Elastomeric Hydrolyzable Porous Scaffolds: Copolymers of Aliphatic

Jul 30, 2005 - Porous scaffolds of 1,5-dioxepan-2-one (DXO), l-lactide (LLA), and ε-caprolactone (CL) were prepared by a solvent casting, salt partic...
0 downloads 19 Views 728KB Size
Download PDF
Biomacromolecules 2005, 6, 2718-2725

2718

Elastomeric Hydrolyzable Porous Scaffolds: Copolymers of Aliphatic Polyesters and a Polyether-ester Karin Odelius, Peter Plikk, and Ann-Christine Albertsson* Department of Fibre and Polymer Technology, Royal Institute of Technology, SE-100 44, Stockholm, Sweden Received March 11, 2005; Revised Manuscript Received June 23, 2005

Porous scaffolds of 1,5-dioxepan-2-one (DXO), L-lactide (LLA), and -caprolactone (CL) were prepared by a solvent casting, salt particulate leaching technique in which the composites were detached from their mold using a novel methanol swelling procedure. By incorporating DXO segments into polymers containing LLA or CL, an increase in hydrophilicity is achieved, and incorporating soft amorphous domains in the crystalline sections enables tailoring of the mechanical properties. The porosities of the scaffolds ranged from 89.2% to 94.6%, and the pores were shown to be interconnected. The materials were synthesized by bulk copolymerization of 1,5-dioxepan-2-one (DXO), L-lactide (LLA), and -caprolactone (CL) using stannous 2-ethylhexanoate as catalyst. The copolymers formed varied in structure; poly(DXO-co-CL) is random in its arrangement, whereas poly(DXO-co-LLA) and poly(LLA-co-CL) are more blocky in their structures. Introduction Since the human body is such a complex system, a large variety of materials with different characteristics are needed in order to fulfill all conceivable applications. Today, homopolymers of resorbable aliphatic polyesters including poly(L-lactide) (PLLA) and poly(-caprolactone) (PCL) are most frequently used in biomedical applications, but the need for new materials with other properties is still large. One way of creating such materials is by copolymerization, which enables the material’s degradation and mechanical characteristics to be tailored, since the copolymers inherit the properties of the homopolymers. Our group has performed extensive work on the homo- and copolymerization of 1,5dioxepan-2-one (DXO). As a homopolymer, PDXO is amorphous and hydrophilic with a glass transition temperature between -35 and -40 °C.1 As a copolymer in the form of poly(DXO-co-LLA) and poly(DXO-co-CL), the material retains mechanical and surface properties that can be optimized according to need.2-4 We have previously shown that the degradation behavior of poly(DXO-co-LLA) commences through hydrolysis.5 The structure and properties of several aliphatic polyesters are under intense investigation for their applications in tissue engineering, a research field often relying on threedimensional porous structures fabricated and used as scaffolds in cell generation.6,7 The structures must offer initial support, contain interconnected pores of the appropriate size, act as temporary guides for the regeneration and formation of new tissue, and be completely degraded and eliminated when the need for the artificial support has diminished. This means that there are immense demands not only on the * Corresponding author. Tel: +46-8-790 82 74. Fax: +46-8-20 84 77. E-mail: [email protected].

structure of the porous scaffolds but also on the materials of which they are composed. Numerous techniques for constructing porous scaffolds are being and have been employed, where the outcome is a threedimensional structure with large surface area and high porosity. The techniques include fabrication of nonwoven structures,8 freeze-drying,9,10 gas foaming,11 phase separation,9,10,12 fused deposition modeling,13 three-dimensional printing techniques,14 and particulate leaching,15,16 as well as combinations of these techniques.17,18 The aim is to create elastic and hydrophilic threedimensional porous scaffolds by utilizing two commonly used monomers, -caprolactone and L-lactide, and an ether ester monomer, DXO. Therefore bulk polymerization of various monomer compositions was performed, yielding copolymers with a wide range of intermolecular structures, sequence lengths, and degrees of randomness enabling easy tailoring of the material properties.

Experimental Section Materials. Stannous 2-ethylhexanoate, SnOct2 (SigmaAldrich, Sweden), was distilled under reduced pressure at 175 °C before use. L-lactide, LLA (Serva Feinbiochemica, Germany) and -caprolactone, CL, (Sigma-Aldrich, Sweden) were purified before use. Diethyl ether (LabScan, Sweden) was dried over molecular sieves, and toluene (Merck, Germany) was dried over a Na-wire. Sodium chloride (NaCl) (Fischer chemicals, Germany) was used after separating the agglomerates by grinding in a mortar. All other chemicals were used as received. Monomers. L-Lactide, LLA, was purified by recrystallization in dry toluene. The monomer was then dried for 24 h

10.1021/bm050190b CCC: $30.25 © 2005 American Chemical Society Published on Web 07/30/2005

Elastomeric Porous Scaffolds Scheme 1. The Three Monomers Used in the Copolymerization Reactions: D, 1,5-dioxepane-2-one (DXO); L, L-lactide (LLA); C, -caprolactone (CL).

under reduced pressure at room temperature prior to polymerization. -Caprolactone (Aldrich) was dried over calcium hydride for at least 24 h at room temperature and was then distilled under reduced pressure. 1,5-Dioxepane-2-one (DXO) was synthesized through a Bayer-Villiger oxidation according to the literature.1 The DXO was then purified by recrystallization from dry ether and two subsequent distillations under reduced pressure. The monomer was dried over calcium hydride for 24 h prior to the final distillation. All monomers were purified before use and were stored in an inert atmosphere. (Monomer structures in Scheme 1.) Polymerization Technique. The desired amounts of monomer, coinitiator, and catalyst were weighed into a previously silanized 25 mL round-bottom flask under nitrogen atmosphere inside a drybox (Mbraun MB 150BG-I). The flask was fitted with a magnetic stirrer bar and was sealed with a three-way valve. The polymerizations were started by immersing the flask into a thermostated oil bath (110 °C) and proceeded for 10 h. The polymer was precipitated in a mixture of cold hexane and methanol (95: 5). Nuclear Magnetic Resonance (NMR). The chemical compositions of the copolymers and the degree of monomer conversion were determined by 1H NMR spectroscopy, comparing the relative intensities of the peaks originating from the comonomers and the resonance peaks from the monomer and polymer. The monomer sequence was determined by 13C NMR spectra. 1H NMR and 13C NMR were obtained using a Bruker AC-400 Fourier transform nuclear magnetic (FT-NMR) operating at 400 and 110.61 MHz, respectively. A 100 mg sample was dissolved in 1 mL deuterochloroform (CDCl3) in a 5-mm-diameter sample tube. Nondeuterated chloroform was used as an internal standard (δ ) 7.26 and δ ) 77.0 ppm). Size Exclusion Chromatography (SEC). SEC was used to monitor the molecular weights of the polymers after polymerization. The polymers were analyzed with a Waters 717 plus autosampler and a Waters model 510 apparatus equipped with two PLgel 10 µm mixed B columns, 300 × 7.5 mm (Polymer Labs., U.K.). Spectra were recorded with an PL-ELS 1000 evaporative light scattering detector (Polymer Labs., U.K.). Millenium32 version 3.05.01 software was used to process the data. Chloroform was used as eluent, at a flow rate of 1.0 mL/min. Polystyrene standards with a narrow molecular weight distribution in the range 580400.000 g/mol were used for calibration. Differential Scanning Calorimetry (DSC). The thermal properties of the synthesized polymers were investigated using a DSC (Mettler Toledo DSC 820 module) under nitrogen atmosphere. To erase the thermal history, the

Biomacromolecules, Vol. 6, No. 5, 2005 2719

specimens were heated from 25 to 180 °C, held there for 2 min, and then cooled to 25 °C at a rate of 10 °C/min. The second scan was used to record the heat of fusion at a heating rate of 10 °C/min. The melting temperatures, Tm, were noted as the maximum values of the melting peaks, and the midpoint temperature of the glass transition was determined as the glass transition temperature, Tg. When evaluating the crystallinty of the copolymers, it was assumed that the only contribution to the heat of fusion in the poly(DXO-co-LLA) and poly(DXO-co-CL) was from the poly(L-lactide) and poly(-caprolactone). Poly(DXO) has previously been shown to be a fully amorphous polymer having a Tg between -35 and -40 °C.1 Tensile Testing. The tensile testing of the polymers was performed with an Instron 5566 equipped with pneumatic grips. The tensile measurements were performed with a load cell with a maximum of 0.1 kN at a crosshead speed of 50 mm/min according to ASTM D 638M-89. The films had a thickness of approximately 1.5 mm, a width of 4 mm, and a gauge length of 35 mm. They were preconditioned before testing according to ASTM D618-96 (40 h at 50 ( 5% relative humidity and 23 ( 1 °C). Five samples from the same porous scaffold were tested for each polymer, and the average thickness of each sample was calculated from five different measurements with a Mitutoyo micrometer. The load at break was chosen as the reference point, since the materials did not yield and in many cases were too weak to induce total stop of the driving motor when breaking. Scanning Electron Microscopy (SEM). Three sample pieces were randomly chosen from each porous scaffold. To ensure minimum loss of structure, the scaffold was immersed in liquid nitrogen and fractured for view of the sample crosssection. Pore size, porosity, efficiency of the salt leaching method, and surface characteristics were evaluated by means of a JEOL JSM-5400 SEM using an acceleration voltage of 15 kV. The samples were mounted on metal studs and sputter-coated with gold-palladium (60%/40%) using a Denton Vacuum Desk II cold sputter etch unit operated at 45 mA for 3 × 15 s. Scaffold Preparation. Three different particulate leaching techniques were evaluated in order to determine the optimum system for scaffold preparation. Sodium chloride (particle size in the range 50-500 µm), ammonium hydrogen carbonate, and poly(ethylene glycol) (Mn ) 1000 g/mol) were used as porogens, and different porogen-to-polymer ratios were investigated. The results show that sodium chloride in a 10:1 weight ratio to the polymer yields the best correlation between porosity and mechanical properties. The copolymers were dissolved in chloroform (CHCl3) to form 5 wt % homogeneous solutions and were subsequently poured over NaCl in molds with a diameter of 9 cm. Each mixture was thoroughly blended using a glass rod and evenly distributed over the bottom of the mold. The mixture was slowly air-dried under a lid for 7 days, after which the scaffold was allowed to dry freely for 7 days, to ensure complete chloroform evaporation. A two-step leaching process was thereafter performed on the now solidified composite. The scaffold and mold were first immersed in methanol for approximately 20 min to wet and swell the

2720

Biomacromolecules, Vol. 6, No. 5, 2005

Odelius et al.

Scheme 2. Example of a Random Copolymer of Poly(DXO-co-CL)

Table 1. Molecular Weight and Composition Characterization of Poly(DXO-co-LLA), Poly(LLA-co-CL), and Poly(LLA-co-CL) mole ratio in feed (%) polymer D

polymer and, in most cases, loosen it from the mold. The free samples were thereafter die-cut into their desired shape and immersed in deionized water to dissolve the salt particles. The water was changed after each 20 min period during 1 h and then once every hour for 5 h, followed by a water exchange 2-3 times every 24 h for 4 days. The porous scaffold was thereafter dried to constant weight in a vacuum and thereafter for another 3 weeks to minimize the amount of residual solvent. In those cases where the scaffolds were fixed to the molds, the scaffolds loosened after 3 days of water immersion, and the same leaching procedure as for the other scaffolds was adopted. To ensure total NaCl leaching, a silver nitrate test was performed, in which the presence of chloride ions in the leaching water is detected by precipitation of AgCl, using a WPA UV-vis spectrophotometer version 1.6, after the instrument had been referenced with deionized water and AgNO3. SEM and gravimetric measurements were also used to confirm total salt leaching. Porosity Determination. The porosity, P, of the scaffolds was determined by measuring the dimensions and mass of the scaffold and was calculated as

[ ]

P) 1-

dp × 100 d

(1)

where dp ) (mp/Vp) is the scaffold density and d is the density of a nonporous film fabricated using the same technique as that used for the scaffold fabrication. Results and Discussion Three different types of copolymers, poly(DXO-co-LLA), poly(DXO-co-CL), and poly(CL-co-LLA), were successfully synthesized as shown schematically in Scheme 2. The reaction conditions were chosen according to previously reported results on similar copolymerizations in which one of the monomers was DXO.19,20 Stannous 2-ethylhexanoate normally acts as a catalyst, starting the polymerizations from impurities in the system.21 In this case, ethylene glycol was used as coinitiator, to ensure better control over the number of chains created and the molecular weight distribution. In all cases, the polymerization temperature was 110 °C, [monomer]/[Sn(Oct)2] ≈ 600, MnTheor ) 80 000 g/mol, and the polymerization time was set to 10 h. The reaction time was chosen to obtain high conversions and at the same time to hinder transesterification reactions that lead to higher molecular weight distributions. Full conversion was however never reached, probably due to the limited diffusion in bulk polymerizations and to the short reaction times. The results are summarized in Table 1.

DL.1 DL.2 DL.3 DL.4 DL.5 DL.6 DC.1 DC.2 DC.3 LC.1 LC.2 LC.3 LC.4 LC.5 L C

L

C

25 75 40 60 50 50 65 35 75 25 90 10 25 - 75 50 - 50 75 - 25 - 10 90 - 25 75 - 40 60 - 50 50 - 75 25 - 100 - 100

polymer composition (%)a D

L

C

14 86 23 77 33 67 48 52 61 39 87 13 26 - 74 50 - 50 75 - 25 - 18 82 - 46 54 - 66 34 - 77 23 - 94 6 - 100 - 100

block length/ sequence length Mnb 133 577 144 400 48 362 109 400 34 700 131 400 186 800 114 797 41 800 155 900 163 800 125 200 138 800 140 500 103 500 172 200

PDIb

Lh D

Lh L

Lh C

1.29 2.0 12.9 1.29 3.1 9.9 1.34 3.7 5.5 1.23 4.1 3.4 1.19 8.0 3.3 1.33 10.6 1.0 1.35 1.4 5.0 1.25 2.7 1.6 1.49 4.1 1.4 1.27 1.4 7.1 1.27 2.7 3.6 1.24 5.9 3.6 1.34 6.5 2.9 1.28 16.7 2.0 1.11 1.29 -

a Molar composition of the copolymers determined by 1H NMR at δ ) D 3.65 ppm, δL ) 5.13 ppm, and δC ) 1.37 or 2.30 ppm. b Determined by SEC (using polystyrene standards).

The molecular weights of the polymers were in most cases higher than expected. This is seen when calibration with polystyrene standards has been used for this type of polymer, and as described by Kricheldorf et al., a twofold increase in molecular weight is common.22 Therefore, the SEC results are mainly used qualitiatively. However, the results indicate that high molecular weight copolymers with narrow molecular weight distributions, ranging from 1.1 to 1.49, have been produced. The chemical compositions of the copolymers were determined using the ratio of the integrated intensities of the 1 H NMR peaks. Thereafter, the weight percent of the different components was determined according to wi (wt % ) )

100 × xi × Mi (100 - xi) × Mj + xi × Mi

(2)

where xi is the molar fraction of component i and Mi and Mj are the molecular weights of components i and j, respectively. In the case of LLA, ML was set to 72.1 g/mol, i.e., the molecular weight of the half-lactide unit. This value was chosen because of the transesterification reactions that occur, possibly giving half of the monomer length as the repeating unit. The reactivities of the three monomer pairs varies considerably when stannous octanoate is used as catalyst, yielding a variety of different copolymers.19,20,23,24 In the case of the poly(DXO-co-CL)-type copolymers consisting of CL and DXO, rC ) 0.6 and rD ) 1.6, the reactivity factors give ideal copolymers which are random in nature, since both the ratios are close to unity. This can be compared to the large difference in reactivity factors found for L-lactide and -caprolactone, rL ) 42 and rC ) 0.36, respectively, and the L-lactide and DXO, rL ) 10 and rD ) 0.1, respectively. In these cases, the copolymers had a tendency to form a more blocklike structure. Sequence analysis was carried out in the carbonyl region using 13C NMR because of its greater sensitivity to sequence effects than 1H NMR.25 Typical spectra are shown in Figure

Elastomeric Porous Scaffolds

Biomacromolecules, Vol. 6, No. 5, 2005 2721

was determined using the following equations: L hD )

IDDD + ILDD +1 IDDL + ILDL

(3)

L hL )

ILLL + IDLL +1 ILLD + IDLD

(4)

where, e.g., IDLD is the intensity of the peak representing the DLD triad.19 The sequence lengths of poly(DXO-co-CL)4 and poly(LLA-co-CL)25,26 were determined in the same manner. In the case of poly(DXO-co-CL), the 13C NMR spectra showed four different peaks, corresponding to the homopolymer peaks and two dyad peaks, i.e., the transitions between D and C and vice versa. This necessitated a small alteration in the standard equation used for poly(DXO-coLLA) as follows: L hC )

ICC ICC +1) +1 ICD IDC

(5)

L hD )

IDD IDD +1) +1 IDC ICD

(6)

In the case of poly(LLA-co-CL), the carbonyl signals LLC and CLL overlap, yielding equations L hC ) L hL )

Figure 1. Expanded 1C NMR of the carbonyl region: (a) poly(DXOco-LLA), (b) poly(DXO-co-CL), (c) poly(LLA-co-CL).

1, where the homopolymer carbonyl peaks for all copolymers are at the highest and lowest fields in the spectrum. It should be noted that the sequence length of a lactic acid unit is the half-lactide unit, so that a random 50/50 copolymer that has not undergone any transesterification reaction should have a sequence length of 2. For the poly(DXO-co-LLA), eight different triads can be distinguished, and the sequence length

ICCC +1 ICCL

ILLL + (ILLC + ICLL)/2 ICLC + (ILLC + ICLL)/2

(7)

+1

(8)

All sequence length values are given in Table 1. It is obvious here that the monomer composition in the feed greatly affects the sequence lengths of the copolymer blocks, as confirmed by 1H NMR. Similar results have been found by other groups, e.g., Choi et al.27 who found LL values of 7.5, 3.7, and 3.4 when the feed ratios of LLA/CL were 80/20, 65/35, and 50/50, respectively, at 140 °C. It has previously also been shown that an increase in temperature increases the likelihood for transesterification reactions and affects the reactivities of the monomers, thereby decreasing the block lengths.24 The block lengths presented here for a monomer feed ratio of LLA/CL 50/50 agree well with values previously reported.28,29 The results show that the incorporation of the less reactive monomer in the copolymer is lower than expected, so that an excess of the less reactive monomer must be added in order to reach the desired copolymer compositions. The compositional analysis of DXO/CL copolymers, Table 1, show that the polymers are random in their nature, in correspondence with literature, and the sequence lengths are therefore only presented for comparison reasons. Characterization Differential Scanning Calorimetry (DSC). The glass transition temperature, melting temperature, and crystallinity of the copolymers were determined and compared using

2722

Biomacromolecules, Vol. 6, No. 5, 2005

Odelius et al.

Figure 2. Tg as a function of the ratio of DXO, in the case of poly(DXO-co-CL), measured by DSC and the proportion of LLA in the cases of poly(LLA-co-CL) and of poly(DXO-co-LLA). Table 2. Thermal Properties of Poly(DXO-co-LLA), Poly(LLA-co-CL), and Poly(LLA-co-CL) polymer composition (%)a polymer D DL.1 DL.2 DL.3 DL.4 DL.5 DL.6 DC.1 DC.2 DC.3 LC.1 LC.2 LC.3 LC.4 LC.5 L C a

L

C

14 86 23 77 33 67 48 52 61 39 87 13 26 - 74 50 - 50 75 - 25 - 18 82 - 46 54 - 66 34 - 77 23 - 94 6 - 100 - 100

∆Hm [J/g]

Tga [°C]

Tm [°C]

23.2 3 -5.2 -20.3 -26.1 -34.7 -60.5 (-61.6) -57.1 (-54.7) -48.5 (-47.1) -51.7 -44.6 -28.3 -9.1 28.1 53.8 -68.7

142.6 132 36.8 15.3 43.9 41.1 154.2 173.3 57.5

L

C

Xc [%]

22.6 24.3 18.6 20.0 77.2 55.3 0.8 0.56 64.1 68.9 81.0 58.1

The number in parentheses is calculated by the Fox equation.

DSC, in order to establish how changes in composition alter their thermal properties (Table 2). The crystallinity, Xc, was calculated according to29 Xc )

∆Hf ∆H0f

(9)

where ∆Hf (J/g of the crystalline polymer) is the enthalpy of fusion of the specimen and ∆H0f is the enthalpy of fusion of a 100% crystalline polymer. For PCL and PLLA, ∆H0f is 139.530 and 93 J/g,31 respectively. A single glass transition could be observed in all thermograms located between the glass transition temperatures of the corresponding homopolymers and indicates a continuous amorphous phase due to the randomness of the copolymers. The short blocks are not long enough to induce two separate amorphous phases as can be seen for lower conversions32 (Table 2 and Figure 2). For the copolymers poly(DXO-co-LLA) and poly(LLA-coCL), the Tg increases with increasing LLA content, although not linearly. This is because, at very high and very low ratios of one of the monomers, the glass transition temperature is

not intermediate between those of the two homopolymers, but is instead increasingly affected by the monomer in abundance. The Fox equation does not apply to these copolymers, because they are not truly random. The generally low Tg of the CL- or DXO-rich copolymers is attributed to an increase in the supple segments, methylene groups, in the chains, which enhances the mobility.33 For poly(DXOco-CL), Tg increases linearly with increasing DXO content, and the Fox equation applies well, indicating a random structure. The results show clear variations in Tg with different monomers, indicating that the Tg and the thermal properties can be tailored by choice of monomer composition. For copolymers of poly(DXO-co-LLA), the crystallinity depends on the block length of LLA; the crystallinity decreases with increasing DXO content, which in turn yields shorter LLA blocks. When the proportion of DXO exceeds 23%, the copolymer becomes amorphous. For poly(DXOco-CL), the degree of crystallinity depends on the length of the CL sequences, which decreases with increasing DXO content. The CL sequences, however, need not be as long as the LLA blocks to crystallize (Table 1). Nevertheless, the crystallinity decreases to almost zero when the mole ratio of DXO is raised from 26 to 50 mol %. At DXO molar ratios above 50%, the copolymer is amorphous. In poly(LLA-co-CL), crystals are built either by LLA or CL depending on the composition. It was seen that, when the molar ratio of CL was about 50% or higher, the melting temperature of the crystals was closer to that of the crystals in the homopolymer, implying that the crystals formed are from the CL blocks. In contrast, when the LLA proportion was high, the melting temperature was close to that of pure PLLA, and the LLA blocks are thus most likely responsible for the crystallinity. For both poly(DXO-co-CL) and poly(DXO-co-LLA), the Tm decreases with increasing DXO content, because of a shortening of the crystalline blocks and thus the formation of smaller and more imperfect crystals with a lower melting point. Scaffold Preparation. To investigate the scaffold preparation technique and the structure and properties of the scaffolds created, a salt particulate leaching technique was applied to the homo- and copolymers. It is well-established that the polymer salt composite has a tendency to adhere to a glass mold, so that the samples must be formed after leaching. Methanol has been used here to solve this problem. Immersing the scaffold in methanol until it is totally soaked facilitates the separation of the scaffold from the mold and enables the design of scaffolds with different shapes and forms to be studied without destroying or damaging the porous structure. This is because methanol has the ability to wet and slightly swell the polymers without leaching the salt. Unfortunately, all copolymers cannot be used as scaffold materials. More specifically, for the copolymers of poly(DXO-co-LLA), a large loss of structure was seen when the amount of DXO reached 61 mol %. These copolymers first adhered to the mold, and after complete salt leaching, the scaffold structure collapsed completely. The same phenomenon was seen for the poly(DXO-co-CL), for which the

Elastomeric Porous Scaffolds

Biomacromolecules, Vol. 6, No. 5, 2005 2723

Figure 3. SEM micrographs of (a) DL.1, (b) DL.2, (c) DL.3, (d) DL.4, (e) DC.1, (f) LC.1, (g) LC.2, (h) LC.3, (i) LC.4, (j) LC.5, (k) L, (l) C3, (m) DL.2, (n) DL.3, (o) LC.4.

problem occurred at even lower contents of DXO. More interestingly, LC.2, a copolymer of poly(LLA-co-CL), lost its structure in a different manner. This copolymer was also first fixed to its mold, but it could easily be removed and leached, whereafter it folded and fixed to itself. No explanation of this phenomenon can be found in the chemical structure of this particular copolymer, but the polymer obviously became too sticky. An ongoing discussion into the possible disadvantages of creating porous scaffolds using salt particulate leaching techniques has been concerned with whether the pores are interconnected and whether the regular shape of the created pores leads to problems in cell seeding and cell culture. It is shown here that the porous scaffolds manufactured by our technique (Figure 3a-l) have interconnected pores, so that

growth throughout the scaffolds should be possible, and cell culture studies are now in progress. It can also be seen that the scaffolds feature regular structures with homogeneously distributed pores and that the shape and structure of the pores are influenced by, but not completely determined by, the salt particles used. This can be seen as small variations in the shape of pores. The size distribution of the pores has been evaluated using image analysis, in which their diameter was plotted against their occurrence. It was seen that the ground salt gave rise to small variations in the pore diameter. Generally, the pore size is between 20 and 300 µm, with the largest amount of pores being in the 20-160 µm range, and the porosity of the scaffolds was in the range 89.294.6% (Table 3). As can be seen in Figure 3m-o (examples of surface micrographs), the surface topography of the porous

2724

Biomacromolecules, Vol. 6, No. 5, 2005

Odelius et al.

Table 3. Porous Scaffold Characteristics

polymer

scaffold stabilitya

salt leaching efficiency [days]b

DL.1 DL.2 DL.3 DL.4 DL.5 DL.6 DC.1 DC.2 DC.3 LC.1 LC.2 LC.3 LC.4 LC.5 L C

++ ++ ++ + + ++ + ++ ++ ( ++

2 3 1 4 2 2 3 2 2 4 2 4 2 4 2 2

pore size of largest occurrence pores [µm]c

porosity [%]

20-100 20-140 20-200 20-200 40-160 20-200 20-200 20-120 20-100 20-160 0-120

92.6 93.4 92.2 90.5 91.1 91.5 94.5 94.6 90.7 89.2 91.5

a The stability of the scaffold after salt leaching is denoted as follows: ++ no fragmentation and total form stability, + no fragmentation but adhere to mold, ( a few fragments comes off and form stability, - the scaffold form is not preserved. b Determined by UV-vis and measured as the number of days after which the salt concentration was lower than 0.001 mg salt per ml water at all wavelengths. c Determined by SEM picture evaluation.

scaffolds shows equivalent porosities and interconnectivities as the cross-sections. Total salt leaching was determined by measuring the concentration of AgCl precipitate in the leaching media. Concentrations as low as 0.001 mg/mL can be detected, and as can be seen in Table 3, for most polymer scaffolds the concentration of precipitate in the water was lower than 0.001 mg/mL after 2 days of salt leaching. In all cases, no salt could be detected after 4 days of leaching. SEM micrographs showed no visible salt residues, and gravimetric measurements confirmed total salt leaching. Porous scaffolds were investigated by tensile testing in order to determine how variations in the copolymer composition alter the mechanical properties of the materials. As mentioned above, the copolymers created here showed a single glass transition temperature indicating one amorphous phase and, in most cases, a crystalline region of one of the monomers. These soft and hard parts are intertwined through chain segments incorporated in both regions and thus create thermoplastic elastomeric-like properties.28 The interpretation of the tensile properties of the scaffolds does not account for the effects from the variations in pore size. The pore size variation is altering the mechanical properties; however, the consequence of this seems to be relatively small compared with the changes of properties due to the differences in composition. In poly(DXO-co-LLA), the hard domain would be made up of LLA moieties, and it was seen that, when the proportion of LLA in the scaffolds was decreased, thereby increasing the amount of supple segments in the chains and decreasing the crystallinity, the material became more flexible with a resulting reduction in strength but an improvement in the maximum elongation (Figure 4). Scaffolds constructed of poly(LLA-co-CL) exhibited a more complex behavior with variation in the LLA content.

Figure 4. (a) Maximum stress dependent on the DXO content in poly(DXO-co-LLA) and poly(DXO-co-CL) and of the LLA content in poly(LLA-co-CL) (b) Maximum strain dependent on the DXO content in poly(DXO-co-LLA) and poly(DXO-co-CL) and of the LLA content in poly(LLA-co-CL). Table 4. Tensile Properties of the Porous Scaffold polymer composition (%) polymer D L C DL.1 DL.2 DL.3 DL.4 DC.1 LC.1 LC.3 LC.4 LC.5 L C

14 23 33 48 26 -

86 77 67 52 18 66 77 94 100 -

74 82 34 23 6 100

stress at max load [MPa]

strain at max load [%]

0.53 ( 0.2 11 ( 4.1 0.33 ( 0.03 58 ( 5.4 0.19 ( 0.02 109 ( 4.2 0.06 ( 0.01 113 ( 20 0.64 ( 0.20 15 ( 6.3 0.46 ( 0.04 53 ( 9.2 0.09 ( 0.01 176 ( 21 0.14 ( 0.04 87( 5.5 0.23 ( 0.04 4.5 ( 0.41 0.12 ( 0.04 1.4 ( 0.48 0.59 ( 0.08 9.6 ( 1.7

modulus [MPa] 13 ( 0.55 1.5 ( 0.21 0.31 ( 0.05 0.12 ( 0.01 13 ( 2.2 4.3 ( 0.65 0.58 ( 0.49 0.24 ( 0.05 11 ( 2.4 9.6 ( 0.48 16 ( 2.4

When the proportion of LLA was low, the scaffolds were strong and fairly flexible. When the LLA content was increased up to 66%, the strength decreased, whereas the flexibility increased considerably. However, when the LLA content was increased even more, the flexibility again decreased. At 94% LLA, the material was brittle and had tensile properties that are slightly enhanced, but which were almost identical to those of pure PLLA. In general, a PCL homopolymer has a higher tensile strength than a copolymer of poly(LLA-co-CL) (Table 4 and Figure 4), probably because the crystals in the copolymer are disturbed by the incorporation of another monomer. This would contribute to the reduced flexibility of the homopolymer compared to most of the copolymers. The higher modulus of the copolymers compared to scaffolds of PLLA homopolymer might be explained by cracks induced in the very brittle PLLA scaffolds by die-cutting, or it might merely be an artifact of the different characteristics of the polymers.

Elastomeric Porous Scaffolds

In the poly(DXO-co-CL), only one type of scaffold could be examined by tensile testing because of the collapse of the porous structure after leaching. This scaffold was hard and brittle with a low elongation and a high tensile strength compared to the poly(DXO-co-LLA) scaffolds with similar molar ratios of DXO in the copolymer. In comparison with pure PCL scaffolds, the strength is similar, but the elongation at break and flexibility are higher, while the modulus is somewhat lower (Table 4). In general, the mechanical properties of porous scaffolds never reach those of a nonporous material, since pores create voids which are similar to fractures. Conclusions Porous scaffolds were successfully fabricated through a solvent casting and particulate leaching technique, in which methanol was used to wet and swell the composite before leaching, thereby successfully separating the scaffold from its mold. The scaffolds are highly porous and possess interconnected pores. Different copolymers were synthesized by ring-opening polymerization of various combinations of DXO, LLA, and -CL. The copolymers show tensile properties varying from brittle and hard to flexible and soft in scaffolds made of different monomers and compositions. Interestingly, poly(LLA-co-CL) shows a maximum flexibility and a minimum strength at 66% LLA. This is attributed to the block lengths of the two monomers, where the LLA segments are too short to crystallize but long enough to inhibit crystallization of CL, and consequently, this yields a soft and flexible material. All the copolymers showed a single glass transition temperature, indicating a continuous amorphous phase constructed of both monomers with a Tg that changed according to the copolymer composition. The variation was linear for poly(DXO-co-CL) and nonlinear for high concentrations of LLA in poly(DXO-co-LLA) and poly(LLA-coCL). This is due to the long LLA blocks that lead to a thermal behavior similar to that of pure PLLA. The crystallinity and melting temperature of the copolymers were also strongly affected by the monomer types and ratios. The large range of properties made available by these copolymers show that it is possible to fabricate scaffolds to meet a range of different applications. Acknowledgment. The authors gratefully acknowledge the Foundation for Strategic Research for financial support of this work. Note Added after ASAP Publication. This article was released ASAP on July 30, 2005. The last sentence in

Biomacromolecules, Vol. 6, No. 5, 2005 2725

paragraph four of the Experimental Section has been revised. The correct version was posted on August 10, 2005. References and Notes (1) Mathisen, T.; Albertsson, A.-C. Macromolecules 1989, 22, 38383842. (2) Ryner, M.; Albertsson, A.-C. Biomacromolecules 2002, 3, 601-608. (3) Stridsberg, K.; Albertsson, A.-C. J. Polym. Sci., Part A: Polym. Chem. 2000, 38, 1774-1784. (4) Andronova, N.; Finne, A.; Albertsson, A.-C. J. Polym. Sci., Part A: Polym. Chem. 2003, 41, 2412-2423. (5) Stridsberg, K.; Albertsson, A.-C. Polymer 2000, 41, 7321-7330. (6) Agrawal, C. M.; Ray, R. B. J. Biomed. Mater. Res. 2001, 55, 141150. (7) Hutmacher, D. W. Biomaterials 2000, 21, 2529-2543. (8) Freed, L. E.; Vunjak-Novakovic, G.; Langer, R. J. Cell. Biochem. 1993, 51, 257-264. (9) Schugens, C.; Maquet, V.; Grandfils, C.; Jerome, R.; Teyssie, P. Polymer 1996, 37, 1027-1038. (10) Zhang, R.; Ma, P. X. J. Biomed. Mater. Res. 1999, 44, 446-455. (11) Mooney, D. J.; Baldwin, D. F.; Suh, N. P.; Vacanti, J. P.; Langer, R. Biomaterials 1996, 17, 1417-1422. (12) Nam, Y. S.; Park, T. G. J. Biomed. Mater. Res. 1999, 47, 8-17. (13) Hutmacher, D. W.; Schantz, T.; Zein, I.; Ng, K. W.; Teoh, S. H.; Tan, K. C. J. Biomed. Mater. Res. 2001, 55, 203-216. (14) Giordano, R. A.; Wu, B. M.; Borland, S. W.; Cima, L. G.; Sachs, E. M.; Cima, M. J. J. Biomater. Sci., Polym. Ed. 1996, 8, 63-75. (15) Gogolewski, S.; Pennings, A. J. Colloid Polym. Sci. 1983, 261, 477484. (16) Mikos, A. G.; Sarakinos, G.; Leite, S. M.; Vacanti, J. P.; Langer, R. Biomaterials 1993, 14, 323-330. (17) Hou, Q.; Grijpma, D. W.; Feijen, J. Biomaterials 2003, 24, 19371947. (18) Harris, L. D.; Kim, B.-S.; Mooney, D. J. J. Biomed. Mater. Res. 1998, 42, 396-402. (19) Lofgren, A.; Renstad, R.; Albertsson, A.-C. J. Appl. Polym. Sci. 1995, 55, 1589-1600. (20) Albertsson, A.-C.; Lofgren, A. J. Macromol. Sci., Pure Appl. Chem. 1995, A32, 41-59. (21) Nijenhuis, A. J.; Grijpma, D. W.; Pennings, A. J. Macromolecules 1992, 25, 6419-6424. (22) Kricheldorf, H. R.; Eggerstedt, S. Macromol. Chem. Phys. 1998, 199, 283-290. (23) Pitt, C. G.; Gratzl, M. M.; Kimmel, G. L.; Surles, J.; Schindler, A. Biomaterials 1981, 2, 215-220. (24) Grijpma, D. W.; Pennings, A. J. Polym. Bull. 1991, 25, 335-341. (25) Kricheldorf, H. R.; Jonte, J. M.; Berl, M. Makromol. Chem., Suppl. 1985, 12, 25-38. (26) Kricheldorf, H. R.; Kreiser, I. J. Macromol. Sci., Chem. 1987, A24, 1345-1356. (27) Choi, E.-J.; Park, J.-K.; Chang, H.-N. J. Polym. Sci., Part B: Polym. Phys. 1994, 32, 2481-2489. (28) Grijpma, D. W.; Zondervan, G. J.; Pennings, A. J. Polym. Bull. 1991, 25, 327-333. (29) Tsuji, H.; Mizuno, A.; Ikada, Y. J. Appl. Polym. Sci. 2000, 76, 947953. (30) Crescenzi, V.; Manzini, G.; Calzolari, G.; Borri, C. Eur. Polym. J. 1972, 8, 449-463. (31) Fischer, E. W.; Sterzel, H. J.; Wegner, G. Kolloid Z. Z. Polym. 1973, 251, 980-990. (32) Vanhoorne, P.; Dubois, P.; Jerome, R.; Teyssie, P. Macromolecules 1992, 25, 37-44. (33) Fukuzaki, H.; Yoshida, M.; Asano, M.; Kumakura, M.; Mashimo, T.; Yuasa, H.; Imai, K.; Yamanaka, H. Polymer 1990, 31, 2006-2014.

BM050190B