Papia Haque

Andrew J Parsons

Ian A Barker

Ifty Ahmed

Derek J Irvine

Gavin S Walker

Christopher D Rudd

Interfacial Properties Of Phosphate Glass Fibres/Pla Composites: Effect Of The End Functionalities Of Oligomeric Pla Coupling Agents

Keywords: Polymer-matrix composites (A), Glass fibres (A), Coupling agents (A), Mechanical properties (B), Polylactic acid

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Introduction

Bioabsorbable phosphate glass fibres/polymer composites are being investigated as alternatives to commonly used metal alloy-based bone fracture fixation devices [START_REF] Navarro | Development of a biodegradable composite scaffold for bone tissue engineering: physicochemical, topographical, mechanical, degradation, and biological properties[END_REF].

According to Knowles [START_REF] Knowles | Phosphate based glasses for biomedical applications[END_REF], these 'Third Generation' implants remain in the body for the short term necessary for mechanical stabilisation of tissue. In addition the degradation products of the bioabsorbable implants perform an active function in the tissue regeneration process. A bioabsorbable fracture plate could potentially transfer load to the healing bone gradually over time.

Among bioabsorbable polymers, polyhydroxy acids and especially poly(lactic acid) (PLA) are well known for a number of biomedical applications; sutures, pins, screws, drug delivery systems and scaffolds for use in tissue engineering applications [START_REF] Navarro | In vitro degradation behavior of a novel bioresorbable composite material based on PLA and a soluble CaP glass[END_REF][START_REF] Mano | Bioinert, biodegradable and injectable polymeric matrix composites for hard tissue replacement: state of the art and recent developments[END_REF]. Hydrolytic degradation, or complete mass loss of PLA takes approximately 12 to 16 months [START_REF] Maurus | Bioabsorbable implant material review[END_REF]. PLA degrades to form lactic acid which is normally present in the body [START_REF] Gunatillake | Biodegradable synthetic polymers for tissue engineering[END_REF].

The reported elastic modulus of PLA ranged from 0.6 to 4 GPa which is higher compared to the familiar bioabsorbable polyhydroxy acids such as polycaprolactone, polydioxanone [START_REF] Daniels | Mechanical properties of biodegradable polymers and composites proposed for internal fixation of bone[END_REF]. However, the elastic modulus of cortical bone ranges from 6 to 20 GPa [START_REF] Daniels | Mechanical properties of biodegradable polymers and composites proposed for internal fixation of bone[END_REF], which is much higher than that of the polymers. Therefore, reinforcement of the polymer has been considered to attain sufficient mechanical properties for fracture fixation [START_REF] Navarro | In vitro degradation behavior of a novel bioresorbable composite material based on PLA and a soluble CaP glass[END_REF].

Phosphate based glasses (PBG) have great potential to be used as reinforcement in bioabsorbable composites for bone implants. PBG show a complete and congruent degradation in aqueous media with zero order release of incorporated oxides [START_REF] Neel | A facile synthesis route to prepare microtubes from phosphate glass fibres[END_REF]. The chemical composition of PBG can be designed to closely match that of the inorganic phase of bone [START_REF] Ahmed | Weight loss, ion release and initial mechanical properties of a binary calcium phosphate glass fibre/PCL composite[END_REF]. Moreover, the composition of PBG can be easily tailored giving variation in mechanical, thermal and biological properties to suit the end application.

The major limitation of PBG fibre reinforced polymer composites encountered thus far is rapid strength loss after exposure to an aqueous physiological environment [START_REF] Rinehart | Real-time dissolution measurement of sized and unsized calcium phosphate glass fibers[END_REF].

Recent research on glass fibres has involved surface treatments and coupling agents to try and enhance the hydrolytic stability of the fibre-polymer interface [START_REF] Jiang | Preparation of poly(ε-caprolactone)/continuous bioglass fibre composite using monomer transfer moulding for bone implant[END_REF][START_REF] Khan | Surface treatment of phosphate glass fibres using 2-hydroxyethyl methacrylate:fabrication of poly(caprolactone)-based composites[END_REF].

Apart from silane, modified oligomers/macromers of the matrix materials have currently been investigated to improve the interfacial bonding between fibre and matrix.

Phosphate glass fibres (PGF) coated with a mixture of modified methacrylate macromer and 2-hydroxyethyl methacrylate (HEMA), in the same macromer matrix gave the composite a bending strength of around 115 MPa [START_REF] Brauer | Degradable phosphate glass fiber reinforced polymer matrices: mechanical properties and cell response[END_REF]. However, the fibres used were 125 µm in diameter, which would have made them weaker and less flexible due to their large size, and therefore not appropriate as reinforcing fibres. Polyacrylic acid (PAA) and glycidyl acrylate modified PAA coupling agents have been applied to E-glass fibres and electron spectroscopy for chemical analysis (ESCA ) studies showed broad and shifted C 1s peak at 288.5 eV (O-C=O) suggesting adhesion of polymers to the fibre surface [START_REF] Lassas | Modification of E-glass fibre surfaces with polyacrylic acid polymers from an aqueous solution[END_REF].

The recent researches mentioned above suggest that some polymeric materials can make chemical bonds with some glass fibres. In this study, the matrix material was PLA and this is why the present study was conducted in order to investigate the effects of different PLA oligomer coupling agents on PGF/PLA composites.

Experimental

Phosphate glass and fibre production

PGF were produced by melt-draw spinning of phosphate glass (composition 40P 2 O 5 -24MgO-16CaO-16Na 2 O-4Fe 2 O 3 in mol%) using a dedicated in-house fibre manufacturing facility and then annealed at 474ºC [START_REF] Niazi | Mechanical and degradation properties of phosphate based glass fibre/PLA composites with different fibre treatment regimes. 8th International conference on durability analysis of composite systems[END_REF]. The fibres were kept in a desiccator prior to use. These annealed and untreated fibres were used as control for the study. The mechanical properties of PGF produced fluctuate slightly from batch to batch due to their variation in diameter and severity of surface flaws. Due to this slight variation, the properties of control and sized fibres were compared using fibres produced from within the same batch. Two batches of the control fibres produced were named as control-1 and control-2.

PLA oligomers synthesis

Acid ended PLA (PLA-acid) and its sodium salt (PLA-Na)

The PLA-acid and PLA-Na both were synthesised by ring opening polymerisation of D, L lactide (Sigma Aldrich, U.K.) and the resulting products were used without further purification. The PLA-acid and PLA-Na were synthesised using a dried lactide and lactic acid ratio of 5g: 0.105g and a dried lactide and lactic acid sodium salt ratio of 5 g: 0.112 g respectively. The mixtures were then heated to 130 o C and stirred for 48 hours and heated to 140 o C and stirred for 24 hours under a nitrogen atmosphere for PLA-acid and PLA-Na respectively. After this time period the mixtures were decanted whilst hot onto aluminium foil and allowed to cool to room temperature. The molecular weight and the polydispersity of the PLA-acid and PLA-Na were obtained 1921 (Mn) and 1.77 (PDI) and 4410 (Mn) and 2.33 (PDI) respectively.

Hydroxyl groups ended PLA

The PLA oligomers with different number of hydroxyl terminals were produced by depolymerisation of Natureworks PLA (PLA-grade 3051D, Natureworks LLC, U.S.A) catalysed by stannous octoate (Sigma Aldrich, U.K.) using different alcohols as initiators. Ethylene glycol, glycerol and sorbitol (Sigma Aldrich, U.K.) ended short chain PLA oligomers were synthesised by degrading dried Natureworks PLA (10 g) using ethylene glycol (0.72g), glycerol (1.068g) and sorbitol (2.108g) and named as EG-PLA, G-PLA and S-PLA respectively. Tin catalyst was used 0.05g for each of the mixtures. The weighed amount of PLA, catalyst and initiator of each batch were added to a 100 mL round bottomed flask equipped with a magnetic stirrer. The mixtures were heated to 180-185 o C for 2 hours until melt and then the reactions were continued at 170°C for 14 hours under a nitrogen atmosphere and with constant stirring. After this time period the mixtures were again heated to 180 °C and stirred for 2 hours. The mixtures were then decanted like PLA-acid and PLA-Na. A further short chain sorbitol ended PLA (S-PLA-1) of lower molecular weight compared to S-PLA was also synthesised in order to investigate the effect of molecular weight on the mechanical properties of sized fibres. The S-PLA-1 was produced in the same way as mentioned above for S-PLA, with the only exception is that the reaction was continued at 170 °C for 24 hours. The resulting oligomers were then dissolved in dichloromethane (Sigma Aldrich, U.K.) and precipitated in methanol (Sigma Aldrich, U.K.) to purify them. The precipitates were then dried under vacuum. The molecular weight and the polydispersity of the EG-PLA, G-PLA, S-PLA and S-PLA-1 obtained were 8210 (Mn) and 1.98 (PDI), 1217 (Mn) and 2.71 (PDI), 2881 (Mn) and 3.2 (PDI) and 115 (Mn) and 3.25 (PDI) respectively.

Gel permeation chromatography (GPC)

The number average molecular weight (Mn) and polydispersity index (PDI) of the oligomers were determined using GPC. GPC was performed utilising a refractive index (RI) detector with HPLC grade tetrahydrofuran (Fisher Scientific, U.K.) as the eluent. Analysis was performed at 40°C with a flow rate of 1 mL/min through two PLgel Mixed-C columns with a calibration range of 2000 -5000 Da calibrated with poly(styrene) narrow standards. All GPC equipment and standards were supplied by Polymer Laboratories (Varian Inc). GPC data was analysed using the Cirrus GPC Offline software package.

Coupling/sizing agent application processes

The fibres were sized with the above PLA oligomers (6 oligomers), 3aminopropyl triethoxy silane (silane) and 2-carboxyethyl phosphonic acid (phosphonic acid) (Sigma Aldrich, U.K.). Silane is a common sizing agent used for various glass fibres and this is why silane has been investigated as one of the sizing agents in this study. Phosphonic acid contains a phosphonic group which has similarities in chemical structure to the PGF and it has been investigated whether this similarity in structure can facilitate the bonding between this acid and the PGF. The interfacial properties of silane and phosphonic acid sized fibres were compared with those of the PLA oligomers sized fibres. Using a dip-coating process, the agents were applied to the control fibres with a fibre:solution ratio of 1.5g:100ml and a sizing agent:solvent ratio of 0.0043moles:100ml [START_REF] Niazi | Mechanical and degradation properties of phosphate based glass fibre/PLA composites with different fibre treatment regimes. 8th International conference on durability analysis of composite systems[END_REF]. This was followed by heat curing and finally rinsing with the sizing solvents to remove any residual artefacts. The solvents used in the sizing processes were (a) ethanol (Sigma Aldrich, U.K.) /deionised water (90:10 v/v ratio) for silane and phosphonic acid, (b) chloroform (Sigma Aldrich, U.K.) for PLA-Na and PLA-acid, and (c) dichloromethane for EG-PLA, G-PLA, S-PLA and S-PLA-1. Fig. 1 shows the chemical structures of all the sizing agents applied.

Single fibre tensile test (SFTT)

The tensile properties of the fibres produced were obtained by single filament testing using a Series S Hounsfield machine, following the international standard BS ISO 11566 [START_REF]Carbon fibre -Determination of the tensile properties of single-filament specimens[END_REF]. 20-30 samples were prepared at a gauge length of 25mm for each treatment type. The crosshead speed of the machine was 1mm/min and the load cell capacity was 5N. PGF are essentially brittle and Weibull distribution is an accepted statistical tool used to characterise the failure mode of brittle fibres. In this study, Weibull parameters were obtained from the tensile strength data calculated using Minitab® 15 (version 3.2.1).

Interfacial shear strength (IFSS) studies

The IFSS was measured from single fibre composites (SFC) by using the single fibre fragmentation method. SFC specimens were produced by placing an axially aligned single fibre between two rectangular pieces of PLA film [START_REF] Niazi | Mechanical and degradation properties of phosphate based glass fibre/PLA composites with different fibre treatment regimes. 8th International conference on durability analysis of composite systems[END_REF] of 80 x 20 x 0.4 mm 3 and hot pressing at 210ºC with a 2.5 kg weight for 1 minute. The resulting specimens were cooled to room temperature and finally cut into dog-bone shapes. These were axially loaded in a tensile testing machine (Hounsfield series S testing, U.K.) with a load cell of 1kN and crosshead speed of 1mm/min. All IFSS values were obtained from an average of 5-10 repeat specimens. After having conducted the tensile tests, the specimens were placed under an optical microscope (Nikon Optiphot, Japan) and the numbers of fibre fragments present in the 25mm gauge length were tallied, in order to calculate the IFSS.

These IFSS values obtained were calculated using the Kelly-Tyson equation below [START_REF] Kelly | Tensile properties of fibre-reinforced metals: Copper/Tungsten and Copper/Molybdenum[END_REF];

τ i = d.σ f /2L c (Equation 1)
where τ i is the IFSS, d is the fibre diameter, σ f is the single fibre tensile strength at the critical fragment length L c .

σ f = σ o (L c /L 0 ) -1/m ( Equation 2)
L c = 4/3(L f ) (Equation 3) 4)

L f = L 0 /N (Equation
where σ o and m are the Weibull scale and shape parameter respectively, for the fibre strength at gauge length L 0 , L f is the average fragment length and N is the number of fibre fragments obtained from the SFC tests [START_REF] Park | Interfacial properties and microfailure degradation mechanisms of bioabsorbable fibres/ poly-L-lactide composites using micromechanical test and non-destructive acoustic emission[END_REF].

Thermogravimetric (TG) analysis

TG analysis was performed to quantify the amount of sizing agents present on the fibre surfaces. The analyses were conducted from room temperature to 450ºC using a SDQ600 Thermogravimetric analyzer from TA instruments (U.S.A.) with a heating ramp rate of 20ºC /min under an air flow of 100ml/min.

X-ray photoelectron spectroscopy (XPS) analysis

XPS analysis investigated the effectiveness of compatiblising / sizing agent application to the fibre surfaces. A Kratos AXIS ULTRA with a mono-chromated Al kα X-ray source (1486.6eV) was operated at 15mA emission current and 10kV anode potential. For the XPS measurement, a small batch of fibres was mounted on a hollow steel bar using a high vacuum adhesive tape at two ends. Drift of the electron binding energy of the peaks due to a surface charging effect was calibrated by utilising the C 1s peak of the contamination (285eV).

Scanning electron microscopy (SEM)

Samples of each of the fibres were sputter coated with carbon and micrographs were taken using a JEOL 6400 SEM at an accelerating voltage of 20 kV using SE (Secondary electron) modes.

Statistical analysis

A t-test was conducted to identify whether there were statistically significant differences between the means of the two sets of data [START_REF] Montgomery | Applied statistics and probability for engineers[END_REF]. The statistical analysis of the uncorrected and corrected IFSS values was performed using a paired t-test. The statistical analysis of the tensile strength and the IFSS values of the fibres and the SFC respectively was conducted using a two sample unpaired t-test. To interpret the t-test results, a two tailed P-value (P) was considered with a 95% confidence interval and the value of significance level 0.05. The difference was accepted as significant when P obtained was less than 0.05.

Results and discussions

Effect of carboxylic acid and sodium ended PLA oligomer coupling agents

The IFSS values obtained for the PLA-acid, PLA-Na, silane and phosphonic acid sized and the control-1 fibres are presented in Fig. 2. All the values were calculated using a 2-parameter Weibull model for the control-1 fibres. The Weibull scale and shape parameters of the control-1 fibres obtained were 340.0 and 6.3 respectively (Fig. 3a). The IFSS obtained for the control-1 fibres was 8.9 ± 2.9 MPa. The highest IFSS value of ± 2.3 MPa was obtained for PLA-Na sized specimens -acid

- ± 2.3 ± ± 2.2
It is observed from Fig. 2 that the IFSS values of PLA-acid and PLA-Na sized fibres were higher than that of the control-1. The low molecular weight PLA-acid and PLA-Na coupling agents showed relatively high IFSS values (i.e. improved bonding), which was suggested to be due to similarities in the chemical structures of the agents and matrix. The PLA-Na showed better bonding between the fibre surface and matrix as compared to the PLA-acid. The processing used and the relatively high level of hydrolytic instability of PLA-acid may result in it decomposing during preparation of the test specimens. However, the IFSS values of the silane and phosphonic acid treated fibres obtained were similar to that of the control-1. A similar result was also obtained by Cozien-Cazuc et al. for both unsized and silane sized P 40 Na 20 Ca 16 Mg 24 PBG fibres within a PCL matrix [START_REF] Cozien-Cazuc | Characterisation of resorbable phosphate glass fibres[END_REF].

The strength profile in relation to the SFTT of the PLA-Na sized fibres was investigated in order to verify if their Weibull parameters were different to that of control-1. The Weibull scale and shape parameters obtained for PLA-Na sized fibres were 419.5 and 5.8 respectively (Fig. 3b), which were different from those of control-1 (Fig. 3a). The IFSS value of PLA-Na sized fibres w therefore corrected using the Weibull parameters of the PLA-Na sized fibres in order to obtain more accurate results.

After correction, the IFSS value obtained was 1 MPa. It was observed that the change in IFSS with and without correction of the Weibull parameters was statistically significant (P < 0.0001) for the PLA-Na sized fibres. 6.2, 77.0 ± 5.8 GPa respectively. However, the difference between the tensile strength and modulus of all the sized fibres was found to be statistically insignificant in comparison to control-2.

Effect of different hydroxyl groups ended PLA oligomer coupling agents

Fig. 5 shows the IFSS values of control-2 and EG-PLA, G-PLA and S-PLA sized fibres within a PLA matrix, along with the Mn of the oligomers. The effects of the end-group choice of these low molecular weight oligomer coupling agents were investigated from their IFSS values. These IFSS values were calculated using the individual Weibull parameters of control-2 and each of the sized fibres. The IFSS of control-2 and EG-PLA, G-PLA and S-PLA sized fibres obtained were 14.3 ± 6.1, 8.9 ± 2.6, 16.5 ± 3.3 and 23.2 ± 10.3 MPa respectively. The IFSS of EG-PLA and S-PLA (P values are 0.019 and 0.017 respectively) were significantly different from that of control-2. The highest IFSS value was obtained from S-PLA sized fibres followed by G-PLA and EG-PLA. The polymer chain end of S-PLA contains 5 hydroxyl groups which is higher than that for G-PLA (2 hydroxyl groups) and EG-PLA (1 hydroxyl group). The interfacial reactions were expected to be either between the hydroxyl groups on the PGF surface and the coupling agents by esterification or hydrogen bonding as shown in Fig. 6, and/or formation of salt complexes with the metal cations present in the glass. Regarding the interfacial reactions, the number of hydroxyl groups in the oligomers is assumed to influence the IFSS values.

Comparing Fig. 4 and 5, it is observed that the tensile properties of the PLA oligomers sized fibres increased with the increase of Mn of the oligomers, though the differences in their tensile properties are statistically insignificant. However, no trend was observed in the level of interfacial adhesion with a difference in their Mn. The increase in chain length of the agents is suggested to contribute to an increase in coverage of the flawed surface of the fibres. However, the functional end groups of these oligomer chains play a vital role in their ability to effectively bond to the PGF surface. The differences in the end groups of the oligomers perhaps negate any effect Mn may have on the IFSS values of the sized fibres.

Effect of different Mn of the PLA oligomers

To understand the effect of Mn of a particular oligomer on the mechanical and interfacial properties of the sized fibres, two oligomers having the same functional end group with variation in chain length were investigated. Another sorbitol ended PLA; S-PLA-1 (Mn~115 Da) was introduced to make this comparison. In Fig. 7, the mechanical properties and the IFSS value of these sized fibres were compared with that of the S-PLA (Mn ~ 2881 Da). The tensile strength and modulus of S-PLA-1 sized fibres obtained were 279.4 ± 73.9 MPa and 73.8 ± 8.0 GPa respectively, which were lower than that of S-PLA sized fibres. However, taking statistical variation into account, no difference was observed in tensile properties of S-PLA and S-PLA-1 sized fibres. The IFSS value of 9 ± 3.5 MPa was obtained for S-PLA-1 sized fibres. The IFSS values obtained for S-PLA and S-PLA-1 sized fibres were significantly different (P = 0.0002).

With the functional end group for S-PLA and S-PLA-1 being the same, the affinity toward the PGF surface was expected to be similar for both these oligomers.

The matrix-philic tails of the oligomers being different in length may have shown variation in the level of affinity with the PLA matrix. S-PLA-1 having very low Mn, the chain length of the agent was not sufficient to afford significant number of interactions and/or entanglement with the matrix for achieving good interfacial adhesion. Therefore, the longer chain length (i.e. higher Mn) was seen to confer a higher shear strength value of around 23 MPa for the sized fibres.

TG analysis

TG analysis of the S-PLA sized and control-2 fibres was conducted to ascertain the mass of sizing agent that was present on the fibres by pyrolysing it from the surface.

Three replicates of each type of fibres were conducted to make an average of their result.

The percentage weight losses recorded were 2.4 ± 0.5 and 0.23 ± 0.02 for S-PLA sized and control-2 fibres respectively. The TG result for S-PLA sized fibres suggested that the amount of sizing agent attached to the fibres was quite low, being approximately 0.1% of the sizing agent present in the sizing solution. Currently the actual coating level achieved is thought to be due to the low concentration (0.043M) of the sizing solution used in the study. Therefore, experimentation is now underway to optimise the concentration of the sizing solutions used in this process. In the study reported here, the sizing solutions used were all fixed at the above concentration in order to obtain and evaluate comparable data.

XPS analysis

High resolution XPS analysis was performed on S-PLA sized and control-2 fibres to examine the existence of sizing agents on the fibres.

The high resolution scan of the 1s core of O (oxygen) for control-2 and S-PLA sized fibres are shown in Fig. 8 labelled a and b respectively. In narrow scan 8a, the oxygen peak was deconvoluted into two peaks, which were seen at 531.5 and 533.3eV binding energy for non-bridging (NBO) and bridging (BO) phosphate oxygen respectively [START_REF] Gresch | X-ray photoelectron spectroscopy of sodium phosphate glasses[END_REF]. The high resolution scan of the O1s core of the S-PLA sized fibres (scan 8 b) showed three deconvoluted peaks at 531.4, 532.3 and 533.9 eV binding energy. In this case, it is not possible to differentiate between the BO and NBO of the phosphate glass and the C-O-C and C=O bonds within the sizing agents [START_REF] Beamson | High resolution XPS of organic polymers, The Scienta ESCA 300 Database[END_REF], as both have absorption peaks over similar ranges of binding energy. The third peak at 532.3 eV has been attributed to the presence of C-O-H of the S-PLA.

SEM analysis

The SEM micrographs (Fig. 9a,b) show the fibre surface of the control-2 and S-PLA sized fibres. The control-2 fibre (Fig. 9a) shows a pristine surface suggesting no sizing agent on it. With regards to Fig. 9b some agent persists on the surface. These images also support the bonding of the sizing agents on the fibres surface well with the observation made above from their TG and XPS analyses.

Conclusion

The PLA oligomer coupling agents being similar in chemical structure to the matrix showed a better quality of interfacial bonding between PGF and the PLA matrix.

The functional end groups with a greater number of hydroxyl groups showed better compatibility (or attachment) toward PGF. With regards to the number of hydroxyl groups present in the oligomers, sorbitol ended PLA showed the highest IFSS. It was also observed that the molecular weight of an oligomer had a substantial effect on the interfacial properties of the sized PGF fibres. An optimum size of PLA chain with a

Fig. 4

 4 Fig.4shows the tensile strength and modulus of the different functionalized

Figure 1

 1 Figure 1 Chemical structures of the coupling agents.

Figure 2

 2 Figure 2 IFSS values of control-1 and silane, phosphonic acid, PLA-acid and PLA-Na

Figure 3

 3 Figure 3 Weibull analysis of (a) control-1 and (b) PLA-Na sized fibres were completed

Figure 4

 4 Figure 4 Tensile properties of control-2 and EG-PLA, G-PLA and S-PLA sized fibres.

Figure 5

 5 Figure 5 IFSS values of control-2 and EG-PLA, G-PLA and S-PLA sized fibres and the

Figure 6

 6 Figure 6 Possible chemical reaction mechanisms between PGF and S-PLA, (a)

Figure 7

 7 Figure 7 Effect of molecular weight of S-PLA and S-PLA-1 on the tensile and

Figure 8

 8 Figure 8 High resolution XPS spectra of O 1s core of (a) control-2 and (b) S-PLA sized

Figure 9

 9 Figure 9 SEM micrographs of (a) control-2 and (b) S-PLA sized fibres.

suitable functional end group is suggested to be a hugely beneficial coupling agent for the PGF/PLA composites investigated.


V Guarino

M Lewandowska

M Bil

B Polak

L Ambrosio

Morphology And Degradation Properties Of Pcl/Hyaff11 ® Composite Scaffolds With Multiscale Degradation Rate

The analysis of scaffold degradation is a promising strategy for understanding the dynamic changes in texture and pore morphology which accompany polymer resorption, and for collecting same fundamental indicators regarding the potential fate of the scaffold in the biological environment. In this study, we investigate the morphology and degradation properties of three composite scaffolds based on poly(caprolactone) (PCL) embedded with benzyl ester of hyaluronic acid (HYAFF11 ® ) phases, and, in turn, different reinforcement systems -i.e, calcium phosphate particles or continuous poly(lactic acid) (PLA) fibres. Scanning electron microscopy (SEM) and µ-tomography supported by digital image analysis enabled a not invasive investigation of the scaffold morphology, providing a quantitative assessment of porosity (which ranged from 63.1 to 82.8), pore sizes (which varied from 170.5 to 230.4 µm) and pore interconnectivity. Thermal analyses (DSC and TGA) and Raman spectroscopy demonstrated the multi-scale degradation of the composite with highly tailoring degradation kinetics depending on the component material phases and scaffold architecture changes, due to their conditioning in simulated in vivo environment (i.e, SBF solution). These results demonstrate that the judicious mixing of materials with faster (i.e, Hyaff11) and slower (i.e, PLA and PCL) degradation kinetics, different size and shape (i.e, domains, particles or long fibres), certainly concurs to design a smart ACCEPTED MANUSCRIPT 2 composite scaffold with time-controlled degradation which can support the regeneration of a large variety of tissues, from the cartilage to the bone.

Introduction

The architecture and the structural organization of scaffolds are critical factors to assure the functionality of tissue engineered constructs and their appropriate and effective application in health care. It is commonly recognized that porosity actively influences the biological behavior of seeded cells during in vitro culture, through the provision of an adequate surface area for cell attachment, stimulating the free circulation of biological fluids and oxygen transport for their maintenance and, subsequently, influencing the in vivo conductive response of the structure after its implantation [START_REF] Ciapetti | Human bone marrow stromal cells: In vitro expansion and differentiation for bone engineering[END_REF][2] [START_REF] Savarino | The performance of poly-e-caprolactone scaffolds in a rabbit femur model with and without autologous stromal cells and BMP4[END_REF]. In this context, the synchronization of polymer degradation with its replacement by natural tissue produced from cells would be desirable. As a consequence, the degradation properties of a scaffold are relevant both to the biomaterial design and to the long-term success of a tissue engineered construct [START_REF] Sung | Galis -The effect of scaffold degradation rate on three-dimensional cell growth and angiogenesis[END_REF].

To date, much effort has been dedicated to the design of a variety of composite materials for tissue engineering scaffolds with tailored degradation properties, based on biocompatible and bioerodible polymers [START_REF] Seal | Polymeric biomaterials for tissue and organ regeneration[END_REF] [START_REF] Kohn | Bioresorbable and bioerodible materials[END_REF]. Recently, several authors have focused their attention on polyester-based composite systems. [START_REF] Manjubala | Biomimetic mineral-organic composite scaffolds with controlled internal architecture[END_REF][8] [START_REF] Miller | Degradation rates of oral resorbable implants (polylactates and polyglycolates): rate modification with changes in PLA/PGA copolymer ratio[END_REF] The influence of several factors on degradation kinetics are considered [START_REF] Santavirta | Immune response to polyglycolic acid implants[END_REF], including molecular factors (e.g., chain orientation, molecular weight, M w , and polydispersity) supramolecular factors (crystallinity, the spatial distribution of chemically reactive filler [START_REF] Andrew | The influence of polymer blend composition on the degradation of polymer/hydroxyapatite biomaterials[END_REF]) and environmental factors (e.g., mechanical stimuli). The various influences of these factors potentially generates a wide range of resorbable properties for custom-made systems. However, the evolution of structural properties in connection with the selective degradation of one of the components of the scaffold matrix has not been adequately explored, despite its significant contribution to the activation and the requisite enabling and support of specific biological mechanisms on an appropriate time-scale.

In the current work, polymeric composite scaffolds based on poly( -caprolactone) (PCL) and esters of hyaluronic acid have been proposed because of the encouraging pre-clinical results which similar systems have exhibited as in vivo scaffolds for tissue regeneration [START_REF] Chiari | A tissue engineering approach to meniscus regeneration in a sheep model[END_REF] [START_REF] Kon | Tissue Engineering for Total Meniscal Substitution: Animal Study in Sheep Model[END_REF]. PCL shows a remarkably slow degradation rate -from 6 months up to 3 years for its complete removal from the host body -as a function of the starting molecular weight [START_REF] Gabelnick | Biodegradable implants: alternative approaches[END_REF] [START_REF] Gunatillake | Biodegradable synthetic polymers for tissue engineering[END_REF]. Firstly, this preserves mechanical integrity over the degradation lifetime of the device, so allowing adequate mechanical support during the post-implantation period. Secondly, the absence of toxic response, due to the reduced release of acidic products, assures enhanced bone ingrowth into the porous scaffolds supported by the progressive increase of pore size as the erosion mechanism advances [START_REF] Rezwan | Biodegradable and bioactive porous polymer/inorganic composite scaffolds for bone tissue engineering[END_REF]. The integration of highly degradable materials, obtained by chemical modification of purified hyaluronan (HA), namely HYAFF11 ® , formed by the partial or total esterification of the carboxyl groups of glucuronic acid with different types of alcohols -provides the opportunity to affect directly the cell activities, either favouring or, conversely, inhibiting the adhesion of certain cell types [START_REF] Campoccia | Semisynthetic resorbable materials from hyaluronan esterification[END_REF][18] [START_REF] Grigolo | Transplantation of chondrocytes seeded on a hyaluronan derivative (hyaff-11) into cartilage defects in rabbits[END_REF]. Here, partially esterified HYAFF11 ® with reduced hydrophilic, negatively charged, carboxyl groups along the polymer backbone, -esterification degree equal of 75% -has been selected for its optimal hydrophilic/hydrophobic character which confers an adequate biological recognition, without drastically penalizing the structural integrity of the composite scaffold at the preliminary stage of the culture [START_REF] Mori | Haluronan-based biomaterials in tissue engineering[END_REF]. Finally, either PLA fibres or bioactive ceramic particles were added to the polymer matrix in order to satisfy specific structural requirements of scaffolds for load bearing applications.

In this work, we focus on the evolution of morphology and degradation properties of proposed composites for underlining the contribution of single material phases and their peculiar chemical and physical properties on the dynamic changes of scaffold features within simulated in vivo conditions.

Materials and Methods

Materials and scaffold preparation

Poly -caprolactone pellets (MW 65 kDa, = 1.13 g/cm 3 Aldrich) were dissolved in a solvent mixture to form a solution by stirring for 2h at 40°C. Two different solvents were used, respectively, tetrahydrofurane (Sigma Aldrich, Italy) and dimethylsulfoxide (Sigma Aldrich, Italy) by 80/20 weight ratio. A benzyl ester of hyaluronic acid with 75% of esterification degree (HYAFF11 ® -75p -= 1.387 g/cm 3 [START_REF] Netti | Thermodynamics of Water Sorption in Hyaluronic Acid and Its Derivatives in Polymeric drugs & drug delivery systems Editors[END_REF]) supplied by Fidia Advanced Biopolymers in powder form, was then mixed to the polymer solution imposing a PCL/HYAFF11 ® weight ratio equal to 80/20.

Depending on scaffold type, NaCl crystals (Fluka) with specific size ranges (212-300 µm or 300-500 µm) were employed as templating agent by using different PCL/NaCl volume ratio. Furthermore, a mixture of Ca 2 NaK(PO4) 2 (CSPP) and, monocalcium phosphate monohydrate (MCPM) supplied by the Technical University of Catalunia (UPC) under the name of R-cem or polylactid acid fibres (PLA -75/24 dtex) supplied by Sofradim have been used as reinforcing agent.

Three different PCL and HYAFF11 ® based composite scaffolds typologies were prepared via phase inversion and salt leaching technique, labelled respectively as PH, PHR and PHF. The PH type was developed as from the PCL/HYAFF11 ® solution described above. Sodium chloride (NaCl) crystals, in the 212-300 µm size range, were then incorporated as templating agent into the polymer solution by using a PCL/NaCl volume ratio of 9/91 v/v. The mixture was placed into Teflon moulds reproducing the final dish shape, 5mm as diameter, of the scaffold for in vitro degradation tests. Finally, ethanol washings for 24h at room temperature were performed to completely extract the used solvents whereas daily washing in bi-distilled water for 7 days were used to leach out NaCl crystals as well as to remove other contaminants.

Filler-loaded scaffolds (PHR) have been prepared as previously described, by further adding the R-cem ceramic powder to the PCL solution by using a PCL/R-cem weight ratio equal to 80/20. As a precautionary measure, R-cem powder and NaCl crystals were premixed before the introduction into the polymer solution to minimize the cluster formation.

Finally, fiber reinforced composite scaffolds (PHF) were obtained by combining the phase inversion/salt leaching and filament winding technology as described elsewhere [START_REF] Guarino | Polylactic acid fibre reinforced polycaprolactone scaffolds for bone tissue engineering[END_REF]. Briefly, PLA fibres impregnated through the PCL/HYAFF11 ® solution, were wound on a polypropylene tubing coated stainless steel mandrel with 1 mm of diameter by using a specific winding parameter set (winding angle WA=45°, winding pitch WP=500µm). In this case, a less viscous polymer solution by changing the PCL/solvents (12/88 wt/wt) and the PCL/NaCl ratios (32/68 v/v) was selected in order to optimize the fibre/matrix adhesion, maintaining the same PCL/HYAFF11 ratio.

Moreover, NaCl crystals with bigger sizes (300-500 µm) were used to minimize the effects of crystal fragmentation on the pore size, expected during the process. Once tubular shaped scaffold was obtained, conventional procedures involving ethanol and water washings, for 24 hours and 7 days, respectively, were employed. Cylindrical samples with 5mm as outer diameter and 10 mm as height were used for porosity investigations (n=5) while dish shaped samples, thickness of 3 mm, for degradation tests (n = 4).

Morphological investigation by SEM/FESEM

Scaffold morphology was investigated via Scanning Electron Microscopy using a scanning electron microscope (Stereoscan 440, Leica, UK). Briefly, specimens were fractured in liquid nitrogen using a razor blade along preferential directions, parallel and perpendicular to the surface, respectively. The resulting transverse and longitudinal sections were gold-coated under vacuum by using an automatic coating sputter set at 15 mA for about 20 minutes (Emiscope SC500, Italy). The porosity was assessed in terms of pore size, shape and spatial distribution by images at different magnifications.

Further investigations on the scaffold morphology have been performed with the support of field emission scanner electron microscopy (FE-SEM, FEI Quanta200, The Netherlands) able to recognize material phases with different chemical and physical properties by the detection of back scattered (BSE) electrons and secondary electrons (SE) at low acceleration voltage (4kV).

Quantitative µ-tomographic analysis

The samples were scanned using a SkyScan 1072 micro-CT desk scanner supported by micro-tomographic reconstruction software (Skyscan) [START_REF] Guarino | The synergic effect of polylactide fiber and calcium phosphate particles reinforcement in poly -caprolactone based composite scaffolds[END_REF] [START_REF] Darling | 3D Microtomographic Characterization of Precision Extruded Polycaprolactone Scaffolds[END_REF]. The reconstruction of the three-dimensional objects was achieved through a serial reconstruction of cross sections or 2-D slice maps [START_REF] Mori | Haluronan-based biomaterials in tissue engineering[END_REF]. The output format for each sample was 300 serial 1024 ×1024 bitmap images with a resolution of 10.14 µm. After the serial reconstruction, the stack of the 2-D slice maps was elaborated to construct a realistic view of the 3D-object with possibilities to "rotate" and "cut" the object model through dedicated software (SkyScan pack-Belgium) as illustrated by Sun et al. [START_REF] Sun | Recent development on computer-aided tissue engineering-a review[END_REF]. By the accurate adjustment of brightness and image contrast and the manipulation of the grayscale thresholds, a 3-D model of the positive (direct reconstruction) and negative spaces (reverse reconstruction) was obtained, on which to calculate basic morphological parameters such as porosity (i.e, porosity degree, pore size) and bulk structure (i.e, strut size and distribution). The detailed assessment of pore morphology in term of pore size and distribution was performed by CTan Skyscan software whereas the evaluation of strut parameters was performed by ANT Skyscan software. Moreover, further morphometric parameters were evaluated to complete the scaffold morphology characterization: the anisotropy degree (DA) and the Structure Model Index (SMI) respectively -which refer to preferential pore orientations, density and pore interconnection [START_REF] Hildebrand | Quantification of Bone Microarchitecture with the Structure Model Index[END_REF]. High DA values (generally over 1) indicated high anisotropy of the trabecular structure [START_REF] Ruegsegger | A microtomographic system for the nondestructive evaluation of bone architecture[END_REF] where low SMI values usually referring to poor interconnection degree of porous network [START_REF] Ohser | -Mesh free estimation of the structure model index[END_REF].

Finally, further analysis on individual 2-D slice permitted to distinguish single radiopaque phases through differences in contrast of local regions, offering an estimation of the spatial distribution of inclusions and reinforcing elements within the polymer matrix.

Degradation studies a) In vitro degradation assays

In vitro degradation studies were carried out in a simulated body fluid (SBF) at a controlled temperature of (37±1) o C. SBF was prepared by dissolving the reagents NaCl,

NaHCO 3 , KCl, K 2 HPO 4 •3H 2 O, MgCl 2 •6H 2 O, CaCl 2 •2H 2 O
and Na 2 SO 4 into distilled water, as described by Kokubo et al [29]. The solution was buffered to pH 7.25 at 37 o C with tris(hydroxy-methyl)aminomethane and hydrochloric acid. Samples were immersed in SBF solution for time periods of 1, 4 and 8 weeks. The ratio of SBF volume to polymer mass was 1:100. The medium was not changed over the whole testing period. After a given time, the samples were extracted from the solution, rinsed gently with distilled water and left to dry at 37 C to a stable mass.

b) Mass loss measurements

The mass loss during immersion in SBF was calculated as Mass loss [%] = [(m 0 -m 1) /m 0 ] × 100 where m 0 and m 1 are the masses of the samples before and after degradation, respectively. At least 4 samples per data point were measured.

c) DSC analysis

DSC measurements were performed with a TA Instrument model DSCQ1000 equipped with a liquid nitrogen cooling unit. The instrument was calibrated using indium and sapphire standards. All measurements were performed in a helium atmosphere using standard crimped aluminum pans. Samples, with average weights of 9.2-10.8 mg, were scanned at a heating rate of 10 o C min -1 . The glass transition temperature, Tg, was measured at the mid-point of the heat capacity inflexion point. The crystallization temperature, Tc, and melting temperature, Tm, were determined from the peak value of the respective endotherms and exotherms. The degree of crystallinity (Xc) was evaluated according to the following equation:

Xc % = ( H m / H o ) X (100/w)
Where w is the weight fraction of PCL or PLA in the composite, H m is the measured enthalpy of melting and H o is the enthalpy of melting of a 100% crystalline (PCL 139.5 J/g) [START_REF] Coombes | Precipitation casting of polycaprolactone for applications in tissue engineering and drug delivery[END_REF] and (PLA 93 J/g) [START_REF] Sarazin | Controlled preparation and properties of porous poly(L-lactide) obtained from a co-continuous blend of two biodegradable polymers[END_REF] d) TGA analysis Thermal gravimetric analysis (TGA) was performed on a TA Instrument Q 500 under nitrogen atmosphere, from room temperature up to 700 o C, with a heating rate of 10 o C/min.

e) Raman spectroscopical analysis

A visible Raman Microscope Spectrometer (Nicolet Almega XR; laser 532 nm; laser power level: 100%; exposure time 30 sec; spectrograph aperture: 50 m pinhole) was used to emboss changes in chemical structure of the scaffolds, after the degradation studies.

Results

Morphology

The morphology of the three composite scaffolds was first assessed by scanning electron microscopy (SEM) (Figs. 1 and2). In each case, high structural porosity, characterized by a bimodal pore size distribution, was observed. In the case of PH and PHR, cross-section images indicate a homogeneous spatial distribution of macropores of undefined shape and with sizes ranging from 150 to 500 micrometers (Fig. 1 A-B). In the case of PHF scaffolds (Fig. 1C), a more ordered structural organization is observed, with a well-defined arrangement of pores between adjacent fibres. PLA fibres are well integrated into the PCL porous matrix and appear to influence pore morphology. In all cases, micropores with sizes of ca 10 µm may be easily recognized, both in the bulk and on the scaffold surface (Fig 2A -B). Marked differences in micro-and macropore morphology are also apparent on comparing scaffolds with and without fibres. More specifically, a reduction in macroporosity as well as concomitant increase in microporosity have been revealed in the case of PHF, ascribable to the different amount of templating agent and polymer solution concentration used during the preparation. A more accurate investigation of scaffold morphology by FESEM allowed investigation of the distribution of different polymeric and ceramic phases into the scaffold matrix as a function of the chemical and physical properties. Topographic images of scaffolds (Fig. 3A -square) highlight the presence of the HYAFF11® hydrophilic phases (dark grey) embedded in the PCL hydrophobic matrix (light grey). The lower brightness of the Hyaff11 phases arises from the higher electron density of polymer due to the presence of benzylic groups, which cause higher scattering of secondary electrons compared to the PCL.

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Higher magnification images (Fig. 3B) show the distribution of R-cem calcium phosphate particles with micrometric and sub-micrometric sizes, and their tendency to form micrometric clusters into the polymer matrix.

A quantitative investigation of scaffold morphology was carried out using microCT for all scaffold types (Fig. 4). First, the stack of X-Ray images radiographs was rendered as Afterwards, the volume rendering (Fig. 4, left side) was carried out to assess basic porosity parameters in the three-dimensional space. Typically, by converting the slices in "black and white" images, a loss of information on the composite materials phases can occur. However, this was not limitative in the case of the quantitative pore investigation performed which has been referred only to the calculation of objects on macrometric scale (i..e, macropores, filler cluster, fibre bundle). In Table 1, a summary of porosity features for the different scaffolds is reported. Porosity degrees ranged from 63.1 to 82.8 %, pore sizes from 170.5 to 230.4 µm, the anisotropy degree from 0.1 to 0.226, moving from fiber reinforced scaffolds to fibreless one. According to the difference in scaffold porosity, the SMI values change ranging from 0.951 to 2.002.

Degradation studies a) In vitro mass loss

The changes in mass loss of the scaffolds as a function of degradation time are shown in Fig. 5. After 1 week of immersion in SBF, the mass loss was comparable for all three materials and less than 2%. After 4 weeks of immersion in SBF, a linear weight loss was observed for all the composites. The highest weight loss -15.5% -was observed for PHR whereas the lowest -10% -for PH scaffold. During the next 4 weeks, the degradation rate of PHR decreased, so reaching, after 8 weeks degradation, comparable mass losses at ca17% for all composites (Table 2).

b) Raman Spectroscopy

Fig. 6 illustrates Raman spectra of PH (Fig. 6A), PHR (Fig. 6B) and PHF (Fig. 6C) scaffolds before and after 8 weeks of in vitro degradation. Bands at 1458 and 2940 cm -1 are assigned to carbon -hydrogen (C-H) stretching vibrations in aliphatic chains of PLA and hyaluronan. The 1740 cm -1 band is associated with vibration of carbonyl groups in lactones. Peaks at 3060, 3080 (valence vibration of C-H band in benzene), and at about 1600 cm -1 (acryl group) are from aromatic ring vibrations. Ester group vibrations are evident as a pair of bands at 1000 -1100 cm -1 and 1100 -1300 cm -1 .

The presented spectral data provide evidence of the hydrolysis of ester bonds in hyaluronic derivatives. As a result, benzoic acid came into the solution, leading to changes in pH during degradation. The peaks assigned to aromatic compounds disappeared after 1 week of degradation in the case of PHF and PH scaffolds and after 4 weeks for PHR scaffold.

c) DSC measurements

DSC thermograms of all the scaffold typologies are shown in Fig. 7A-C. Additionally, the curves regions showing glass transition are presented in Fig. 7D-F. The values of characteristic temperatures (T g , T m ) as well as enthalpies (of melting and crystallization) and the crystallinity degree are summarized in Table 2. The Tg temperatures of microporous PCL/HYAFF11 ® -based scaffolds are close to the values quoted for PCL, i.e. -60 o C (Table 2). The T g of PLA of about 55 o C [START_REF] Södergard | Properties of lactic acid based polymers and their correlation with composition Prog[END_REF] cannot be exactly determined for PHF scaffolds, as it is very close to the T m of PCL (Fig. 7C). Relatively minor changes were observed in T g values over 8 weeks of immersion in SBF for all studied materials.

All the thermograms displayed one sharp melting peak at around 60 o C (Fig. 7A-C) associated with the melting of PCL crystalline phase. The total amount of the crystalline phase of PCL segment increased during soaking in SBF for all samples (Table 2). The melting temperature remained almost unchanged during in vitro soaking. In addition, for PH and PHR scaffolds, a second melting peak at 143-149.5 o C appeared after 1 and 4 weeks immersion in SBF, respectively (Fig. 7A,B). In the case of PH-based scaffold, a small and sharp peak located at about 162 o C and associated with the melting of PLA was also observed (Fig. 7C).

d) TGA measurements

TGA curves of PH and PHR showed a two-step degradation process (Fig. 8A,B) whereas for PHF three stages of degradation were observed (Fig. 8C). Weight losses and decomposition temperatures for the composites after different periods of immersion in SBF are summarized in Table 3

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13 HYAFF11 ® content of these composites varies from 20% in PH to 17% in PHR, a correlation can be drawn between HYAFF11 ® content and weight loss. During soaking in SBF, the weight loss in the first step of degradation decreased and, after 8 weeks, only one step degradation at the temperature range 292-499 o C was observed for PH. For PHR the weight loss dropped to 4% in the first stage degradation after 4 weeks immersion in SBF, whereas, after 8 weeks, a 6% weight loss was observed. The same situation was also observed for the PHF scaffolds. After 1 week soaking in SBF, a three-step degradation process was observed (Fig. 8C). The weight loss decreased from 9% to 7% in the first degradation step, but increased from 8% to 10% in the second.

After 4 weeks of in vitro degradation, the weight loss further increased to 13% in the second step of thermal degradation. After 8 weeks immersion only a one step degradation process, associated with PCL thermal degradation, was observed.

Discussion

The current challenge in scaffold design is to fabricate reproducible bioactive and bioresorbable 3D scaffolds with tailored porosity and pore morphology, capable of maintaining structure integrity for a predictable period, even under load-bearing conditions. Through the time-controlled resizing of pores which occurs during degradation, materials with a predictable degradation rate may guarantee greater penetration of extracellular substance, improve the nutrient/metabolites exchange and offer a better connection between neighbouring cells, so encouraging and promoting the growth of new tissue. Apart from creating the physico-chemical environment and the structural integrity required for tissue regeneration, scaffolds also should act as a local regulator to control the release of signals, preferably without the use an additional drug carrier [START_REF] Mouriño | Boccaccini AR Bone tissue engineering therapeutics: controlled drug delivery in three-dimensional scaffolds[END_REF].

Here, three biodegradable scaffolds, based on PCL as a major component, and with tailored microstructure and multi-scale degradation rates, were investigated to estimate the dynamic changes occurring in the pore network during the resorption. Recently, an upward trend in PCL use in biomaterials for tissue engineering has been observed. This is due to the high processability of such materials into composite structures with superior mechanical and biocompatible properties compared to the polymer on its own [START_REF] Luciani | PCL microspheres based functional scaffolds by bottom-up approach with predefined microstructural properties and release profiles[END_REF][35] [START_REF] Guarino | Porosity and mechanical properties relationship in PCL based scaffolds[END_REF].

Three different typologies of composite scaffolds have been proposed by the integration into the PCL matrix of different material phases: Hyaff11, as a hydrophilic cue; R-cem (particles) as reinforcement systems by using conventional phase inversion/salt leaching technique alone (i.e, PH, PHR); or incorporation of PLA long fibers by filament winding technique (i.e, PHF). All scaffold prototypes showed a bimodal porosity, characterised by macropores arising from salt crystal dissolution (Figure 2), and micropores formed by the controlled removal of solvent via non-solvent exchange (Figure 2). These characteristics assure an open-pore network, highly capable of supporting cell adhesion and proliferation, preserving the correct and requisite fluid transport, and sustaining the external mechanical stress until the regeneration of new functional tissue is completed. The investigation of morphology by X-ray microcomputed tomography (microCT, or µCT), whose use has recently been adopted for the study of soft materials and natural tissues [START_REF] Mawatari | Quantitative analysis of three-dimensional complexity and connectivity changes in trabecular microarchitecture in relation to aging, menopause, and inflammation[END_REF] [START_REF] Lin | Microarchitectural and mechanical characterization of oriented porous polymer scaffolds[END_REF] in place of invasive intrusion-based techniques [START_REF] Gupta | Characterization of a novel bioactive composite using advanced X-ray computed tomography[END_REF], allowed a quantitative assessment of scaffold porosity features (pore size, structural porosity and degree of anisotropy) as summarized in Table 1. The data here clearly show the contribution of polymer phases with different chemical and physical properties on determining the final pore morphology of the scaffold. Porosity and pore sizes of PH scaffolds are firstly affected by the response of the HYAFF11 ® phases to the water retention, leading to a slight underestimation of pore sizes (around 170 µm) due to swelling of the hydrogel through absorption of environmental moisture.

The moderate increase of pore sizes in the case of PHR scaffolds may be explained by the effect of rigid R-cem calcium phosphate particles, which partially constrain the water induced swelling of the HYAFF11 ® phases. However, it should not to be discounted that the reduced loss of R-cem particles due to the high water retention of the hydrogel during the drying of the scaffold may also contribute to this effect. In the case of PHF scaffolds incorporating PLA fibre, a drastic decay of the porosity degree, from 81.7% to 63.1%, was observed, when a smaller weight amount of sodium chloride crystals was used in order to reduce the solution viscosity, and thus optimize the fibre/matrix adhesion, during the winding process.

Accurate microCT investigation also affords the opportunity to identify the spatial distribution of polymer phases with different chemical properties and so evaluate their contribution on the composite degradation behaviour. In particular, by regulating the voltage/current characteristics, microCT enabled the identification of distinct single phases within the composite (Fig. 4), and their spatial distribution. In turn, this enabled an informed optimization of the composite preparation process. By comparison domains with different radiopacity, it has been possible to identify the contributions of: individual rates of degradation are considerably different [40][41]. The most significant differences in mass loss were observed after 4 weeks of immersion. It is assumed that most of the mass loss at the first 4 weeks of degradation was due to degradation products of HYAFF11 ® -p75 leaving the system by dissolving into the aqueous medium [START_REF] Milella | Chemico-physical properties of hyaluronan-based sponges[END_REF]. The decreasing amount of weight loss attributable to thermal decomposition of HYAFF11 ® segments [START_REF] Tsuji | Hydrolytic degradation of poly(3caprolactone) in the melt[END_REF], as immersion time in SBF increases, seems to confirm this. This is also in good agreement with Raman spectroscopy analysis, which showed an absence of bands associated with aromatic ring vibrations in all those composite samples examined after either 1 or 4 weeks immersion in SBF. The rapid degradation rate of HYAFF11 ® -p75 correlates very well with the findings of Avitabile et al [START_REF] Avitabile | Biocompatibility and biodegradation of intravitreal hyaluronan implants in rabbits[END_REF],

who reported that HYAFF11 ® -p75 based implants were fully resorbed after 15 days implantation in rabbit.

The higher degradation rate of PHR composite may be due to the presence of the highly hydrophilic R-cem particles. The accelerating effect of various ceramic fillers on PCL degradation rate has been previously reported [45][46]. Hydrophilic fillers can facilitate water absorption, leading to faster degradation of HYAFF11 ® macromolecules, explaining why PHR composites show the highest mass loss of all the composites studied during the first 4 weeks of immersion in SBF. The mass loss can arise from the diffusion of the ceramic filler from the sample into solution, so creating voids and microcracks within the sample. After this initial period, its degradation rate clearly diminished (Fig. 5).

A slower degradation rate directly results from an increased degree of crystallinity, as revealed by DSC analysis (Table 2). Polyester degradation starts with water uptake [START_REF] Södergard | Properties of lactic acid based polymers and their correlation with composition Prog[END_REF].The diffusion of water into loosely packed amorphous regions is much easier than its ingress into crystalline domains, and accordingly the first stage of hydrolytic degradation is located in the amorphous region of the sample. The initial degree of crystallinity for PH was more than 10% higher than that for PHR and PHS composites. Also, the crystallinity of PH increased by 10% after immersion in SBF for 1 week (Table 2), whereas for the crystallinity of PHR and PHF samples remained unchanged over the same period. The slow increase in crystallinity enabled faster diffusion of water into the PHR and PHF matrices during the next 3 weeks of soaking in SBF, and thus faster hydrolysis leading to higher mass loss (Fig. 5). As degradation proceeds by random scission of ester bonds, the remaining chain segments gain more space and mobility, leading to reorganization of the polymer chains and an increase in crystallinity. In the case of PH and PHR, shorter more labile polymer chains retained in the polymer matrix generated a biphasic crystalline phase, manifesting itself in the presence of additional melting peaks at around 147 o C, which appeared after immersion in SBF (Figs. 6 A-B).

.

Conclusion

Three biodegradable scaffolds based on a PCL and HYAFF11 ® matrix, with tailored chemical and physical properties, were developed for bone and/cartilage tissue regeneration. The combined use of SEM and computer microtomography (mCT), supported by digital image analysis, enabled a non-destructive investigation of the diverse porous microstructures of the proposed scaffolds. This investigation demonstrated the contribution of single composite phases to the basic morphological features. Furthermore, in vitro degradation studies have demonstrated the need to reach a compromise between the material chemistry and the composite structural design to predict ab initio the dynamic changes which may occur as the hydrolytic degradation mechanism progressively evolves, paving the way for a more informed use of polymeric composite scaffolds with multiscale degradation in bone tissue regeneration.

Captions to Figures

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  planar 2D slices (Figs 4, right side) in order to describe the scaffold architecture and to correlate structural differences with chemical composition changes. These slices confirm the preliminary indications obtained by SEM regarding pore size and homogeneous spatial distribution of pores. By comparing 2D slices from PH, PHR and PHF scaffolds (Fig.4, right side), the presence of localized domains with different radiopacity as a function of material density are recognized. In Fig 4, white regions indicate pores, whereas slightly contrasted regions (cyan regions) identify the PCL matrix, which shows a lower susceptibility to X-ray adsorption due to its reduced density. In turn, blue and orange regions are representative of HYAFF11 ® inclusions and calcium phosphate particles, respectively, as a function of the increasing material density.

  . The maximum temperature (peak) of the first stage thermal degradation related to the HYAFF11 ® decomposition, ranged from 222.8 o C to 255.0 o C. The second and third stage of thermal degradation observed in the temperature range of 333 o C -361 o C, and 399.6 o C -403.8 o C, can be attributed to the decomposition of PCL and PLA macromolecules, respectively. Initial samples of PH and PHR showed a weight loss of 15% and 14%, respectively, in the first step. As the

  I)a non-homogeneous spatial distribution of calcium phosphate particles in PHR scaffolds with the clusters formation (orange region), and II) more dense polymer phases of HYAFF11 ® (blue region), which seem preferentially to cover the boundaries of the ceramic particles due to the adsorption of the water molecules during the scaffold preparation process.Our studies on in vitro hydrolytic degradation of PCL/HYAFF11 ® -based scaffolds indicate the strong influence of the extent, interrelationship and distribution of the phases of the various polymer components -indeed, the influence of what might be considered the "design" generated by these chemically distinct phases. Although all the polymeric components (i.e., PCL, PLA, and HYAFF11 ® ) of the composites studied undergo degradation via random hydrolytic chain scission of the ester linkage, their

Figure 1 :

 1 Figure 1: S.E.M. images of cross-sections of PCL/HYAFF11 ® -based composite scaffolds:. A) PH, B) PHR, and C) PHF.

Figure 2 :

 2 Figure 2: Microporosity of PCL/HYAFF11 ® based composite scaffolds by SEM analysis. Micropores (A) along PH scaffold struts and (B) onto PHF outer surface.

Figure 3 :

 3 Figure 3: Morphology of PCL/HYAFF11 ® based composite scaffolds by FE-SEM analyses. (A) Detection of hydrophilic HYAFF11 ® (dark gray) patches (by arrows)

Figure 4 :

 4 Figure 4: MicroCT analysis: 3D Rendering (left) and 2D slice maps of material phase density variations (right) of proposed PCL/HYAFF11 ® -based composite scaffolds: (A)

Figure 5 :

 5 Figure 5: Mass loss of the composite scaffolds as a function of degradation time in SBF solution.

Figure 6 :

 6 Figure 6: Raman spectra of (A) PH, (B) PHR and (C) PHF scaffolds before and after degradation.

Figure 7 :

 7 Figure 7: DSC thermograms of PCL/HYAFF11 ® composites before and after immersion in SBF. (A-D) PH, (B-E) PHR and (C-F) PHF.

Figure 8 :

 8 Figure 8: Thermogravimetric curves of PCL/HYAFF11 ® composites before and after immersion in SBF. (A) PH, (B) PHR and (C) PHF.

Figure

  Figure7A,B,C

Table 1 . Porosity features obtained on the different composite scaffolds by the MicroCT analysis

 1

			Average
		Total Porosity		Anisotropy	Structure Model
	Type		macropore sizes
		degree (%)		Degree	Index (SMI)
			(10-500 m)
	PHS2	81.7	170.4	0.215	1.947
	PHRS2	82.8	198.7	0.226	2.002
	PHF2	63.1	230.5	0.1	0.951

Table 3 . Weight losses and decomposition temperatures for the composites after different time of immersion in SBF

 3

								Second
			First degradation						Third degradation
	Type	Degradation time, weeks	temperature ( o C)	Weight loss, %	degradation temperature ( o C)	Weight loss, %	temperature ( o C)	Weight loss, %
			Start Peak End		Start Peak End		Start Peak End
		0	177 233 303	15				-	303 400 504	77
		1	142 242 296	13		-		-	296 400 507	80
	PHS2
		4	154 222 292	7				-	292 402 529	86
		8	-	-	-	-				-	292 404 499	91
		0	177 239 311	13				-	311 401 576	69
		1	150 243 300	14		-		-	300 403 548	67
	PHRS2
		4	161 231 269	4				-	269 402 500	84
		8	148 254 286	7				-	286 399 510	76
		0	174 233 265	9	265 311 334	8	334 403 490	75
		1	199 247 266	7	266 308 339	10	266 308 339	75
	PHF2
		4	-	-	-	-	263 339 352	13	352 409 501	85
		8	-	-	-	-	-	-	-	-	271 409 479	99

Acknowledgements

This study was financially supported by the IP STEPS EC project, FP6-500465.

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E Stodolak

T Gumula

R Leszczynski

J Wieczorek

S Blazewicz

A composite material used as a membrane for ophthalmology applications

Keywords: polymer-matrix composites (PMCs), short-fibre composites, mechanical properties, porosity, casting

New biomaterials for intracorneal ophthalmologic implants were designed, manufactured and characterized. A composite material in the form of a membrane was manufactured in a two-stage process. The first stage of the process depended on preparation of multidimensional (MD-type) fibrous polymer composite. A stable terpolymer polytetraflouroethylene-co-polyvinylene fluoride-co-polypropylene (PTFE-PVDF-PP) was used as a composite matrix, and sodium alginate-based biopolymer (NA) in the form of short fibres and/or powder were used as porogenic constituents. The composite materials were subjected to physicochemical treatment in order to remove a water soluble biopolymer. The treatment led to about 50% of open porosity within the polymer matrix. Depending on the membrane type the mean pore size determined with SEM microphotographs was 15 to 25 µm. Permeability and durability of the membranes in simulated eye fluid (culture medium enriched with albumin) was tested. The size and shape of the pores before and after the permeability test were compared (SEM), and they depend on the porogen form. Mechanical parameters of the composite materials such as; tensile strength, Young's modulus, and strain to failure were measured. A membrane derived from fibres and particles showed better mechanical properties than a membrane derived from porogen particles. Microstructure and mechanical properties make the membranes a good candidate for ophthamological implants.

Introduction

Glaucoma is a group of eye diseases characterized by an increase in intraocular pressure that may lead to pathological changes in the optic nerve and subsequently to the vision loss and blindness. Most patients with glaucoma require only medications to control the eye pressure. When medicine therapy fails in lowering the eye pressure a surgery is required. It often involves the use of laser treatments or membrane implants. The objective of any glaucoma surgery is to enable the eye fluid to circulate within the eye easier. Worldwide, it is estimated that about 6.9 million people have visual impairment from glaucoma, and 6.7 million suffer from blindness [START_REF] Eldaly | Non-penetrating filtration surgery versus trabeculectomy for open angle glaucoma[END_REF]. The main role of implants in glaucoma treatment is lowering the internal eye pressure by improving the flow of an aqueous humour from inside of the eyeball. Such the membrane implants should facilitate filtration after the surgical treatment [START_REF] Nuengjamnong | Membrane fouling caused by extracellular polymeric substances during microfiltration processes[END_REF][START_REF] Nizankowska | Prevalance of open angle glaucoma and ocular hypertension as a risk factor for primary open angle glaucoma in Wroclaw population[END_REF].

The requirements for any material used for ophthalmologic implants are in general sense the same; the material must be biocompatible, biostable and should have the microstructure suitable for fluid transport [START_REF] Shaarawy | Comparative study between deep scleretomy with and without collagen implant: long term follow-up[END_REF][START_REF] Wevill | A pilot study of deep sclerectomy with implantation of chromic suture material as a collage implant: medium -term results[END_REF]. Such implants are to date made of natural polymers, e.g. collagen, hiauronians or synthetic polymers, e.g. expanded polytetraflouroethylene (e-PTFE), polypropylene (PP), silicones. The main drawback of resorbable membranes made from collagen or hiauronians, is their instability and low permeability [START_REF] Chiou | An Ultrasound Biomicroscopic Study of Eyes After Deep sclerectomy with Collagen Implant[END_REF]. On the other hand, such materials are characterized by good biocompatibility and are able to mimic the structures of natural tissues. For a long-term use the materials applied for the cornea permanent implants have to withstand static and dynamic pressure changes, and be stable in biological environment [START_REF] Shaarawy | Comparative study between deep scleretomy with and without collagen implant : long term follow-up[END_REF]. In order to ensure the proper fluid permeability an ophthalmologic membrane should display an optimum interconnected pore system, total porosity and the mean pores size. Too small pores may be easily blocked by molecules present in the aqueous humour, while larger pores may undergo deformation due to high intraocular pressures [START_REF] Shaarawy | Comparative study between deep scleretomy with and without collagen implant : long term follow-up[END_REF][START_REF] Kontturi | Adsorption of globular proteins on polymeric microfiltration membranes[END_REF]. The optimal membrane permeability may be obtained for various combinations of pore size and total porosity. Typical pore size for synthetic polymer membranes is in the range of few micrometers to 100 µm [START_REF] Zalewska | Non-penetrating deep sclerectomy with SK-GEL implant in the patients with pseudoexfoliation glaucoma[END_REF].

Composite materials are more and more frequently used for design and production of implant materials in the wide range of biomedical applications [START_REF] Ramakrishna | Biomedical applications of polymer-composite materials: a review[END_REF][START_REF] Feng | Preparation and histological evaluation of biomimetic three-dimensional hydroxyapatite/chitosan-gelatin network composite scaffolds[END_REF][START_REF] Toshihiro | Preparation and mechanical properties of polylactic acid composites containing hydroxyapatite fbers[END_REF].

They acquired usefulness due to their capability of controlling mechanical properties of the pure polymers and adjusting them to the properties of the damaged tissue. The composites having differently oriented fibrous reinforcement component (1D, 2D, 3D and MD) within the matrix can be used to fabricate load-bearing implants for bone surgery [START_REF] Blazewicz | Experimental study of mechanical properties of composite carbon screws[END_REF][START_REF] Gasser | About composite materials and their use in bone surgery, Injury[END_REF][START_REF] Fujihara | Performance study of braided carbon/PEEK composite compression bone plates[END_REF][START_REF] Gross | Biodegradable composite scaffolds with an interconnected spherical network for bone tissue engineering[END_REF]. The proper design and choice of the composite constituents for ophthalmologic biomaterials enable to obtain a biomimetic structure similar to the natural tissue required for filtration of the nourishing fluid flow into and out of the eye, i.e. drainage channel. In the case of MD composite the effective dispersion of the short fibres in a polymeric matrix enables modification of the whole volume of the material and ensures creation of a continuous fibres network [START_REF] Xie | In vivo bone regeneration using a novel porous bioactive composite[END_REF]. The suitable selection of the fibrous phase as a porogen and further physicochemical treatment of the composite material enables the formation of interconnected pores within the material [START_REF] Stodolak | Membrane composite materials for medical application -primary material and biological study[END_REF]. Such a material may be an alternative for commonly used polymer materials.

The aim of this work was to elaborate a biocompatible composite membrane that can be used in the glaucoma treatment. Microstructure, mechanical parameters and biocompatibility of the materials in the form of membrane were optimized. The effect of a simulated biological environment on the elaborated composite materials was investigated. The membrane was made of a stable synthetic polymer, i.e. PTFE-PVDF-PP constituting a composite matrix, and a fibrous phase made of sodium alginate being a water soluble porogenic component for the creation of the desired porosity in the polymer matrix. PTFE-PVDF-PP terpolymer was chosen for a matrix due to its suitable mechanical properties (flexibility, creep behaviour, tensile strength and durability), biocompatibility, surface physicochemical properties, i.e. hydrophobicity and easier processability comparing to its pure polymer constituents, particularly to PTFE. The terpolymer was successfully used in other biomaterial applications [START_REF] Stodolak | Membrane composite materials for medical application -primary material and biological study[END_REF][START_REF] Zaczy Ska | The study of cytotoxity of fibrous composite materiale[END_REF][START_REF] Stodolak | In vitro behavior of PP-PVDF-PTFE terpolymer modified with alginate fibres[END_REF][START_REF] Blazewicz | Fibrous, bioactive composite materials for bone regeneration[END_REF][START_REF] Stodolak | The effect of surface parameters of fibrous composite materials on cellular response[END_REF][START_REF] Stodolak | Composite membranes materials for ophthalmological implants[END_REF].

The study was devoted to preparation and characterisation of the membrane materials having suitable mechanical and fluid transport properties in the implant site.

Materials and methods

PTFE-PVDF-PP terpolymer (Aldrich Chemical Co., USA) and sodium alginate (Biopolymer, AS Protanal LF 20/60, with dominating blocs of guluronic acid, ~ 65 %, Norway) were used to prepare the composite membranes. Molecular weight of the terpolymer from solution viscosity measurements was approx. 300 kDa. A 1% w/v solution sodium alginate in water has a viscosity at 20°C of 1,2 cP. Two types of composite materials were made. The first type was prepared by using 2 wt. % of sodium alginate (NA) in the form of a powder. The second type of samples was prepared by using 2 wt. % of sodium alginate in the form of two-fraction additive, i.e. short fibres and small particles (NA powder). The two-fraction additive was prepared by grinding of short alginate fibres in a rotating mill. The aim of the fibres fragmentation was to improve their ability to create a continuous interconnected pore network in the membranes. The alginate fibres were prepared by Department of Man-Made Fibres of Technical University in Lodz using the wet method described elsewhere [START_REF] Bogu | Properties of Sodium Alginate Spinning Solutions with Ceramic Nanoadditives[END_REF].

Weighted amounts of the porogen (2 wt.%) were mixed with the polymer solution (acetone, 1:10). Then the mixture was poured on a Petri dish and dried at ambient Permeability of the membranes was investigated using a special test device allowed to simulate variable pressure conditions (dynamic mode) inside the eye. The eye simulator device illustrated in Fig. 2 was designed to compare changes in the membranes subjected to intraocular pressure in a normal eye and the pressure existing in an eye with glaucoma disease. During the experiments the test device consisted of piston (driven by an electric motor) pumped the fluid through the membrane to simulate pressure conditions taking place during a severe attack of glaucoma, i.e. the cyclic pressure exerted on the membrane was in the range of 0-55 mmHg (5h/37 o C). As a biological environment a bovine serum modified with 5% of albumins was used (SEF). Like in a case of the static conditions, the membrane was placed between two chambers -one filled with distilled water and the second one filled with the serum.

The permeability was determined by measuring the ionic conductivity and pH of the serum. After the tests in the eye simulator, the membranes microstructure and their thickness were measured. Mechanical properties of the composite samples and membranes, namely tensile strength (R m ), Young's modulus (E), elongation at the maximum force (ε F max ), and work up to failure (W) were measured using an universal testing machine Zwick 1435.

All presented results are an average from five measurements, and the error was estimated basing on t-Student test at significance level of 0.05.

Results

The microscopic observw3ations of the sodium alginate short fibres after grinding (fragmentation) in the laboratory rotating mill indicate that the fibrous phase consisted of short fibres (diameter of 0.17 ± 0.05 m, length of 1.50 ± 0.28 mm; 95 % of the population), and much smaller particles (diameter of 0.75 m; 5 % of the population), which were products of their fragmentation (Fig. 3). Diversification of the porogen particles sizes ensures the formation of the network of interconnected pore system. The pore sizes in polymer matrix correspond to the sizes of the fibres and products of their fragmentation. The alginate powder used as a porogen in the other membrane type consisted of particles with similar sizes (Fig. 3, diameter 0.45÷0.75 m), which were the origin of pores of analogous characteristics.

After removal of the porogen components (powder, fibres) both membranes exhibited surface porosity and similar thickness of about 180 -200 µm. Alginate fibres and powder -based membrane (M1) exhibits two types of pore fractions, i.e. large, elliptical pores formed by washing out the fibrous porogen (small amount), and round and small in size pores formed by washing out the NA powder (larger amount).

Microphotographs of the cross-section areas of the membranes revealed that distribution of pores was uniform in whole volume of samples (Fig. 4). A well-developed system of interconnected pores was particularly visible in the case of M2 membrane, in which the porogen consisted of a mixture of fibres and particles of a different size. The presence of the interconnected pores was confirmed by ion conductivity test performed in static conditions (Fig. 5).

A qualitative analysis of several microphotographs revealed that both membranes have similar porosity, slightly higher for M1 membrane (c.a. 54%) as compared to M2 membrane (c.a. 49%). The estimated mean pore size of the membrane surface was close to the size of the removed porogen particles/fibres and corresponded to c.a. 25 µm, and c.a. 15 µm for M1 and M2 membranes, respectively. These values of pore sizes make the membranes suitable for transport of aqueous humour containing macromolecules of proteins such as albumin and inorganic salts. Such pores do not affect permeability of ions, and the changes of ion conductivity observed during the static test (Fig. 5) may be attributed to removal of porogen particles/fibres which were not fully washed out during the membrane fabrication. The differences between the residual porogen in M1 and M2 membrane resulted probably from the difference in the form of the porogen. Relatively larger fibrous porogen present in M2 membrane was washed out faster and after 4 days the conductivity remained constant. Increasing ion conductivity in the case of M1 membrane indicated that smaller porogen particles were constantly washed out during the test.

A test performed in static conditions revealed that Na + and Cl -ions migrated through M1 membrane faster, which may be related to its a higher pore size and a better developed network of interconnecting pores (Fig. 5).

Permeability of the membranes was investigated in dynamic in vitro experiment using the set-up to simulate pressure conditions and fluid transport in the eye (Fig. 2). The membrane placed in the chamber with the simulated eye fluid (SEF): bovine serum containing albumins was subjected to variable pressure loading (0-55 mmHg). Despite relatively high pressure exerted on the membranes they did not lose their permeability which indicates, that the pores were not blocked by the protein molecules, and the fluid circulation was not hindered. Also the ion conductivity of the solution used to verify the permeability remained at the same level during the whole time of the experiment (Fig. 6). The runs of conductivity measurements shown in Fig. 6 for water and bovine serum confirm the stable fluid transport through the membranes.

Both membrane materials were stable in in vitro conditions since no chemical and physical changes were observed after 24 months of incubation in water at 37 0 C. These preliminary results indicate that the membrane materials studied may find application as permanent implants in glaucoma treatment.

The mechanical tests of the membranes showed inferior mechanical properties in comparison with the pure polymer foil (the reference sample). In the case of the membranes the decrease in tensile strength (R m ), Young's modulus (E), work to failure (W), and strain to failure (ε F max ) were observed, Table 1. These changes resulted from higher porosity of the membrane materials.

Dynamic tests performed on the membranes subjected to constant and variable pressures indicated that initially determined mechanical properties did not change after the test. Moreover, the shape and diameter of the pores were stable. The membrane thickness remained stable after the tests simulating an eye environment. Changes in microstructure of the membranes under the applied load were negligible and did not influence their permeability.

The mechanical fatigue properties of the membranes were tested under static and dynamic pressure loads. The average diameter of pores was higher than that observed in membrane materials without pressure loading, i.e. 30-35 µm, and 35-40 µm for M1 and M2 membrane, respectively. Neither of the membranes was fractured nor damaged after the fatigue mechanical tests.

Conclusions

Application of water soluble sodium alginate fibres for polymer modification is a useful method for the preparation of a membrane material with a controllable pores microstructure. The membrane material based on the sodium alginate fibres showed the highest strain to failure and an enhanced tensile strength. The lack of physical changes observed in in vitro conditions suggests that the membranes are stable.

Both types of the membranes performed a relatively high strength and elasticity (high strain to failure), which make them suitable for planned application. Time test simulating an eye environment (fluid, pressure) demonstrated that the elaborated materials did not loss their properties. It may be expected that such membranes will be suitable for a long-term performance implant. The obtained data, however indicate that further study should be continued to evaluate the manufactured glaucoma implant on in vivo model. M2 material is a better candidate for a membrane for glaucoma treatment due to its mechanical properties, sufficient pore size and fraction, possibility for removing completely the porogen as well as stability of the mechanical characteristics in in vivo conditions.

  temperature for 48h in order to evaporate the solvent and finally the samples in the form of thin foils were obtained. In order to remove the porogen and to create open porosity within the materials the polymer foils were incubated in distilled water at 80 o C for 84h. The scheme of preparation of the composite membranes differing in the porogen component is shown in the Fig 1. Using this method two kinds of membranes were manufactured; M1 -a membrane derived from the composite containing NA particles only (NA powder), M2 -a membrane derived from the composite containing NA fibres and the particles (NA powder) prepared by additional fragmentation (grinding) of the fibres. The porogen morphology, i.e. sodium alginate fibres/particles was studied under SEM (Jeol 5400). Particle size distribution of the powder was determined by DLS method (Nanosizer-ZS), and fibres size (diameter, length) was measured with the optical microscope (Lanametr, PZO, Poland). Permeability of the membranes in static conditions was tested by measuring the electrolytic conductivity of distilled water separated by the membrane from NaCl solution (0.1M NaCl/37 0 C/3 weeks).

Fig 1 .Fig 2 .Fig 3 .Fig 4 .Fig 5 .Fig 6 .Fig 7 .

 1234567 Fig 1. Schema of preparation of composite membrane M1 and M2

Tabel 1 .

 1 Mechanical properties of membrane materials: M1 and M2 R m -tensile strength, ε F max -elongation at the maximum force, E -Young's modulus, W -work up to break

	Material	Rm [MPa]	Fmax [%]	E [MPa]	W [J]
	Membrane M1	0.28 ± 0.06	258 ± 39	0.12 ± 0.06	0.22 ± 0.06
	Membrane M2	0.35 ± 0.02	485 ± 45	0.19 ± 0.04	0.37 ± 0.08
	Polymer foil	0.5 ± 0.06	726 ± 53	0.49 ± 0.14	0.41 ± 0.03

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Acknowledgements

This work was supported by The Polish Ministry of Science and Higher Education, grant No. N518 028 32/1769. Ewa Stodolak is a beneficent of programme START of Foundation for Polish Science in 2009.