel permeation chromatography (GPC) using a Waters Model 510 pump, a HP Ti-Series 1050 autosampler, a Waters Model 410 Differential Refractomer and a Viscotek H502 Vis˚ Waters cometer Detector with 10 5 –10 4 –10 3 –500 A Ultra-Styragel columns placed in series. Chloroform was used as eluent at a flow rate of 1.5 ml / min. Narrow polystyrene standards were used for calibration. Sample concentrations of 0.4–0.5 wt / vol% and injection volumes of 30 ml were used. All determinations were performed at 25 8C. The thermal properties of the polymer samples were evaluated by differential scanning calorimetry (DSC). Samples (5–15 mg) placed in aluminum pans were analyzed with a Perkin-Elmer Pyris1 at a heating rate of 10 8C / min. Although this heating rate might result in a temperature overshoot, it is often applied in the thermal characterization of degradable polymers. All samples were heated from 40 8C below their glass transition temperature (T g ) to 40 8C above their T g or melting temperature (when present). The samples were then quenched rapidly (300 8C / min) until 40 8C below their T g and after 5 min a second scan was recorded. Unless mentioned otherwise, the data presented were collected during the second heating scan, as this provides information on the structure–properties relationships of the polymers independent of thermal history. The glass transition temperature was taken as the midpoint of the heat capacity change and the peak melting temperature (T m ) was determined from the melting endotherm. Tensile tests were carried out at room temperature using a Zwick Z020 universal tensile testing machine according to the ASTM standard D882-91. The machine was operated at a crosshead speed of 50 mm / min without extensometers, and an initial gripto-grip separation of 50 mm. The determined Young’s modulus (calculated from the initial slope of the stress–strain curves) gives only an indication of the stiffness of the different polymers.

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of both polymers were melt-pressed at 140 8C for 9 min to a thickness of 200 mm and rapidly cooled to 15 8C in the mold with cooling water. Subsequently, disks with a diameter of 8 mm were punched out of the films and sterilized by two incubation steps in 70 vol% ethanol solution for 15 min, followed by a rinsing step of 30 min in sterile water. Although this method is not a standardized and accepted sterilization method in surgery, this procedure did not lead to infections during our seeding experiments. Cardiomyocytes were seeded on the copolymer surface (5310 4 cells / cm 2 ) and the adhesion and proliferation were studied by phase contrast microscopy. The culture of human Schwann cells on surfaces of TMC-CL-based (co)polymers is described in a separate study [11]. Briefly, primary human Schwann cell cultures were established from pieces of sural nerve that remained after nerve transplantation to restore brachial plexus lesions. The pieces of nerve were stripped and stored in Schwann cell culture medium in a CO 2 incubator. After 4–6 days 1-mm 3 nerve pieces were explanted into gelatin coated 75cm 2 culture flasks. After 1 week these pieces were transplanted into fresh culture flasks; this procedure was repeated three times. Immunostaining confirmed that 95–98% of the cells emerging from the tripletransplanted pieces of sural nerve were Schwann cells. Films were prepared by casting polymer solutions (4.0–6.5 wt / vol%) in chloroform onto glass plates. After drying the films under reduced pressure at room temperature, disks with a diameter of 17 mm were punched out. The disks were sterilized by two incubation steps in 70 vol% ethanol solution for 15 min, followed by a rinsing step of 30 min in sterile water. Schwann cells were seeded (8310 2 cells / cm 2 ) and cultured on poly(TMC), poly(TMC-CL) either with 10 or 82 mol% of TMC and poly(CL) disks. At selected time points, cultures were gently rinsed with PBS, fixed with Cryofix  and stained with 0.25 wt / vol% Coommassie blue solution.

2.5. Cell culture Rat cardiomyocytes (embryonal rat cardiomyocyte cell line; ATTC: CRL-1446) were cultured using previously described conditions [10] on compression molded TMC-DLLA copolymer films. These copolymers contained 20 or 50 mol% of TMC. Films

2.6. Preparation of porous scaffolds and nerve guides Cylindrical porous scaffolds (5514 mm, height58 mm) based on TMC-DLLA copolymers with either 20 or 50 mol% of TMC were prepared by compres-

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sion molding in combination with NaCl porosifying particles. Mixtures of different polymer / salt ratios and different salt size ranges were melt-pressed at 140 8C for 10 min and rapidly cooled in the mold to 15 8C with cooling water. The salt particles were leached out in water for 5 days (with daily renewal of the water) and the porous samples were subsequently dried. Two-ply porous nerve guides were prepared by a combination of dip-coating (inner layer) and fiber winding (outer layer) techniques. To prepare the inner layer a 3 wt / vol% solution of poly(TMC) in chloroform was prepared and sugar crystals (,20 mm) were added in a 15 / 85 (w / w) polymer / salt ratio. A glass mandrel (551.5 mm) was immersed in this suspension, subsequently pulled out and the chloroform was allowed to evaporate at room temperature (18–22 8C) overnight. The outer layer fibers were spun from 10 wt / vol% poly(TMC-CL) (10:90 mol%) solutions in chloroform containing 0.17 wt / ] vol% poly(ethylene oxide) (PEO, Mw 58310 6 g / mol, Aldrich, Germany) and wound on the rotating previously dip-coated tube. The sugar was leached out for 2 weeks in water (the water was renewed several times) and the remaining structure was dried at room temperature under reduced pressure. The structure of the prepared porous scaffolds was evaluated using a Hitachi S800 field emission scanning electron microscope (SEM) operating at 5 kV. Prior to microscopical evaluation the structures were coated with a gold layer using a Polaron E5600 sputter-coater. SEM showed that the porous scaffolds did not contain residual salt or sugar.

3. Results and discussion

3.1. Polymer synthesis and characterization The synthesis of the statistical copolymers poly(TMC-CL) and poly(TMC-DLLA) was performed by ring-opening polymerization in the melt, using stannous octoate as a catalyst (Fig. 1). The homopolymers and a series of copolymers were prepared. The residual monomer content and the copolymer composition were determined by 1 H NMR analysis of the crude polymerization products. The monomer conversion was high (.94%) or almost complete and the obtained compositions were in accordance with the ratio of monomers charged. Under the applied polymerization conditions high ] molecular weight polymers (Mn .100 000 g / mol) were obtained. This is important in order to obtain materials with good mechanical performance even after heat processing. The starting molecular weight influences the resorption time of the implant. Independent of the degradation mechanism, the onset of the loss of mechanical properties and mass is delayed at higher initial molecular weights [12]. The polydispersity values of all polymers ranged from 1.6 to 2.3. The thermal properties (Fig. 2), and consequently the mechanical properties, of the TMC copolymers can be varied to a large extent by adjusting the comonomer composition. TMC and DLLA (co)polymers are amorphous with a T g varying between 217 8C for poly(TMC) and 53 8C for poly(DLLA). Therefore, the polymers obtained vary from rubbers

Fig. 1. Synthesis of statistical poly(trimethylene carbonate-co-´-caprolactone) and poly(trimethylene carbonate-co-D,L-lactide) copolymers.

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of TMC-CL- and TMC-DLLA-based copolymers previously described in literature [13,14]. The available data do not cover the whole range of copolymer compositions and do not provide information on the mechanical properties of TMC-based copolymers during degradation. We investigated the degradation behavior of solid compression molded samples of selected TMCDLLA and TMC-CL (co)polymers in PBS at 37 8C. ] In Table 1 the composition, molecular weights (Mw ] and Mn ), Young’s modulus and maximum strength (smax ) of the materials under study are presented. Fig. 2. Thermal properties of TMC-CL and TMC-DLLA copolymers after purification as a function of TMC content. (j) T g of TMC-DLLA copolymers; (♦) T g and (d) DH of TMC-CL copolymers.

to stiff materials as the content of TMC decreases. CL-based copolymers are flexible materials as their T g values are all below room temperature. At CL contents above 70 mol% all copolymers are semicrystalline. Higher CL contents result in increased melting temperatures (T m ) and heats of fusion (DH ). As a consequence, an improvement in the toughness of the polymers was observed. Poly(TMC) showed unexpected high values for the tensile strength as it presents strain-induced crystallization [8].

3.2. In vitro degradation of TMC-DLLA and TMCCL copolymers There are few studies on the degradation behavior

3.2.1. TMC-DLLA copolymers TMC-DLLA (co)polymers are hydrophobic [9] and the initial water uptake of the DLLA-based copolymers was below 2 wt%. The water uptake did result in a decrease of the T g of the polymers. First DSC scans show the thermal properties of the (co)polymers equilibrated in water at 37 8C (see Table 2). The increased chain mobility due to water uptake and the conditioning temperature (37 8C) changed the mechanical behavior of the samples and their dimensions. The elastic modulus of all materials, except that of poly(TMC), decreased after water uptake (see Table 2). In particular, the behavior of the copolymer containing 20 mol% of TMC is worth noticing. At room temperature this material is in the glassy state allowing easy processing and handling. After water uptake, T g decreases below the conditioning temperature of 37 8C and the material now

Table 1 Characterization of TMC-DLLA and TMC-CL compression molded samples used in the in vitro degradation study ] ] Polymer TMC Mw 310 25 Mn 310 25 Young’s (mol%) (g / mol) (g / mol) modulus (MPa) Poly(TMC) Poly(TMC-DLLA) Poly(DLLA) Poly(TMC-CL) Poly(CL)

100 79 50 20 0 82 31 10 0

5.8 3.4 5.7 6.1 7.3 4.8 3.8 2.7 3.4

3.2 1.7 2.6 2.9 2.9 2.6 2.0 1.5 1.8

6 5 16 1900 1900 5 3 140 320

smax (MPa) 12 2 10 51 53 2 0.4 23 32

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Table 2 Physical properties of TMC-DLLA compression molded samples at equilibrium water uptake a TMC content (mol%)

Water uptake (%)

Tg b (8C)

Young’s modulus (MPa)

smax (MPa)

100 79 50 20 0

1.4 1.6 1.6 1.2 1.1

219 29 11 33 46

6 4 13 1100 1400

18 1 11 38 50

a

Equilibrium reached after conditioning in PBS, at 37 8C for 24 h. b First heating scan (DSC).

shows rubbery behavior. Poly(TMC) specimens show a small reduction in length (smaller than 10%) during the degradation period. In contrast copolymers with either 79 or 50 mol% of TMC shrink and the maximum observed sample length reductions are 30 and 25%, respectively. Poly(DLLA) and poly(TMC-DLLA) containing 20 mol% of TMC do not show any dimensional change. ] Fig. 3 shows the Mn of TMC-DLLA (co)polymers as a function of degradation time. After more than a year in PBS poly(TMC) did not show any significant ] ] decrease in Mn . In contrast, the Mn of poly(DLLA) ]0 (M n 5291 000 g / mol) decreased continuously reaching 8% of its initial value after 60 weeks. The copolymers degraded much faster than the parent

] Fig. 3. Molecular weight (Mn ) as a function of time for TMCDLLA copolymers during degradation at 37 8C and pH 7.4. (d) 100% TMC; (.) 79% TMC; (♦) 50% TMC; (m) 20% TMC and (j) 100% DLLA.

homopolymers. This can be explained by two factors. Ester bonds are more labile to hydrolysis than carbonate bonds as shown by NMR analysis. In time the mol% of TMC in the copolymers tended to increase, indicating that chain cleavage is not random but preferential for ester sequences. Also, T g reduction caused by water uptake increases chain mobility. Especially when T g is reduced to values below the conditioning temperature of 37 8C (see Table 2), the access of water to the labile bonds is facilitated, increasing the hydrolysis rates. At later stages of degradation an increase in water absorption was observed possibly due to an increase of end groups. The mechanical properties of the non-degraded wet polymer specimens vary with composition, ranging from brittle materials in the case of the glassy polymers to highly elastic rubbers for the low T g polymers with high contents of TMC (Table 2). In time, a decrease of the mechanical properties of the copolymers was observed as a result of decreasing molecular weights. After 3 months the copolymer with 50 mol% of TMC was fragmented with only limited mass loss. The shape of the copolymer with 20 mol% of TMC was still intact after 4 months, although it broke brittle when handled. The copolymer with 79% TMC lost its tensile strength after 5 months of degradation. For poly(DLLA) a steady decrease of the mechanical properties was observed during the period of study, which can be related to the decrease in molecular weight. Samples lost their mechanical strength after 60 weeks in PBS. In the case of poly(TMC) only a relatively small deterioration of the tensile strength (53%) was observed during this 60-week study. For all polymers no weight loss was observed during the initial stages of degradation. However, after loss of tensile strength a sharp increase in mass loss was observed which coincides with a narrowing of the molecular weight distribution of the samples [15]. Copolymers with either 50 or 80 mol% of TMC in particular, undergo total resorption in less than 11 months.

3.2.2. TMC-CL copolymers The degradation of the TMC-CL-based copolymers was much slower than that of the TMC-DLLAbased copolymers. Initially, poly(CL) and poly-

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(TMC-CL) with 10 mol% of TMC were white and opaque, while the other CL-based copolymers and poly(TMC) were transparent. This is in agreement with the crystalline structure of the former polymers and the amorphous structure of the latter (Fig. 2). When placed in PBS the water absorption was low in accordance to previous reports [8]. All transparent amorphous polymers became translucent from day 1 and, except for poly(TMC), the dimensions of the samples changed in time—the samples became more undulated, twisted and shrunk (specimens of the copolymers with 82 and 31% TMC decreased in length up to 20 and 70%, respectively). This can be explained by a relaxation effect due to an increase of chain mobility at 37 8C as compared to room temperature in combination with a plasticizing effect of water. The semi-crystalline copolymers showed no changes in their shape or dimensions during the period of study. ] Fig. 4 shows Mn as a function of degradation time for the TMC-CL copolymers. All polymers exhibit a much lower hydrolysis rate than TMC-DLLA copolymers. CL-based copolymers degrade faster than poly(TMC). In this case a plasticizing effect of the absorbed water has no significant effect on the rate of hydrolysis because all polymers have a T g below 217 8C (T g of the TMC homopolymer). In accord] ance with the small decrease in Mn during the time

] Fig. 4. Molecular weight (Mn ) as a function of time for TMC-CL copolymers during degradation at 37 8C and pH 7.4. (d) 100% TMC; (.) 82% TMC; (앳) 31% TMC; (n) 10% TMC and (j) 100% CL.

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of the study, no significant changes of water uptake, molecular weight distribution or mass were observed for the TMC-CL copolymers [15]. The mechanical performance of the poly(TMCCL) copolymers in the wet state was not significantly different from the performance of the dry samples [8]. The mechanical properties of these polymers as a function of degradation time followed a similar trend as the molecular weight, with all polymers possessing suitable mechanical performance even after 1 year of degradation. The copolymer with 31 mol% of TMC showed poor mechanical properties from the start. The first DSC heating scan can provide information on the thermal history of the sample, the plasticizing effect of solvents (in the present case water) as well as the extent of degradation of the sample. In the case of the semi-crystalline TMC-CL copolymers some changes in the thermal properties of the polymers occurred during the incubation time. During degradation of the samples in PBS at 37 8C it was observed that the T m and the heat of fusion increased. This increase in crystallinity due to annealing was much higher for poly(CL) than for the CL copolymer with 10 mol% TMC. For poly(CL) T m and DH increased linearly from initial values of, respectively, 60 8C and 67 J / g to 70 8C and 82 J / g in 60 weeks. The initial crystallinity of the copolymer containing 10 mol% of TMC was lower (T m 542 8C and DH547 J / g) only increasing to 51 8C (T m ) and 53 J / g (DH ). Concerning the use of these polymers in the preparation of porous scaffolds for tissue engineering, one should consider the fact that crystalline debris formed during degradation may cause an undesired late inflammatory response negatively influencing the tissue growth and normal function [16]. Therefore, a minimal degree of crystallinity is desired. In summary, materials with different degradation profiles can be readily obtained by changing the type of comonomer and the copolymer composition. Very slowly degrading polymers, that can maintain good mechanical properties for more than 1 year are poly(TMC) and the TMC-CL copolymer with 10 mol% of TMC. Materials that undergo complete resorption in less than a year are copolymers of TMC and DLLA containing either 20 or 50 mol% of TMC.

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3.3. TMC copolymers as scaffolds for soft tissue engineering TMC copolymers have suitable physical, mechanical and degradation characteristics for the preparation of biodegradable, elastomeric porous scaffolds for soft tissue engineering. In the next section we will discuss the design of optimal scaffolds for heart tissue engineering and nerve regeneration, respectively.

3.3.1. Heart tissue engineering Replacing damaged heart tissue with newly grown material from autologous origin can be a very promising treatment to prevent post-ischemic heart failure complications. In the tissue engineering approach to cardiomyoplasty, cardiomyocytes are incorporated into a scaffold that is implanted in or onto the scar tissue. After the cells have organized, proliferated and excreted their own extracellular matrix, the scaffold should degrade and resorb in the body. The material strength needs to be retained for at least the period during which a cardiac tissue structure with sufficient strength has formed. In the rat model, this is estimated to take between 1 and 3 months [17]. After this period, the polymeric scaffold should resorb with minimal tissue response. In order to allow the contraction of the growing tissue and to withstand the contractions of the surrounding myocardium after implantation, the scaffold should be very flexible and able to sustain cyclic loading. TMC-DLLA copolymers with 20 and 50 mol% of TMC are totally amorphous and very flexible under physiological conditions, and resorb completely within 10 months. They are therefore excellent materials for the preparation of scaffolds for heart tissue engineering. Since cardiomyocyte survival depends on their attachment to the substrate, first, the culture of rat cardiomyocytes on films of selected TMC-DLLA copolymers was investigated [18]. The obtained results indicate that rat cardiomyocytes attach and proliferate well on the copolymer surfaces (Fig. 5). The use of three-dimensional scaffolds based on copolymers of TMC and DLLA for heart tissue engineering is currently under investigation. Threedimensional porous scaffolds have been prepared by compression molding and particulate leaching tech-

Fig. 5. Phase-contrast micrographs (magnification: 350 and 3200) of rat cardiomyocytes (cell line) on the surface of a poly(TMC-DLLA) (20:80 mol%) compression molded film at 8 days of culture.

niques. We are aiming at highly porous structures (porosities .95%) with pore sizes in the range of 50–150 mm (comparable to the length of a cardiomyocyte). This can be readily achieved by adjusting the size of the salt particles used and optimizing the polymer to particle ratio (Fig. 6). Presently, the growth of cardiomyocytes in the three-dimensional scaffolds is being investigated.

3.3.2. Artificial nerve grafts When the extent of nerve damage, resulting from

Fig. 6. Poly(TMC-DLLA) (20:80 mol%) porous scaffold (93% porosity, pore size between 70 and 130 mm) prepared by compression molding in the presence of 95 wt% NaCl particles of 100–250 mm and subsequent salt leaching.

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trauma or injury, prohibits direct approximation of two nerve stumps, autologous nerve grafting is considered to be the surgical treatment of choice. Artificial nerve guides offer a promising alternative to this technique, preventing the sacrifice of a donor nerve and reducing operation time. The nerve guide should be flexible, but relatively strong and easy to handle in microsurgery. In order to regenerate large nerve defects in clinically relevant situations it is anticipated that a period up to a year is required for

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the nerve to regenerate and to mature. During this period the device should keep its form and not fragment in order to avoid the formation of scar tissue that could jeopardize the regeneration process. To enhance the rate of nerve regeneration, the tube will be seeded with Schwann cells. These cells play an important role in the regeneration process, producing extracellular matrix proteins and a range of neurotropic and neurotrophic factors essential for the growth of neurons. The attachment and prolifer-

Fig. 7. Light micrographs (magnification: 3100) of human Schwann cells SCs (primary culture) on the surface of different solvent cast TMC-CL (co)polymer films at 15 days of culture. (A) poly(CL); (B) poly(TMC-CL) (10:90 mol%); (C) poly(TMC-CL) (82:18 mol%) and (D) poly(TMC).

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ation of human Schwann cells on TMC-CL (co)polymer films was previously investigated [11]. The results indicate that human Schwann cells attach and proliferate on the various (co)polymer surfaces (see Fig. 7), especially on the surface of the (co)polymers with higher TMC content. In view of the previously mentioned graft requirements and considering the polymer processability and Schwann cell culture results, poly(TMC) and poly(TMC-CL) (10:90 mol%) seem interesting candidates for the development of such a graft. These polymers were processed into porous two-ply tubes (Fig. 8) by means of salt leaching (inner layer) and fiber winding (outer layer) techniques. The microporous poly(TMC) inner layer has a barrier function, allowing exchange of fluids but avoiding the ingrowth of fibrous tissue. The fibrous poly(TMC-CL) (10:90 mol%) outer layer provides the mechanical support and strength to the graft. These tubes can be easily handled in surgery without kinking. The performance of these grafts seeded with Schwann cells in guiding severed nerves towards recovery of nerve function is presently being evaluated in vivo (sciatic nerve, rat model). The in vivo degradation of these copolymers as well as the detailed preparation of nerve guides and heart tissue grafts from these materials will be published in forthcoming papers.

4. Conclusions Copolymerization of TMC with either DLLA or CL provides materials with different mechanical properties and degradation profiles. Copolymers of TMC-DLLA are amorphous polymers which can be resorbed completely in vitro. Copolymers with 20 or 50 mol% of TMC were completely resorbed in less than a year. (Semi-crystalline) TMC-CL copolymers degraded much slower. These materials retained suitable mechanical properties for more than 1 year. Furthermore, TMC-based polymers could be readily processed into porous scaffolds with different shapes, pore sizes and porosities. The physical, mechanical and degradation characteristics of these materials and their suitability for the culturing of different cell types, make TMC-based copolymers good candidates for use in soft tissue engineering.

Acknowledgements ˆ A.P. Pego is grateful to the PRAXIS XXI programme (Portugal) for her research grant (BD/ 13335 / 97). We would like to thank C.L.A.M. Vleggeert-Lankamp and E. Marani of the Leiden University Medical Center (artificial nerve grafts) and X.J. Gallego y van Seijen and M.J.A. van Luyn of the Groningen University (heart tissue engineering) for collaborating in this research. We are also grateful to M. Smithers for the SEM work.

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Fig. 8. Two-ply nerve guide consisting of a microporous inner layer (pore size between 1 and 10 mm) based on poly(TMC) and a macroporous outer layer (pore size between 20 and 60 mm) based on 10 mol% TMC-CL copolymer.

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