María J.
Sánchez-Fernández
a,
Manon
Peerlings
a,
Rosa P.
Félix Lanao
b,
Johan C. M. E.
Bender
b,
Jan C. M.
van Hest
c and
Sander C. G.
Leeuwenburgh
*a
aDepartment of Dentistry–Regenerative Biomaterials, Radboud Institute for Molecular Life Sciences, Radboud University Medical Center, 6525 EX Nijmegen, The Netherlands. E-mail: Sander.Leeuwenburgh@radboudumc.nl
bGATT Technologies BV, 6525 ED, Nijmegen, The Netherlands
cDepartment of Bio-Organic Chemistry, Institute for Complex Molecular Systems, Eindhoven University of Technology, 5600 MB Eindhoven, The Netherlands
First published on 13th July 2021
To create a novel generation of barrier membranes with bone-adhesive properties, three-component membranes were successfully developed using a solvent-free approach by combining an occlusive polyester backing layer with a bone-adhesive fibrous gelatin carrier impregnated with calcium-binding alendronate-functionalized poly(2-oxazoline)s (POx-Ale). The mechanical properties of these novel membranes were similar to other commercially available barrier membranes. In contrast, the adhesion of our membranes towards bone was by far superior (i.e. 62-fold) compared to conventional commercially available Bio-Gide® membranes. Moreover, alendronate-functionalized membranes retained their bone-adhesive properties under wet conditions in phosphate-buffered saline (PBS) solutions with and without collagenase. Finally, the in vitro degradation of the membranes was studied by monitoring their weight loss upon immersion in PBS solutions with and without collagenase. The membranes degraded in a sustained manner, which was accelerated by the presence of collagenase due to enzymatic degradation of the carrier. In conclusion, our results show that surface functionalization of barrier membranes with alendronate moieties renders them adhesive to bone. As such, the biomaterials design strategy presented herein opens up new avenues of research on bone-adhesive membranes for guided bone regeneration.
We propose that the next generation of biodegradable barrier membranes should become adhesive to bone via a molecular mechanism to ensure fixation of the membrane. This approach will minimize gingival ingrowth, maximize the extent of bone regeneration, and improve the overall clinical handling and reliability of barrier membranes. To render barrier membranes adhesive to bone, the adhesive material should be rationally designed based on the chemical composition of bone,17 which consists of an organic collagenous matrix (≈30 wt%) reinforced by finely dispersed calcium phosphate mineral nanocrystals (CaP) (≈70 wt%). We propose that synthetic polymers such as poly(2-oxazoline)s (POx) are particularly attractive to this end since these polymers can be functionalized with pendant chemical groups of high affinity with the organic and/or inorganic components of bone tissue. In our study, we selected bisphosphonates as primary bone-bonding moiety, since bisphosphonates such as alendronate (Ale) are anti-osteoporotic molecules known for their exceptionally strong chemical affinity for calcium cations in hydroxyapatite, the mineral phase of bone.18–20
Herein, we designed novel barrier membranes consisting of three components: (i) a polyester backing layer and (ii) a fibrous gelatin carrier layer impregnated with (iii) calcium-binding alendronate-functionalized poly(2-oxazoline)s (POx-Ale). We hypothesized that this approach would allow for independent tuning of bone adhesiveness (by controlling the amount of alendronate groups) and prevention of gingival ingrowth (by controlling the degradation rate of the polyester backing layer). Ideally, this approach would ensure short-term bone adhesiveness of the membranes (in terms of days) with a long-term occlusive function (in terms of several months) of the resulting membranes. Poly(2-oxazoline)s were selected as a core polymer in regards of their high functionalization possibilities along the entire polymeric backbone, and their tunable versatility by copolymerization.21–23 POx polymers were functionalized with both alendronate (P(EtOx-Ale)) and hydroxyl moieties (P(EtOx-OH-Ale)) in view of the strong binding affinity of both moieties for calcium in hydroxyapatite bone mineral nanocrystals.24,25 A commercially available fibrous gelatin carrier (GELITA TUFT-IT®) was selected in view of the favorable biological response of gelatin-based medical devices. Furthermore, the selected fibrous gelatin carrier combines the beneficial features of gelatin-based membranes with the easy handling properties of a non-woven material.26 In addition, the fibrillar architecture of the gelatin carrier facilitated impregnation of the bone-adhesive polymers as a dry powder. Dry powder deposition was selected to improve the future shelf life of the resulting membranes. In addition, unlike impregnation with organic or inorganic solvents, dry deposition does neither affect the structural integrity nor the mechanical properties of the carrier, since this method does not create any internal mechanical stress.27 Finally, the backing layer was prepared from biodegradable polyesters to enhance the mechanical properties at early stages of bone regeneration, while enabling biodegradation of the membranes at later stages of bone regeneration. Specifically, the backing layer consisted of a blend of poly(L-lactide-co-ε-caprolactone) (P(LA-CL)), poly(D,L-lactic-co-glycolic acid) (PLGA), and poly(2-propyl-2-oxazoline) functionalized with pendant amine groups (P(PropOx-NH2)), which was processed using hot-melt extrusion. P(LA-CL) copolymer provided the desired strength and flexibility to the backing,28 whereas PLGA was added to tune the long-term degradation rate of the resulting membranes.29,30 P(PropOx-NH2) was incorporated to the polyester backing layer to facilitate its attachment to the impregnated carrier since the amine groups present in the backing could interact by hydrogen bonding with carboxylic acid groups present in the carrier as well as hydroxyl and bisphosphonate groups from the impregnated POx polymers.
The membranes were characterized in detail regarding their wettability and mechanical properties. Moreover, adhesion of these membranes to apatite-coated model surfaces as well as bone was assessed by lap-shear tests. Finally, the degradation kinetics of these membranes were evaluated in vitro in phosphate-buffered saline (PBS) with and without collagenase to study the degradation mechanism in more detail.
1H NMR and 31P NMR spectroscopy were used to determine the degree of modification of the different substitutions in the polymers.24 NMR spectra were recorded on a Bruker Avance III (400 MHz) spectrometer in the indicated solvent at 25 °C. 1H NMR data are reported as chemical shifts (given in parts per million (ppm) with respect to tetramethylsilane as standard), multiplicity (br = broad), integration, and assignment. 1H NMR and 31P NMR spectra are presented in Fig. S1 (ESI†). The number average molecular weights (Mn) were recorded on a Bruker Microflex LRF matrix-assisted laser desorption ionization time-of-flight mass spectrometry (MALDI-TOF MS) system as described previously.24 All mass spectra were obtained in the positive ion mode. α-Cyano-4-hydroxycinnamic acid (CHCA) was used as a matrix in THF (10 mg mL−1). Polymer samples were dissolved in THF/MeOH (1:1, 10 mg mL−1), and analyte solutions were prepared by mixing 10 μL of matrix and 1 μL of the polymer sample. Samples were applied using the dried droplet method. To determine the dispersity of the polymers (Đ), size exclusion chromatography (SEC) was performed on an automated Shimadzu HPLC system as described previously,24 with a PLgel 5 μm MIXED-D column at 50 °C, using N,N-dimethyl acetamide (DMA) containing 50 mM LiCl as the eluent at a flow rate of 0.6 mL min−1. Poly(methyl methacrylate) (PMMA) was used as standards.
The POx polymers were analyzed with regard to the amount of alendronate (Ale), hydroxyl (OH), and carboxyl (COOH) functional groups present in the polymers. Table 1 summarizes the analytical data of the POx polymers.
Polymer | 1H NMR (mol%) | MALDI-TOF MS | Đ | ||||
---|---|---|---|---|---|---|---|
EtOx | OH | COOH | Ale | M n (kDa) | |||
P(EtOx) | 100 | 50 | |||||
P1e | P(EtOx70-Ale30) | 70 | 3 | 27 | 19.6 | 1.11 | |
P2e | P(EtOx70-OH10-Ale20) | 70 | 11 | 1 | 18 | 14.9 | 1.11 |
The phase transition temperature of both POx-Ale P1e and P2e polymers as well as P(EtOx) was determined using differential scanning calorimetry (DSC, Mettler Toledo). The measurements were carried out at a heating rate of 10 °C min−1 under nitrogen atmosphere (sample masses ≈ 2 mg). The glass transition temperatures (Tg) were obtained as the midpoint of the intersection of the tangent before and after baseline shifting. All temperatures were determined from the second heating scan.
A field emission scanning electron microscope (SEM, Zeiss Sigma 300) was used to evaluate the morphology of the membranes at an accelerating voltage of 3 kV. To this end, the membranes were placed on carbon tape and sputter-coated with a 10 nm layer of gold (Edwards Pirani 501) at 1 kV of accelerate voltage and a current of 20 mA for 120 s. The infrared spectra of the membranes were obtained by Attenuated Total Reflectance-Fourier Transform Infrared (ATR-FTIR) spectroscopy (PerkinElmer, Spectrum Two). The density of the membranes was calculated as the average of the ratio between sample weight and volume in sextuplicate (n = 6).
The wettability of the membranes was quantified by measuring the water contact angle upon initial contact using a contact angle goniometer (Theta Lite) employing the sessile drop method. For each measurement, 2 μL of Milli-Q water was placed onto the membranes at room temperature. The spreading of the droplet was recorded using a high speed video camera and the angle was determined using Laplace Young fitting. The angle was measured after 1 second of contact as the average between the left and right angles in triplicate (n = 3).
The crystallographic structure of apatite-coated disks was determined by means of thin-film X-ray diffraction (thin-film XRD, X’Pert3 PANalytical, Philips). The measurements were performed at 45 kV acceleration and 30 mA current, scanning from 25 to 35° 2θ at a rate of 0.02° s−1. Reflection peaks were characteristic of crystalline apatite (Fig. S3A, ESI†).33,34 The infrared spectrum of the coating was obtained by ATR-FTIR spectroscopy, which also confirmed the apatitic nature of the coatings (Fig. S3B, ESI†).35 The surface morphology of the disks was examined using scanning electron microscopy, which confirmed the homogeneous coverage of the substrate (Fig. S4A, ESI†). The thickness of the coating was determined using a profilometer (Innowep, universal-surface tester).
Flat bone specimens were immersed in 500 mL of Sakura reagent TDE™ 30 and decalcified using the Sakura TDETM 30 electrode system for 24 h. Thereafter, they were kept frozen at −20 °C until further use. The demineralization process was confirmed using XRD (Fig. S3A, ESI†) and ATR-FTIR spectroscopy (Fig. S3B, ESI†) and the samples were inspected visually (Fig. S4BC, ESI†).
A field emission scanning electron microscope was used to evaluate the morphology of the membranes after 14 days of immersion in PBS with or without collagenase.
Fig. 2A–F show the surface morphology of the three-component membranes. The fibrillar architecture of the pure gelatin carriers was clearly visible in Fig. 2A, whereas the smooth appearance of the polyester backing layer was recognizable in Fig. 2B. The blank membrane (Fig. 2C), consisting of a gelatin carrier on top of a backing layer, was covered by a fibrillar gelatin layer, although the membrane structure was compacted upon adhering the backing layer to the carrier. All POx-covered membranes (Fig. 2D–F) revealed a fibrillar surface architecture. Successful impregnation of POx polymers into the fibrous gelatin carrier was evidenced by abundant coverage of gelatin fibers with POx polymers, forming a highly intertwined POx-gelatin structure. In addition, irregularly shaped POx deposits were heterogeneously dispersed throughout the superficial layer of the POx-coated membranes.
The backing layer was attached by heating at 150 °C, which was above both the glass transition temperature (60 °C) and melting point (104–114 °C) of P(LA-CL), the main component of the backing layer,38,39 and above the glass transition temperature of the POx polymers, as depicted in Table 2 and Fig. S6A (ESI†). Generally, POx polymers did not melt at temperatures below 200 °C.
Polymer | T g (°C) | |
---|---|---|
Onset | Midpoint | |
P(EtOx) | 52.19 | 54.45 |
P(EtOx-Ale) | 54.54 | 71.15 |
P(EtOx-OH-Ale) | 54.53 | 62.94 |
In Fig. S6B (ESI†) the infrared spectra of the different membranes are combined. The small peak at 1728 cm−1 corresponds to carbonyl stretching from esters in P(EtOx-Ale) and P(EtOx-OH-Ale). The peaks at 1625, 1541, and 1238 cm−1, on the other hand, are attributed to the amides I, II, and III as present in both the gelatin carrier and the three types of POx polymers. The peaks at 1426 cm−1 and 1164 cm−1 present in all membranes correspond to bending of C–H bonds and stretching of C–N bonds, respectively. The peak at 915 cm−1 accounts for stretching of PO and P–OH bonds in P(EtOx-Ale) and P(EtOx-OH-Ale).40
Fig. S7 (ESI†) shows the water contact angle of the different membranes after 1 s of contact. Membranes containing P(EtOx-Ale) and P(EtOx) gave a contact angle of 0 °C, absorbing the water drop within 1 s. These membranes were significantly more hydrophilic than the blank and membranes impregnated with P(EtOx-OH-Ale), which revealed contact angles of 96.3 ± 0.3 and 74.2 ± 8.2, respectively. After 6 s of contact, the water drop was absorbed by all types of membranes.
Fig. 3 (A) Representative stress–strain curves of the different membranes. (B) Tensile strength and tensile modulus of the membranes. Values represent the mean ± standard deviation (n = 6). |
The tensile strength and modulus of the different membranes are depicted in Fig. 3B. Both the maximum strength and tensile modulus were 3-fold higher for membranes comprising a backing layer (strength: 2.0 ± 0.3 MPa, modulus: 42.0 ± 8.9 MPa for blank) compared to pure fibrous gelatin carriers (strength: 0.7 ± 0.2 MPa, modulus: 13.3 ± 3.4). In general, the impregnation of POx polymers into membranes comprising a polyester backing did not further reinforce these membranes. However, membranes impregnated with hydroxyl-functionalized P(EtOx-OH-Ale) polymers showed 2-fold higher tensile strength and tensile modulus (strength: 3.9 ± 0.7 MPa, modulus: 76.2 ± 14.7 MPa) compared to hydroxyl-free P(EtOx-Ale) (strength: 2.3 ± 0.2 MPa, modulus: 44.3 ± 6.9 MPa) and P(EtOx) polymers (strength: 1.7 ± 0.4 MPa and modulus: 50.5 ± 16.9 MPa). Overall, the tensile strength of hydroxyl-functionalized P(EtOx-OH-Ale) polymers was comparable to values reported for commercially available membranes such as Bio-Gide® (4.6 ± 0.9 MPa) and Ossix Plus® (5.1 ± 2.5 MPa).36
The work of adhesion, as a measure for energy dissipation during lap shearing, is shown for the various types of membranes in Fig. 4C and D. Membranes impregnated with P(EtOx) showed a significantly higher work of adhesion (up to 83 J m−2) to CaP-coated Ti disks (Fig. 4C) compared to all other experimental groups. In contrast, the work of adhesion of all POx-impregnated membranes to bone (Fig. 4D) was comparable (up to 75–79 J m−2). Generally, the work of adhesion of impregnated membranes to both types of mineralized substrates was much higher compared to mineral-free control substrates. The POx-free blank membranes and Bio-Gide® revealed extremely low amounts of energy dissipation during lap-shearing to both mineralized and mineral-free control substrates.
All POx-impregnated membranes failed due to cohesive failure upon lap-shear testing. Fig. S8 (ESI†) shows FTIR spectra of the surface of bone specimens after lap-shear testing of all experimental groups. These spectra revealed that bone samples were covered with membrane remnants for POx-impregnated membranes, which were absent on bone samples after detachment of POx-free gelatin carriers, corresponding to adhesive failure for blank controls.
The mechanical parameters of the tested membranes are summarized in Table 3.
Prototype | Tensile strength (MPa) | Tensile modulus (MPa) | Shear strength (kPa) | Work of adhesion (J m−2) | ||
---|---|---|---|---|---|---|
CaP coated Ti | Bone | CaP coated Ti | Bone | |||
Blank | 2.03 ± 0.31 | 41.95 ± 8.89 | 3.99 ± 2.87 | 2.92 ± 1.61 | 3.10 ± 3.42 | 8.62 ± 5.24 |
P(EtOx) | 1.73 ± 0.37 | 50.50 ± 16.88 | 63.51 ± 21.36 | 15.18 ± 6.40 | 83.12 ± 36.01 | 78.44 ± 58.22 |
P(EtOx-Ale) | 2.25 ± 0.21 | 44.29 ± 6.94 | 42.24 ± 11.36 | 25.85 ± 12.56 | 20.02 ± 16.09 | 75.44 ± 49.80 |
P(EtOx-OH-Ale) | 3.86 ± 0.72 | 76.22 ± 14.65 | 35.88 ± 12.60 | 19.66 ± 12.27 | 31.27 ± 24.20 | 79.23 ± 49.07 |
Finally, to mimic the clinical application scenario as closely as possible, the adhesion of the various membranes was also evaluated under wet conditions in PBS for 24 and 72 h (Fig. 5A and B). Evidently, only membranes comprising alendronate groups remained adhesive to bone after 24 h immersion in PBS, which weakened with increasing immersion time. Moreover, these membranes revealed much lower degrees of swelling compared to alendronate-free groups (Fig. S9, ESI†). To investigate the biodegradation effect on the adhesion under et conditions, the adhesion of the membranes to bone was also evaluated in PBS containing collagenase for 24 and 72 h. As expected, the addition of collagenase to the immersion medium reduced the adhesion of membranes comprising both POx-Ale polymers, indicating that degradation of the fibrous gelatin layer of the membranes was accelerated by the presence of collagenase.
Fig. 5 Membrane adhesion to bone under wet conditions in (A) PBS for 24 h, (B) PBS for 72 h, (C) PBS with collagenase for 24 h, and (D) PBS with collagenase for 72 h. |
Fig. 6 Weight loss (%) of membranes after immersion in (A) PBS and (B) PBS with collagenase. Values represent the mean ± standard deviation (n = 3). |
To obtain more insight into the biodegradation profile of the membranes, scanning electron microscopy was used to study the surface morphology of the membranes after 14 days of immersion in PBS with and without collagenase. As shown in Fig. 7, this SEM evaluation confirmed that blank membranes consisting of a gelatin carrier and polyester backing retained their fibrous layered structure after 14 days of soaking in PBS buffer. In PBS with collagenase, these blanks were severely degraded, as indicated by abundant formation of pores and disappearance of the fibrous surface morphology. On the contrary, the surface morphology of the polyester backing was not affected upon immersion in PBS. POx-coated membranes lost their fibrous surface morphology, revealing a smoother surface with the presence of pores after immersion in PBS with or without collagenase. With collagenase present in the solution, the pores were much larger (50–67 μm), confirming that the degradation rate of the membranes was accelerated due to enzymatic degradation of the carrier.
Fig. 7 Scanning electron micrographs of the membranes after 14 days of immersion in PBS solutions with and without collagenase. Scale bars correspond to 50 μm. |
POx-polymers were homogenously dispersed throughout the fibrous gelatin carrier using dry powder deposition, yielding a loading efficacy of approximately 80%. Although polymers were homogeneously distributed throughout the carriers, the membranes showed a variability in thickness and density. However, this variability was comparable to values reported for other commercially available membranes and can be attributed to their animal-derived origin, which typically comes along with batch-to-batch variability. The polyester backing layer was attached at a temperature well above the glass transition of all polymers used, resulting in increased mobility of the polymer chains and decreased viscosity of the backing layer,41,42 thereby allowing for diffusion of P(LA-CL) and PLGA into the fibrous architecture. Consequently, abundant interconnections were formed between POx polymers, gelatin fibers and the polyester backing. Moreover, the membranes became thinner upon attachment of the polyester backing bottom layer by heating while compressing. The non-porous morphology of the backing is regarded as beneficial to minimize premature cell and tissue ingrowth into bone defects covered by these membranes, thereby providing space for osseous regeneration.
The wettability of the membranes was determined by water contact angle measurements, since we assumed that hydrophilic membranes would promote the formation of molecular interactions between membranes and substrates.43 Upon initial contact, water was completely absorbed by the most hydrophilic membranes comprising P(EtOx) and P(EtOx-Ale) polymers, since these polymers are rich in polar side groups. However, after 6 seconds of contact, all membranes absorbed the droplet due to the porous gelatin membrane.
Subsequently, we performed tensile tests in order to study the mechanical properties of the membranes. From the stress–strain curve, we observed that the carrier showed a higher strain %, which corresponded to ductile behavior. During testing, fibers were continuously pulled from the fibrous layers without any detectable failure event. Both the tensile strength and tensile modulus increased significantly when a polyester backing layer was attached to the fibrous gelatin carrier. As such, the polyester backing may act as both cell-occlusive and reinforcing layer. The incorporation of moderate amounts of LA into P(LA-CL) blends significantly increased their tensile strength and rigidity, which improved their overall mechanical performance.44,45 The impregnation of POx polymers into membranes comprising a backing layer did not lead to further reinforcement. However, membranes impregnated with hydroxyl-functionalized P(EtOx-OH-Ale) polymers showed superior tensile strength and modulus values compared to other experimental groups. This could be attributed to the interaction between the hydroxyl side groups with amines and carboxylic acid groups present in the carrier as well as amines derived from P(PropOx-NH2) present in the polyester backing layer.
The adhesion of the membranes onto mineralized substrates (i.e. CaP coated Ti disks and bone specimens) was evaluated by in vitro lap-shear tests, which are generally accepted as test methods to measure adhesion to bone. Generally, all POx-impregnated membranes showed excellent adhesion to both apatite-coated Ti substrates and bone, as evidenced by high shear strength and work of adhesion values. As expected, cohesive failure was confirmed for all POx-impregnated membranes by the presence of membrane remnants on the substrates after lap-shear testing due to the layered architecture of the membranes. The cohesion of the layers might be improved by tuning the hot melt extrusion formulation of the backing. All POx-impregnated membranes adhered more tightly to CaP-coated disks compared to bone substrates, which was attributed to the surface roughness of CaP-coated disks (Ra ≈ 1.5 μm), which was higher than the smooth bone samples (Ra ≈ 0.8 μm). The adhesion strength and work of adhesion of membranes comprising POx polymers to mineral-free control substrates (i.e. Ti and demineralized bone specimens) were significantly lower, which confirms that alendronate, hydroxyl, and amides groups present in POx polymers formed strong and specific bonds with calcium-containing substrates. In contrast, POx-free membranes as well as Bio-Gide® controls did neither adhere to mineralized (CaP-coated Ti and bone) nor mineral-free control groups (Ti and demineralized bone), since these membranes lack functional groups with bone-bonding ability. Membranes impregnated with alendronate-free P(EtOx) and P(EtOx-Ale) showed higher shear strength to CaP-coated Ti than membranes impregnated with P(EtOx-OH-Ale) polymers due to their higher hydrophilicity (lower water contact angles) causing better contact and more molecular interactions between membranes and substrate. We attributed the higher work of adhesion values of P(EtOx)-impregnated membranes than with other POx groups to the hydrophilic nature of P(EtOx).46 Importantly, all POx-impregnated membranes showed similar shear strength and work of adhesion values. Nevertheless, the standard deviations are large which compromise the statistical analysis. To test the adhesion of the membranes under clinically relevant wet conditions, we submerged the membranes adhered to bone substrate into a PBS bath with and without collagenase for 24 and 72 h. Strikingly, we observed that all P(EtOx-Ale)-impregnated membranes remained strongly adhesive to bone after 24 h, whereas two out of three P(EtOx-OH-Ale)-impregnated membranes remained adhesive to bone. We attributed this difference under wet conditions to the higher alendronate content and reduced swelling of membranes impregnated with P(EtOx-Ale). However, after 72 h, only one out of three P(EtOx-Ale) and P(EtOx-OH-Ale)-impregnated membranes remained adhesive to bone. We attribute the reduced bone adhesion after 72 h to dissolution of POx-Ale polymers caused by their high hydrophilicity. As expected, in the presence of collagenase, the adhesion of both hydroxyl and hydroxyl-free POx-Ale membranes weakened, indicating that the degradation of the membranes was accelerated in the presence of collagenase. Although it should be realized that the selected collagenase concentrations 20 μg mL−1 were relatively high, these results confirm that bone adhesiveness of our membranes decreases with enhanced biodegradation. In contrast, all blank membranes and membranes impregnated with P(EtOx) detached from bone in PBS solutions with and without collagenase, proving that long-term adhesion under wet conditions could only be achieved by formation of coordination bonds between alendronate-functionalized POx polymers and calcium in bone. In addition, the high swelling degree of P(EtOx) also weakened the interfacial adhesion, rendering this polymer unsuitable for applications requiring long-term adhesion to bone under wet conditions.
Finally, the in vitro degradation of the membranes was studied by monitoring their weight loss upon immersion in phosphate-buffered saline solutions with and without collagenase. Ideally, to ensure mechanical stability, the degradation rate of the barrier membranes should match with the rate of bone healing. Hence, the integrity and stability of membranes should be guaranteed in the range from four weeks, when bone remodeling phase starts, up to several months.17 In our study, carriers gradually disintegrated in PBS due to their high swelling ability, while they partially dissolved in PBS solutions containing collagenase, suggesting that the degradation mechanism of the membranes was mainly dependent on enzymatic degradation of the fibrous gelatin carrier. The backing layer revealed an initial weight loss of about 15%, followed by a stable phase without any substantial degradation. After 14 days, 21 and 23 wt% of the original backing layers were degraded in PBS solutions with and without collagenase, respectively, which was caused by hydrolytic cleavage of the esters bonds and dissolution of uncrosslinked P(PropOx-NH2). The backing is mainly composed of P(LA-CL), which is a slow-degrading polymer with typical in vivo degradation times of more than two years.47,48 However, PLGA was added to the backing to accelerate degradation rate of the resulting membranes by increasing the amount of hydrophilic glycolic acid monomers in the copolymer.28 This accelerating effect of glycolic acid on polyester degradation was previously reported for PLGA with different L:G ratios of 50:50 (complete degradation after 2–3 months) vs. 85:15 (6 months).49,50 Consequently, variation of parameters such as copolymer ratio and the polyester percentage in the backing will allow to tune both degradability and mechanical strength of the resulting membranes tailoring to the specific requirements of the application. Blank membranes lost ∼20% more weight after 14 days in PBS with collagenase than in collagenase-free PBS, which confirmed that the gelatin carrier degraded proteolytically. The backing layer did not only reinforce the strength of the membranes in dry state, but also it helped to keep the integrity of the membranes even after 14 days immersed in PBS solution, avoiding the delamination of the carrier. Based on the carrier/backing weight ratio and the total weight loss of the backing, we estimated that 10 wt% of gelatin remained present in the blank membranes upon 14 days of soaking in collagenase-free PBS, whereas in PBS containing collagenase, it degraded almost completely. Strikingly, only blank membranes retained their fibrous morphology after 14 days of soaking in PBS, which corresponded to the lower amounts of weight loss of blank membranes vs. POx-impregnated membranes. We attribute the accelerated degradation of gelatin in POx-impregnated membranes to a faster dissolution of POx-gelatin complexes in PBS due to their higher hydrophilicity. Overall, it can be concluded that sequential degradation of gelatin and POx polymers on short-term followed by long-term degradation of the polyester backing provides the opportunity to independently tune the duration of the initial bone-adhesive phase vs. the long-term occlusive phase where the membrane should act as a barrier against soft tissue infiltration. We speculate that sequential instead of simultaneous degradation of membrane components will allow to reduce toxicity and inflammatory responses due to sudden release of high concentrations of degradation by-products.17
Footnote |
† Electronic supplementary information (ESI) available: Synthetic route of alendronate-functionalized POx polymers (Scheme 1); experimental procedures and characterization of the synthesized polymers (S1); NMR spectra of synthesized polymers (Fig. S1); images of a blank membrane and backing layer (Fig. S2); thin-film X-ray diffractograms and FTIR spectra of CaP-coated titanium disks, bone, and demineralized bone specimens (Fig. S3); scanning electron micrographs of CaP-coated titanium disks and longitudinal and cross-sectional views of bone specimens before and after demineralization (Fig. S4); experimental setups for tensile and lap-shear testing (Fig. S5); thermograms of POx-polymers and FTIR spectra of various experimental groups (Fig. S6); water contact angle measurements of the membranes (Fig. S7); FTIR spectra of bone specimens after lap-shear testing of various membranes (Fig. S8); swelling of membranes under wet conditions (Fig. S9); and physical characterization of the membranes (Table S1). See DOI: 10.1039/d1tb00502b |
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