In vitro and in vivo behavior of self-reinforcedbioabsorbable polymer and self-reinforced bioabsorbablepolymer/bioactive glass composites

Henna Niiranen,1 Tuomo Pyhalto,2 Pentti Rokkanen,2 Minna Kellomaki,1 Pertti Tormala1

1Institute of Biomaterials, Tampere University of Technology, P.O. Box 589, FIN-33101 Tampere, Finland2Department of Orthopaedics and Traumatology, University of Helsinki, Topeliuksenkatu 5,FIN-00260 Helsinki, Finland

Received 22 August 2002; revised 9 February 2004; accepted 11 February 2004Published online 29 April 2004 in Wiley InterScience (www.interscience.wiley.com). DOI: 10.1002/jbm.a.30043

Abstract: The aim of this study was to investigate the invitro and in vivo properties and degradation of (1) self-reinforced (SR) lactide copolymer, P(l/dl)LA 70:30, and (2)SR composites of the same polylactide and bioactive glass13-93. The following three polymer and polymer–bioactiveglass samples were studied: SR-PLA70, SR-PLA70 �BaG15s, and SR-PLA70 � BaG20c. In vitro behavior wasstudied in a phosphate-buffered saline for 87 weeks at 37° �1°C and a pH of 7.4 � 0.2. In vivo behavior was studied byimplanting the rods in the dorsal subcutaneous tissue of rats(SR-PLA70 � BaG20c) or rabbits (SR-PLA70 and SR-

PLA70 � BaG15s) for 48 weeks. The degradation of thespecimens was evaluated by measuring the changes in me-chanical properties, crystallinity and molecular weight ofpolymer, water absorption, weight loss, and structuralchanges. Results showed that the addition of bioactive glassfiller modified the degradation kinetics and material mor-phology. © 2004 Wiley Periodicals, Inc. J Biomed Mater Res69A: 699–708, 2004

Key words: bioactive glass; bioabsorbable; composite; self-reinforced; degradation

INTRODUCTION

Self-reinforced (SR) bioabsorbable polymers with suf-ficient initial strength and suitable strength retentionhave been studied since the 1980s for internal fixationdevices in bone fractures.1,2 Bioabsorbable materialsare the method of choice in cases in which only tem-porary support for healing tissue is needed. However,most of these polymers do not have properties toaccelerate or facilitate tissue healing. Bioactivity andthus bone bonding ability would be advantageous forimplants used in bone tissue applications.

Since the 1990s, various ceramic filler materials havebeen added to bioabsorbable polyesters to improveosteoconductivity and mechanical properties, and tocompensate for the acidity of the degradation prod-ucts. Filler materials such as carbonated apatite3 car-bonated calcium phosphates, calcium carbonate,4,5

and Bioglass�6 have been studied to discover whetherthey stabilize or control the acidity of degradationproducts of biodegradable polyesters in vitro. In addi-tion, ceramic fillers have an effect on the structure,mechanical properties, strength retention, degradationrate, water absorption (WA), weight loss (WL), andbioactivity of the composites.7–11

To enhance bone fracture healing, we manufacturedSR composites containing bioactive glass 13-93 as anosteoconductive filler in bioabsorbable polymer ma-trix. Bioactive glass 13-93 has been shown to bond tobone through a CaP layer in vivo and it has been foundto be resorbable in soft tissues.12 Using this particularfiller, bioactive glass 13-93, and the poly(l/dl)lactide70:30 matrix bioactive/bioabsorbable composites canbe made by conventional melt-extrusion. The mechan-ical and structural properties can be further variedwith a self-reinforcing process.

Several studies have shown that the degradationmechanism of aliphatic polyesters, such as polylac-tides, appears to be relatively complex.13–20 There aremany factors that affect the hydrolytic degradation ofnon-reinforced and SR bioabsorbable polymers.21,22

Composite structures with ceramic additives are evenmore complicated to predict, and therefore it is essen-

Correspondence to: H. Niiranen; e-mail: henna.niiranen@tut.fi

Contract grant sponsor: Academy of Finland (BiomaterialResearch Group, Finnish Centre of Excellence Programme2000–2005); contract grant number: 44863

© 2004 Wiley Periodicals, Inc.

tial to determine how the filler addition will influencethe properties and degradation characteristics of thecomposite. The aim of the present study was to inves-tigate the material properties and degradation behav-ior of SR bioabsorbable polymer (SR-PLA70) and itsbioactive glass-containing composites (SR-PLA70 �BaG) in vitro and in vivo.

MATERIALS AND METHODS

Materials

The bioabsorbable polymer used for PLA70 rods and usedas a matrix material for composite rods was a lactide stereo-copolymer, poly(l/dl)lactide 70:30, with the trade nameResomer LR 708 (Boehringer Ingelheim, Ingelheim am Main,Germany). The polymer is medical grade (highly purified),has an inherent viscosity of approximately 6.0–6.5 dL/g,and has a residual monomer content of �0.5% (both re-ported by the manufacturer).

The bioactive filler for the composites was bioactive glass13-93, with composition 6 wt % Na2O, 12 wt % K2O, 5 wt %MgO, 20 wt % CaO, 4 wt % P2O5, and 53 wt % SiO2.12 Eithercrushed or spherical particles were used (Vivoxid Ltd.,Turku, Finland) and in both cases the particle size distribu-tion was 50–125 �m.

Manufacturing the implants

The raw materials were vacuum dried at 80°C for 1 day toremove the excess water and air. Non-reinforced PLA70billets were produced by extrusion in an N2 atmosphereusing a Gimac microextruder (Mac.gi, Castronno, Italy) witha screw diameter of 12 mm and extrusion temperature (die)of 221°C. Previously mixed and dried PLA70 and BaG pow-ders were extruded and thus blended using the samemethod. The glass content in the composites is shown inTable I.

Solid-state die-drawing was used to produce SR rods fromthe initial extruded billets.1 The drawing temperature was60°–70°C and the achieved draw-ratio (DR) was approxi-mately 4.5 for SR-PLA70, 4.5 for BaG20c, and 3.5 for BaG15s.The finished SR rods had a diameter of 2 mm and length ofeither 38 mm (SR-PLA70 � BaG20c) or 42 mm (SR-PLA70and SR-PLA70 � BaG15s). The implants were sterilized with

gamma irradiation (minimum dose 25 kGy). All the materi-als studied and abbreviations used are listed in Table I.

Degradation in vitro and in vivo

In vitro behavior of SR-PLA70 and SR-PLA70 � BaG sam-ples was studied in phosphate-buffered saline. The ionicconcentrations of phosphate-buffered saline were Na� 156mM, Cl� 101 mM, HPO4

2� 24 mM, and H2PO4� 6 mM. The

samples were immersed for 87 weeks at 37° � 1°C, pH 7.4.Vsolution to Vsample ratio was �20 in order to diminish thedecrease of pH between the solution changes. The buffersolution was changed periodically every other week and thepH of the solution was measured (three parallel samples). Ifthe pH changed more than �0.2, the buffer solution waschanged weekly.

For the evaluation of composite degradation in vivo,BaG20c implants were implanted into the subcutaneous tis-sue of rats. SR-PLA70 and BaG15s were implanted into thesubcutaneous tissue of rabbits. The in vivo study on rats wasperformed as a pilot study. This article concentrates on thematerial property changes in vivo compared with in vitro.The more specific in vivo results, animal models, and histo-logical results will be published in a separate article. Thefollow-up times were 3, 6, 12, 24, and 48 weeks. Implantswere stored in 0.9% saline after removal from the animalsand the mechanical properties (wet specimens) were mea-sured within 24–48 h. After mechanical testing, all the invitro and in vivo samples were rinsed with ion exchangedwater and ethanol and vacuum dried before further analy-ses.

Mechanical tests

Both the flexural and shear properties were measured atroom temperature using an Instron materials testing ma-chine (Instron 4411; Instron Ltd., High Wycombe, UK). Flex-ural properties were evaluated using a three-point bendingtest according to standard (SFS-EN ISO 178, 1997). Shearstrength was measured according to the modified standard(BS 2782, method 340 B, 1978) by means of a specific tool.Bending and shear strength and bending modulus values forcylindrical rods were calculated using Equations (1–3). Flex-ural results are the means of four parallel samples in vitroand four or two parallel samples in vivo. Shear data is themean of six repeat samples in vitro and six or four repeatsamples in vivo. Results are given with standard deviation inparentheses.

Bending strength, �f [MPa]

�f �8 � Fmax � L

� � d3 (1)

Bending modulus, E [GPa]

E �4 � slope � L3

3 � � � d4 �1

1000 (2)

Shear strength, [MPa]

TABLE IMaterials Studied

Abbreviationfor GlassType andContent

GlassContent

(%)

Shape ofGlass

Particles DR

SR-PLA70 0 4.5SR-PLA70 � BaG15s BaG15s 15 � 1 Spheres 3.5SR-PLA70 � BaG20c BaG20c 19 � 2 Crushed 4.5

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�2 � Fmax

� � d2 (3)

where Fmax applied maximum force in testing [N], L support span [mm], d diameter of the rod [mm].

Molecular weight measurements and thermalanalysis

Gel permeation chromatography (GPC) analyses werepreformed using a Waters apparatus (Waters, Milford, MA)equipped with one PLgel 5-�m guard column and twoPLgel 5-�m mixed-C columns (Polymer Laboratories, Am-herst, MA). Narrow polystyrene standards were used forcalibration and chloroform was used as a solvent and eluent.The bioactive glass was removed from dissolved samplesusing a 0.2-�m filter before injection. Millennium32 softwareversion 3.05.01 was used for analysis. Mark-Houwink pa-rameters � 0.73, K 5.45 � 10�4 for PLA70 and compositespecimens and � 0.73, K 1.12 � 10�4 for polystyrenestandards were used to calculate the viscosity-average mo-lecular weight (Mv). Averages of Mv and Mn were calculatedfrom two injections (estimated error of each value is within10%).

A PerkinElmer DSC-7 (Norwalk, CT) differential scanningcalorimeter (DSC) was used to determine the transition tem-peratures of polymers. Indium standard was used for cali-bration. The heating rate was 20°C min�1. The level ofcrystallinity was estimated from the values of the heat offusion (melting), using 93.7 Jg�1 for 100% crystallinePLLA.23

WL and WA

The percentages of WA and WL were determined fromthe wet (immediately after hydrolysis), dry (after vacuumdrying), and initial weights of the specimens (three sam-ples). WA and WL were deduced from the equations WA% 100(Ww � Wd)/Wd, WL% 100(Wi � Wd)/Wi, where Wi

was the initial weight of the specimen and Ww and Wd werethe weights of wet and dried specimens.

Scanning electron microscopy (SEM)

SEM was used to study the structure of the compositesbefore and after hydrolysis. Cross- and longitudinal sectionsof the specimens were taken by breaking the rods in liquidnitrogen. The samples were coated with a thin layer of goldbefore examination using a JEOL T100 scanning electronmicroscope (JEOL Ltd., Tokyo, Japan) at an accelerationvoltage of 15 kV.

RESULTS

Mechanical properties

The initial bending strength of gamma irradiatedSR-PLA70 rods was 192 (5) MPa. The bending

strengths of SR composites were 32% (BaG15s) and28% (BaG20c) lower compared with plain SR-PLA70.Bending moduli of BaG15s and BaG20c were 27% and24% lower than the 4.3 (0.2) GPa of SR-PLA70. Exten-sion (at maximum load) was approximately 7% for allsamples studied. The initial shear strengths were 25%(BaG15s) and 30% (BaG20c) lower than 125 (2) MPa(SR-PLA70).

The mechanical properties of all materials studiedhad decreased significantly after 24 weeks in vitro andin vivo (Table II). At 24 weeks in vitro, the bendingstrength of the SR-PLA70 rods had decreased by 27%.Similar rods in vivo had disintegrated and flexuralproperties could no longer be measured. The bendingstrength of SR-PLA70 rods in vivo had decreased by20% from the initial value already 6 weeks after im-plantation. At 24 weeks, the bending strength of theSR glass composites had decreased by 18%/39% (invitro, in vivo) (BaG15s) and 49%/70% (in vitro, in vivo)(BaG20c). The shear strength retention of the materialstudied in vitro and in vivo is presented in Figure 1. At24 weeks, the shear strengths had decreased as fol-lows: 27% in vitro (SR-PLA70), 11%/50% (in vitro, invivo) (BaG15s), and 37%/68% (in vitro, in vivo)(BaG20c).

Hydrolysis

With SR-PLA70, the pH of the buffer solution re-mained constant (above 7.3) for the first 50 weeks invitro, after which the pH decreased to 6.8. The BaG15sand BaG20c composites showed a slight increase insolution pH at the beginning of hydrolysis. Fromweek 2 until the end of hydrolysis, the pH remainedpractically constant (above 7.3).

After 24 weeks in vitro, the SR-PLA70 samples be-gan to swell. By the end of 48 weeks, the diameter ofthe dried samples had increased 50–60%. Glass-con-taining composites retained their diameter until week75. Thereafter, measurements were not possible be-cause of the fragility of the specimens.

With SR-PLA70, WA increased steadily reaching theapproximate 20% level at the 48th week. Between 48and 87 weeks, a remarkably more rapid increase in

TABLE IIMechanical Properties (B.S., Bending Strength; S.S.,Shear Strength) at 24 Weeks In Vitro and In Vivoa

B.S. (MPa)In Vitro

B.S. (MPa)In Vivo

S.S. (GPa)In Vitro

S.S. (GPa)In Vivo

SR-PLA70 140(5) n/m 91(4) n/mBaG15s 106(4) 79(10) 84(3) 79(7)BaG20c 71(5) 41(3) 55(1) 46(11)

a Values are averages of six samples [mean (STD)]. n/m,not measurable, disintegrated.

SELF-REINFORCED BIOABSORBABLE POLYMER 701

WA up to 300% was observed. For SR composites, WAwas 15% for BaG15s and 30% for BaG20c samples after36 weeks and remained at that level until 48 weeks invitro. Then, WA attained 50% for BaG15s and 100% forBaG20c samples at 87 weeks in vitro.

Until 48 weeks, no WL was detected in SR-PLA70.Thereafter, a remarkable WL was seen simultaneouslywith a considerable increase in water uptake and de-crease in pH. The WL of BaG15s and BaG20c compositeswas observed after 24 weeks. With BaG20c, the WL wasslightly greater. When plotting the material WL as afunction of viscosity average molecular weight (Mv), itwas observed that with Mv values between 22,000 and5000 Da, the WL of BaG15s and BaG20c composites was

up to 20%. The WL of SR-PLA70 was not detected untilthe Mv was �5000 Da. The WL of SR-PLA70 was 20% atan Mv value of 3000 Da.

Molecular weight

Viscosity average molecular weight (Mv) retentionof the in vitro samples is shown in Figure 2. The Mvbegan to decrease as soon as the rods were eitherimplanted or immersed in a buffer solution. The Mv ofSR-PLA70 was 50% of the initial value at around 18weeks in vitro compared with 12 weeks in vivo. The Mv

Figure 1. Shear strength retention of SR-PLA70 and SR composites of BaG15s and BaG20c.

Figure 2. Retention in viscosity average molecular weight (Mv) of SR-PLA70 and SR composites of BaG15s and BaG20c asa function of in vitro time.

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of composites was 50% of the original Mv at 24 weeksin vitro and for BaG15s at approximately 22 weeks andfor BaG20c at approximately 18 weeks in vivo. Thedifference in the polymer degradation of SR-PLA70and BaG15s in vitro and in vivo is seen clearly in Figure3, where the molecular weight distribution (MWD)curves after 48 weeks are presented. With all the ma-terials studied at the beginning of hydrolysis and until24 weeks, the polydispersity was around 2, afterwhich it increased to 5–7. By the end of the hydrolysis,the polydispersity had decreased to 2–3.

Initially, monomodal MWD of the SR-PLA70 showeda shoulder at week 52 in vitro (Fig. 4). The MWD ofBaG20c had transferred to bimodal after 75 weeks in vitroand that of BaG15s after 87 weeks in vitro. The bimodal-ity was observed only with BaG20c at 52 weeks in vivo.Bimodality in MWD was observed when the Mv of thesamples was 3700–2700 Da and Mn 1400–1100 Da. Atthe same time, the crystallinity was approximately 20%and the glass transition temperature (Tg) had decreasedto 51°–53°C. Bimodality was not observed with all thematerials in vivo, presumably because the above-men-tioned levels of Mv or Mn were not reached because ofthe limited number of samples and weeks studied.

Thermal properties

Initially, SR-PLA70 and SR composite samples werecompletely amorphous and a DSC thermogram re-vealed only the Tg of polymer chains at 56°–57°C.

During the hydrolysis, Tg decreased to 51°–53°C. Amelting peak (Tm) at 120°–123°C was detected within6–24 weeks, showing the formation of crystalline res-idues in vitro and in vivo. All the materials showedcold crystallization (peak at 92°–95°C) at 24 weeks invitro and in vivo. This peak disappeared whereas thedegradation proceeded and the overall crystallinityincreased to �15%. Crystallinity of the samples in vitroand in vivo is shown in Table III. After 63 weeks invitro, the crystallinity of SR-PLA70 was highest (41%)and lowest with BaG15s (7%). After 48 weeks in vivo,the crystallinity of BaG15s and BaG20c was equal(21%) and lower compared with SR-PLA70 (59%). Atthe later stage of hydrolysis, the shoulder on the melt-ing peak or double peak was also observed in DSCthermograms around 106°–110°C. The double meltingpeak was more pronounced with the SR composites.By the end of the hydrolysis, the melting peak hadshifted from 120°–123°C to 110°C.

Structure

The SEM examination of the SR-PLA70 specimenbefore hydrolysis had a smooth external surface andfibrillated, nonporous, homogeneous cross-sectionalstructure. Unlike SR-PLA70 samples, SR compositeswere initially internally macroporous with a fibrillatedpolymer structure [(Fig. 5(a,b)]. Spindle-shaped poressurrounded the crushed glass particles or spheres andwere formed during die-drawing. Glass spheres also

Figure 3. GPC chromatograms of SR-PLA70 and BaG15s after 48 weeks in vitro and in vivo.

SELF-REINFORCED BIOABSORBABLE POLYMER 703

formed surface porosity, as has been reported in ourearlier studies.24,25 The surface porosity was not soobvious with crushed glass particles.

The first structural changes due to the degradationof SR-PLA70 were observed after 24 weeks in vitro,when the first small pores were formed underneaththe surface skin of the rod. A progressive loss oforiented, fibrous structure was also detectable. In vivo,the first structural changes and small pores were seenas soon as after 3 weeks. Compared with in vitro, the

microporosity was more substantial after 24 weeks invivo. Multiple micropores occurred beneath the sur-face skin layer in the 48th week in vitro; small poreswere also seen at the core of the rod and no orientedstructure was left (Fig. 6).

With BaG20c composites, the degradation of struc-ture and micropore formation were seen in still fibril-lated and oriented polymer matrix beneath the surface(Fig. 7) after 12 weeks in vitro. After 12 weeks in vivo,microporosity was also observed in the internal partsof the sample (Fig. 8). No micropores close to thesurface were seen in BaG15s samples before 36 weeksin vitro. With BaG15s composites, also, the microporeformation was more generous in vivo. At 63 weeks invitro, crack-like pores, perpendicular to the longitudi-nal axis and oriented structure, were seen throughoutboth SR-composite specimens.

SEM studies also confirmed the appearance of apatiteprecipitation at the surface of both composites, startingon the glass particles and spreading to the polymersurface, already after 1 week’s immersion. As the hydro-lysis proceeded, apatite was also seen in the internalparts of the composite. On the rod surface and also in theinternal parts, apatite precipitation was found to spreadover the polymer matrix close to the glass particles.

DISCUSSION

The DR affects the final strength and modulus of SRbioabsorbable polymer21 and thus the mechanical

Figure 4. Bimodal GPC chromatograms of SR-PLA70 and SR composites of BaG20c and BaG15s after 52, 75, and 87 weeksin vitro, respectively.

TABLE IIICrystallinity Changes of SR-PLA70 and SR Composites

of BaG15s and BaG20c with Degradation In Vitroand In Vivo

Weeks SR-PLA70 BaG15s BaG20c

In vitro0 0 0 03 0 0 06 0 0 1.112 0 0 2.624 3.6 0.4 4.848 12.5 4.4 17.463 41 7.1 17.887 52.4 22.8 30.8

In vivo0 0 0 03 0 0 06 0 0.6 112 1.6 1.4 3.324 5 4.3 6.148 58.7 20.9 21

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properties of SR composites are also related to theDR.25 However, the presence of glass particle fillermodifies the SR fibrillated polymer structure, eachparticle forming a reported spindle-shaped cavityaround the particle along the drawing axis. The glasscontent and the number of macropores formed there-fore greatly affect the mechanical propertiesachieved.25 The initial bending strength, modulus, andshear strength of both BaG15s and BaG20c compositeswere lower compared with SR-PLA70 because of theabove-mentioned macropores and the loss of adhesionbetween the filler and the matrix. With SR composites,the shear strength seemed to be more dependent onthe glass content changes than on changes in DR, sothat the increase in glass amount decreased the shearstrength despite the greater DR of BaG20c. With theself-reinforcing process, the brittleness of ceramic fillercomposites is avoided and the toughness of SR com-posites is comparable to that of SR-PLA70.

Figure 5. (a) Longitudinal section of BaG20c. (b) Longitu-dinal section of BaG15s. Scale bar, 100 �m.

Figure 6. Longitudinal section of SR-PLA70 after 48 weeksin vitro. Scale bar, 100 �m.

Figure 7. Longitudinal section of BaG20c after 12 weeks invitro. Scale bar, 100 �m.

Figure 8. Longitudinal section from the middle of BaG20cafter 12 weeks in vivo. Scale bar, 100 �m.

SELF-REINFORCED BIOABSORBABLE POLYMER 705

In the strength retentions, there was a clear differ-ence in vitro and in vivo. All of the specimens studiedlost their mechanical properties faster in vivo than invitro. When comparing the strength retention of plainSR-PLA70 rods and SR composites, BaG20c degradedfaster and BaG15s slower than SR-PLA70 in vitro andin vivo. Thus, glass clearly affects the degradation rate.The degradation rate of composites seems to dependon the amount of glass and the macroporous structureof the SR composite.

The decrease in buffer solution pH with SR-PLA70indicates that soluble acidic degradation productsformed within the degraded PLA70 matrix were re-leased into the medium. With SR composites, a slightincrease in solution pH at the beginning of hydrolysisis due to the dissolution of the glass and exchange ofNa� or K� ions of glass with H� or H3O� ions of thesolution. This concurs with our earlier in vitro stud-ies,26 in which we found the pH decrease to occur after26 weeks with non-reinforced PLA70, whereas thebuffer solution pH of the non-reinforced PLA70 �BaG composites remained �7 throughout the fol-low-up period (87 weeks). The release of Na�, K�, andCa2� ions from bioactive glass (13–93) neutralizes theacidic degradation products released from the de-graded polymer matrix of composites. Neutralizationmay take place, for example, by formation of the neu-tral salts of carboxylic acids. The neutralized degrada-tion of both non-reinforced and SR PLA70 � BaGcomposites may reduce the risk of sterile sinuses andbone resorption around the implant.

Dimensional stability of the SR-composite rodscompared with SR-PLA70 is one of the advantages.The glass particles may retard the relaxation of the SRstructure, thus helping the composites to retain theinitial dimensions of the implant rods.

At the beginning of hydrolysis, the WA of SR com-posites was more pronounced compared with that ofSR-PLA70. Initially, macroporous composite structureand the presence of glass and matrix interfaces facili-tated the water penetration in the SR composites.Thus, the increase in amount of glass was studied toincrease the WA of SR composites. While the hydro-lysis and degradation proceeded, polymer matrix mi-croporosity was formed, thus further increasing theWA of SR composites. Degradation and the formationof micropores also facilitated the water penetrationinto the SR-PLA70 and a substantial increase in WA ofthe SR-PLA70 rods after 48 weeks in vitro was there-fore observed.

With SR-PLA70, a WL was seen after 48 weeks dueto the changes in a structure, WA, and release of acidicdegradation products. With SR composites, the WLwas more pronounced until 48 weeks. The greater WAat the early stage of immersion probably induced thedegradation of SR composites in terms of WL. Ini-tially, macroporous structure, bioactive glass resorp-

tion, glass neutralization effect and the hydrophilicsalts of carboxylic acids, products of composite deg-radation, may also promote WL in composites.

The molecular weight (Mv) loss was found to befastest at the beginning of the hydrolysis. In terms ofmolecular weight, the plain SR-PLA70 degraded fasterthan the polymer matrix in composites. With all theimplant types studied, degradation in vivo was fasterthan in vitro. The appearance of bimodal MWD wasconsidered to indicate late-stage degradation of poly-mer. In earlier studies, the bimodal MWD of intrinsi-cally amorphous aliphatic polyesters has been re-ported to be a consequence of heterogeneousdegradation caused by the existence of two domainsdegrading at different rates (surface, center)13–15 or thecrystallization of degradation products due to the se-lective degradation of amorphous domains.16 In thisstudy, the bimodality of MWD of the materials isattributed to morphological changes due to the selec-tive degradation of copolymer, preferential degrada-tion in dl-lactic units and crystallization of fragmentsof l-lactic units rather than surface–center differenti-ation and heterogeneous degradation.

Thermal analysis showed that in all cases the crys-tallinity increased over time. The overall crystallinityof BaG15s composites was lower at all times studiedcompared with that of SR-PLA70. At the later stage ofhydrolysis, the crystallinity of BaG20c composites wasalso clearly lower compared with that of SR-PLA70.The lower crystallinity of composites is probably dueto the presence of bioactive glass, which probablyinhibited the recrystallization of the polymer matrix.At the later stage of hydrolysis, a shoulder on themelting peak or a double peak was also observed inDSC thermograms comparable to the MWD bimodal-ity. This confirms that the bimodality might be attrib-uted to the morphological changes and the existenceof different crystallite fragments with crystallites ofdifferent sizes. With composites, the double meltingpeak was more pronounced, indicating that the addi-tion of glass modifies the hydrolytic degradation andthus the morphology of the polymer matrix. The de-gree of crystallinity was observed to be higher in vivothan in vitro. This indicates that polymer degradationwas more pronounced in vivo.

Structural observations indicated that the hydrolyticdegradation was faster near the surface than in themiddle of the rods. At the middle of the rods, nohollow structures were formed in SR-PLA70 or in SRcomposites. Extruded, non-reinforced PLA70 andPLA70 � BaG13 composite rods and thin non-rein-forced PLA70 plates yielded similar results.17,26 How-ever, these results do not correlate with previous re-sults that the degradation of initially amorphousnonfilled polylactides and lactide–glycolide copoly-mers proceeds heterogeneously and is faster in theinner part rather than at the surface. Heterogeneous

706 NIIRANEN ET AL.

degradation was found to cause the hollow center ofthe specimen15,18,19 and if not hollowed out, the inte-rior was cracked and turned to granules because of thecrystallization of degradation products.13,20 The differ-ence in degradation morphology may be related to thepurity and monomer content of the raw material usedand the manufacturing method of the specimens. Themechanical properties, WL, WA, and dimensional sta-bility of all materials studied are in accordance withthe structure of materials and the changes observed,such as pore formation and loss of orientation. Thestructural studies also confirmed the apatite formationof BaG15s and BaG20c composites. Apatite precipita-tion is taken as a sign of potential osteoconductivityand bone-bonding ability of the composites, as indi-cated earlier in the literature.27,28

Compared with earlier studies on bioceramic/bio-absorbable polymer composites, the in vitro degrada-tion rate of the composites is closely related to thematerial combination used. The incorporation of bio-active filler, bioactive glass S53P4, into the poly(�-caprolactone-co-dl-lactide) matrix was seen to acceler-ate degradation.10 The presence of bioactive glassincreased WA, and the average molecular weights ofthe polymer component of the composite sampleswere observed to decrease much faster compared withthe neat polymer sample during the 24 weeks ofstudy.10 However, retarded and homogeneous degra-dation was found with composites of hydroxyapatite/poly-dl-lactide (HA/PDLLA) and coral/poly-dl-lac-tide (coral/PLA50).9,29 The studies with HA/PDLLAcomposites showed slower degradation rate, andlower WA and WL than unfilled PDLLA. pH did notvary during the follow-up period of 12 weeks.9 Coralreduced the degradation rate of PLA50 and degrada-tion of the composite was homogeneous.29 The degra-dation rate of the composites seems to depend amongother things on the macroscopic structure, degrada-tion mechanism of bioabsorbable polymer and theefficiency of bioceramic filler in eliminating autocatal-ysis, and neutralizing the acidic degradation products.

CONCLUSIONS

Bioactive glass 13-93 increased the hydrophilic na-ture of the composites by glass/matrix interfaces andpromoting the macropores in the structure of the com-posites at die-drawing. However, the faster WA of thecomposites did not accelerate the degradation of thepolymer matrix according to the decrease in molecularweight and the increase in crystallinity. On the con-trary, the WL of the composites was more pronouncedin the early stage of the hydrolysis compared withSR-PLA70 which may indicate erosion-like degrada-tion of the composites. Erosion-like degradation has

also been reported with HA and MgO-filled l/d co-polymer films.30 On the basis of the data obtainedfrom GPC, DSC, and the more pronounced bimodal-ity, the glass modified the degradation kinetics andthe material morphology (crystallinity) during thedegradation of the composites. The hydrolytic degra-dation of BaG15s and BaG20c composites was foundto depend on their macroscopical structure, theamount of glass particles, and thus the appearanceand degree of the initial macroporosity. It was alsoobserved that the addition of glass was beneficial interms of dimensional stability, neutral degradation,and apatite formation of the composite rods.

The degradation was more pronounced in vivo thanin vitro both with SR-PLA70 and SR-PLA70 � BaGcomposites. This indicated that the soft tissue environ-ment (enzymes, implant movement) contributes to thedegradation process. Further in vivo studies regardingbone implantation are in progress.

We suggest that the bioabsorbable polymers con-taining osteoconductive filler may have great potentialas bone-repairing materials. The porous structure, in-creased hydrophilicity, nonacidic degradation, andenhanced osteoconductivity of the SR composites mayprovide a favorable substrate for osteoblast adhesion,bone tissue formation, and ingrowth. Surface porositymay also be beneficial for creating the interlockingmechanism, thus helping the permanent stabilizationof the implant.

The authors thank Ms. Anja Marttila for assistance withthe in vitro studies and testing of the in vivo implants.

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