vidros empleados como biomateriales
TRANSCRIPT
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Bioresorbable and bioactive polymer/Bioglass1 composites withtailored pore structure for tissue engineering applications
Aldo R. Boccaccinia,*, Veronique Maquetb
aDepartment of Materials and Centre for Tissue Engineering and Regenerative Medicine Imperial College London, London SW7 2BP, UK bCentre for Education and Research on Macromolecules (CERM), Interfacultary Centre for Biomaterials, University of Liege,
B-4000 Lie ge, Belgium
Received 17 December 2002; accepted 16 May 2003
Abstract
An overview about the development of porous bioresorbable composite materials for applications as scaffolds in tissue engi-
neering is presented. A thermally induced phase separation method was developed to fabricate porous foam-like structures of
poly(lactide-co-glycolide) (PLGA) containing bioactive glass particle additions (up to 50 wt.%) and exhibiting well-defined, orien-
ted and interconnected porosity. The in vitro bioactivity and the degradability of the composite foams were investigated in contact
with phosphate buffer saline (PBS). Weight loss, water absorption and molecular weight measurements were used to monitor the
polymer degradation after incubation periods of up to 7 weeks in PBS. It was found that the presence of bioactive glass retards the
polymer degradation rate for the time period investigated. The present results show a way of controlling the in vitro degradation
behaviour of PLGA porous composite scaffolds by tailoring the concentration of bioactive glass.
# 2003 Elsevier Ltd. All rights reserved.
Keywords: Bioresorbable polymers
1. Introduction
Tissue engineering presents an alternative approach to
the repair and regeneration of damaged human tissue,
avoiding the need for a permanent implant. The under-
lying principle involved is the regeneration of living tis-
sue, where a loss or damage has occurred as a result of
injury or disease [1, 2]. Tissue engineering can be there-
fore simply defined as the ‘‘science of persuading the
body to heal by its intrinsic repair mechanisms’’ [3]. The
scientific challenge encompasses understanding the cells
themselves, their mass transport requirements and bio-
logical environment as well as the development of sui-table scaffold materials, usually porous, that act as
templates for cell adhesion, growth and proliferation.
The ultimate goal is to return full biological and
mechanical functionality to a damaged tissue or organ.
Several physicochemical and biological requirements
have to be fulfilled by the scaffold, depending on
the particular tissue under consideration, which are
closely dependent on the scaffold porosity and porous
structure.
In a recent review paper, Hutmacher has summarized
the requirements for scaffolds intended for musculoske-
letal tissue engineering [4]. Ideally, scaffolds for this
application should have the following characteristics: (i)
three dimensional and highly porous with an inter-
connected pore network for cell growth and flow trans-
port of nutrients and metabolic waste, (ii) biocompatible
and bioresorbable with a controllable degradation and
resorption rate to match cell/tissue growth in vitro and/
or in vivo, (iii) suitable surface chemistry for cell
attachment, proliferation, and differentiation and (iv)mechanical properties to match those of the tissues at
the site of implantation. In particular for bone tissue
applications, it has also been reported [5], that a suitable
scaffold should bond to the host tissue without the for-
mation of scar tissue, i.e. it should exhibit bioactivity
and osteoconductivity. Moreover a suitable scaffold
should be made from versatile processing techniques
that can produce irregular, usually complex, shapes to
match that of the defect in the tissue of the patient.
A wide variety of both natural and synthetic materi-
als, or a combination of them, are being investigated for
0266-3538/$ - see front matter # 2003 Elsevier Ltd. All rights reserved.
doi:10.1016/S0266-3538(03)00275-6
Composites Science and Technology 63 (2003) 2417–2429
www.elsevier.com/locate/compscitech
* Corresponding author. Tel.: +44-207-5946731; fax: +44-207-
5843194.
E-mail address: [email protected] (A.R. Boccaccini).
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design and construction of scaffold for tissue engineering.
These include naturally occurring polymers, e.g. hydro-
gels like gelatin, agar, fibrin or collagen [6–9], synthetic
bioresorbable polymers, e.g. poly(lactide acid), poly(-
glycolic acid), polycaprolactone and poly(propylene
fumarate) [3,10–15] as well as bioactive porous cera-
mics, e.g. synthetic foam-like bioactive glass and cal-cium phosphate structures [16–20] and naturally
occurring ceramics, such as coral [21]. Also metallic
foams are being considered for tissue engineering
scaffolds [22].
Bioresorbable synthetic polymers have attracted
increasing attention for their use as tissue engineering
scaffolds in the last ten years [3,10–15,23]. Many prac-
tical advantages arise when using synthetic scaffolds
because precise control of material composition and
micro- and macrostructure, including porosity, is possi-
ble. This allows adequate control of scaffold properties,
thus creating optimal conditions for cell survival, pro-
liferation, and subsequent tissue formation. Polyesterssuch as poly(lactic acid) (PLA), poly(glycolic acid)
(PGA) and poly(lactic acid-co-glycolic acid) (PLGA)
are being mainly considered for scaffold applications
[3,10–15,23]. These polymers have already demon-
strated promising results in clinical use, for example as
resorbable surgical sutures and meshes or in drug deliv-
ery systems, and they have the United States Federal
Food and Drug Administration (F.D.A.) approval for
clinical use. The specific use of synthetic bioresorbable
polymers in tissue engineering applications has been
reviewed by Hutmacher [4], Maquet and Jerome [24]
and Agrawal et al. [23].As mentioned above, porosity and pore structure are
key parameters determining the properties and the
applicability of scaffolds for tissue engineering. Fig. 1
shows a summary of the different functions related to
the pore structure in a tissue engineering scaffold. In
general, scaffold porosity, pore morphology and pore
orientation must be tailored to the particular tissue
under consideration.
Porous polymeric tissue engineering scaffolds of
3-dimensional (3-D) structure have been prepared by
numerous techniques, including solid–liquid and liquid–
liquid phase separation [25–31], solution casting [32], gel
casting [33], gas saturation [34], gas foaming [35,36],
fibre bonding, fabric forming and related textile based
processes [37–40], sintering of polymeric microspheres
[41], combined solvent casting and extrusion [42], emul-
sion freeze-drying [43], several rapid prototyping tech-niques, e.g. 3-D printing and fused deposition modelling
[4,6,44–48] and various solvent casting/particulate
leaching methods using different porogen additives [49–
55]. Overviews on fabrication methods of 3-D porous
polymeric scaffolds for tissue engineering are given by
Hutmacher [4], Agrawal et al. [23], Yang et al. [56] and
Thomson et al. [57]. In recent studies, advanced meth-
ods have been developed for the optimal designed of
porous scaffolds based on solid free form manufactur-
ing and conventional sponge scaffold fabrication [58].
Selected ceramics, such as hydroxyapatite (HA), tri-
calcium phosphate (TCP) and some compositions of
silicate and phosphate glasses and glass–ceramics, reactwith physiological fluids and form tenacious bonds to
hard and soft tissues through cellular activity [2,5].
These materials are therefore known as ‘‘bioactive’’ [59].
If biodegradability and bioactivity are to be combined
in an optimised tissue engineering scaffold for bone tis-
sue engineering, then the design of composite materials
offers an exceptional opportunity: by combining bior-
esorbable polymers and bioactive ceramic phases scaf-
folds with tailored physical, biological and mechanical
properties can be produced. Moreover the addition of a
ceramic or glass phase to a biodegradable polymer may
be exploited to favourable alter the in vitro and in vivopolymer degradation behaviour, which, in polylactides,
is strongly acidic. It has been proposed that bioactive
glass particles used as inclusions or coatings in biode-
gradable polylactides should lead to the rapid exchange
of protons in water for alkali in the glass which should
then provide a pH buffering effect at the polymer sur-
face, thus preventing acceleration of polymer degrada-
tion [60]. Moreover, as discussed elsewhere [61],
dissolution of the bioactive glass should result in
nucleation and growth of a crystalline HA layer on the
surface of the polymer scaffold, which should further
affect the polymer degradation behaviour in addition to
provide the required osteoconductivity. Thus, a combi-nation of porous bioresorbable polymers and bioactive
ceramic or glass phases is expected, for the following
reasons, to result in promising composite scaffolds for
(bone) tissue engineering [62]: (i) a better cell seeding
and growth environment can be achieved because of the
good osteoconductivity properties provided by the
bioactive phase, (ii) the acidic degradation by-products
from polyesters may be buffered, and (iii) the mechan-
ical properties may be improved by using the traditional
composite approach (inclusion of a stiffer particulate
ceramic phase in the polymeric matrix).Fig. 1. Schematic diagram showing the different functions of a tissue
engineering scaffold depending on its porosity and pore structure.
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In a broader sense, the microstructure of composite
materials, including porosity and pore structure, can be
engineered in such a way that the resorption rate of the
composite scaffold in the body can be designed to match
the formation rate of new tissue. In this regard, also
functionally graded materials with variable porosity and
graded bioactive phase content have been developed[63] in order to mimic the defect site in terms of biolo-
gical, mechanical and morphological properties.
A review covering the fabrication, properties and
applications of bioresorbable and bioactive porous
composites for bone tissue engineering scaffolds have
been presented elsewhere [64]. In the present paper, new
developments based on PLGA foams with tailored
porosity and bioactive glass particle inclusions are pre-
sented. Both the fabrication and in vitro characterisa-
tion of the foam composites are described in detail. Two
bioresorbable copolymers with different lactide/glyco-
lide ratios were investigated. This study complements
our recent work which was mainly based on poly(D,L-lactide) (PDLLA) foams and composites [61,65].
2. Experimental
2.1. Composites processing
Poly(lactide-co-glycolide) (PLGA) copolymers with
two different lactide:glycolide (LA:GA) ratios, i.e.
LA:GA=75:25 (i.v.=0.6) and LA:GA=50:50
(i.v.=0.2), Resomer RG 756 and 502, respectively, were
provided by Boehringer-Ingelheim (Germany). Thesecopolymers will be designated as PLGA75 and 50,
respectively. Dimethylcarbonate (DMC) of 99% purity
(Sigma Aldrich) was used as a solvent. Polymers and
solvent were used without further purification. The
bioactive material used was a bioactive glass powder
(Bioglass1 grade 45S5, US Biomaterials Co., Alachua,
FL, USA). The powder had a mean particle size <5
mm. The composition of the bioactive glass used was (in
weight percentage): 45% SiO2, 24.5% Na2O, 24.5%
CaO and 6% P2O5, which is the original composition of
the first bioactive glass developed by Hench and co-
workers [59].
A 50/50 mixture (wt./wt.) of PLGA75 and PLGA50was used as the polymer component. The Bioglass1
content was varied from 10 to 50 wt.%. For the sake of
comparison, neat (unfilled) polymer foams were also
prepared and characterised.
The preparation of Bioglass1-filled polymer foams,
referred here as composite foams, followed a thermally
induced phase separation process, also termed freeze-
drying, which has been described in detail elsewhere
[30,31]. The original process was conveniently modified
for the incorporation of Bioglass1 particles as follows.
The polymer was dissolved in DMC to produce a poly-
mer weight to solvent volume ratio of 5%. The mixture
was stirred overnight to obtain a homogeneous polymer
solution. Determined amounts of Bioglass1 powder,
calculated to result in final proportions of 10, 25 and 50
wt.% of Bioglass1 in the composites, were added into
the polymer solution, resulting in homogeneous poly-
mer-Bioglass
1
particles mixtures. Each mixture wasthen transferred into a 600 ml lyophilisation flask and
sonicated for 15 min in order to improve the dispersion
of the Bioglass1 particles into the polymer solution.
The flask was then rapidly immersed into liquid nitro-
gen and was maintained at 196 C for 2 h. The frozen
mixture was then transferred into an ethylenglycol bath
at 10 C and connected to a vacuum pump (102 Torr)
for solvent sublimation (at 10 C for 48 h, and then at
0 C for additional 48 h). The foam samples were sub-
sequently completely dried at room temperature in a
vacuum oven until reaching a constant weight, as
determined by using an electronic balance.
2.2. Characterisation and in vitro studies
Neat polymer foams and polymer/Bioglass1 compo-
site samples were characterised by using scanning elec-
tron microscopy (SEM) in order to assess the porosity
structure. The apparent density of the foams (ra) was
determined by mercury pycnometry measurements as
follows: a sample of weight ws was placed in a pycn-
ometer, which was completely filled with mercury and
weighted to obtain wsl, then a was calculated according
to Eq. (1):
a ¼ ws
wl wsl þ ws
Hg ð1Þ
where wl is the weight of the pycnometer filled with
mercury, and Hg is the density of mercury (13.5 g
cm3). The density of the solid, non-porous polymer/
Bioglass1 skeleton, sk, was measured by helium pyc-
nometry using a AccuPyc 1330 pycnometer (Micro-
metrics Co.). The porosity of the foam (e) was
calculated according to Eq. (2):
" ¼sk a
sk100% ð2Þ
The compressive mechanical properties of the com-posite foams were measured with a rheometer (Ares,
Rheometric Scientific). Specimens of 10 mm diameter
and 4–5 mm height were compressed at a cross-head speed
of 1 mm/min. The compressive modulus (F) was deter-
mined from the initial linear region of the normal force
versus compression strain plot. The compressive mod-
ulus in Pa was calculated according to Eq. (3), where m
is the slope of the initial linear region of the curve:
F ¼m 104
16 ð3Þ
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For degradation studies in normal physiological con-
ditions, selected samples of polymer foams and compo-
sites were sterilised by UV exposure under a laminar
flow for 10 min and placed in sterile Falcon tubes con-
taining 50 ml of pre-filtered (0.22 mm porosity) phos-
phate buffer saline (PBS: 0. 13M, NaCl: 0.9%, NaN3:
0.02%, pH: 7.4). The samples were incubated underslow tangential agitation at 37 C and allowed to
degrade. The pH of the buffer was monitored during the
experiment. At each time point, 3–4 samples of each
foam composition were removed from the buffer, and
weighted wet after surface wiping. They were abun-
dantly rinsed with deionized distilled water (ddH2O) in
order to remove the soluble inorganic salt, and weighted
after freeze-drying. From the weight measurements,
water absorption (WA%) and weight loss (WL%) were
calculated according to Eqs. (4) and (5), respectively:
WA ¼ W
a W
oð Þ
Wo100% ð4Þ
WL ¼ Wo Wtð Þ
Wo
100% ð5Þ
where: Wo, Wa, and Wt are the samples’ weights before
immersion, after removal from the buffer and after
freeze-drying, respectively. Three to four samples of
each composition were measured and the results aver-
aged.
The changes in average molecular weight (Mw) were
determined by size exclusion chromatography (SEC)using a Helwett-Packard HP-1090 SEC apparatus
equipped with three Ultrastyragel columns (102 –105A ˚ ).
Tetrahydrofuran was used as an eluent (flow rate: 1 ml/
min) and calibration was performed using monodisperse
polystyrene standards (Polymer Laboratories Ltd.,
Shropshire, UK). After degradation for different time
periods in PBS, selected samples were characterised
using SEM, X-ray diffraction (XRD) analysis and
Raman spectroscopy.
Information on the elementary composition of Bio-
glass1 particles at the surface of the composite scaffolds
was obtained using environmental scanning electron
microscopy (ESEM) (Philips FEG XL-30) combinedwith energy dispersive X-ray analysis (EDXA). EDXA
was carried out to determine the Ca/P ratio of the Bio-
glass1 particles prior and during in vitro degradation
Fig. 2. SEM micrographs showing cross-sections of (a, b) an unfilled PLGA foam at two different magnifications; (c) a 100/25 PLGA/Bioglass1
composite at low magnification and (d) a 100/50 PLGA/Bioglass1 composite at high magnification.
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and to monitor HA formation on the composite sur-
faces. The mean Ca/P ratios were determined from five
separate measurements in different areas of the compo-
site samples.
3. Results
Highly porous PLGA/Bioglass1 composites foams
have been prepared by solid-liquid phase separation and
subsequent solvent sublimation, as summarised inTable 1. The density of the composite increases with
Bioglass1 content. In parallel, the porosity decreases
when Bioglass1 content is increased. The porosity of
the composite foams was in general very high ( 90%)
even for the higher Bioglass1 content (50 wt.%). Figs. 2
(a–d) show SEM micrographs of unfilled and composite
foams. Unfilled PLGA foams are characterised by an
open, well-structured porous network composed of
interconnected tube-like macropores parallel to each
others and oriented along the heat-transfer direction
imposed by the unidirectional cooling process (Fig. 2a).
At a higher magnification (Fig. 2b) it is possible toobserve that the foam exhibits a ladder like substructure
(parallel microtubes with thin partitions). This kind of
pore architecture is typical of PLA and PLGA foams
prepared by uniaxial thermally induced phase separa-
tion process in solvents such as dioxane, benzene and
dimethylcarbonate at suitable polymer concentrations
(i.e. 5 wt.:v%) [24,30,31].
The pore morphology of the foams filled with Bio-
glass1 particles is slightly different to that of pure
PLGA foams. A typical SEM micrograph of a 100/25
PLGA/Bioglass1 composite foam shows a co-con-
tinuous structure of interconnected irregular pores of
size ranging from 10 to 100 mm (Fig. 2c). Bioglass1
particles can be easily identified on the polymer matrix
surface, more often assembled into aggregates. The dis-
persion was found to be more homogeneous for high
Bioglass1 contents, e.g. 100/50 PLGA/Bioglass1, as
seen in Fig. 2d. A qualitative good adhesion was foundbetween the PLGA matrix and Bioglass1 particles,
however no quantitative information about the inter-
facial bonding strength between polymer and Bioglass1
was obtained in this study.
Fig. 3. Compression modulus of the PLGA foam and 100/50 PLGA/
Bioglass1 composite foam.
Table 1
Density and porosity of PLGA/Bioglass1 composite foams
Composition(wt.%) Apparent density (g/cm3) Porosity (%)
PLGA/Bioglass1: 10 0/0 0 .073 0.001 94.3 0.4
PLGA/Bioglass1: 100/10 0.098 0.007 91.5 0.9
PLGA/Bioglass1: 100/25 0.124 0.024 90.4 1.8
PLGA/Bioglass1
: 100/50 0.130 0.020 89.9 0.4
Fig. 4. (a) Water absorption (WA), (b) weight loss (WL), and (c)
changes in pH of the incubation medium (PBS) versus incubation time
in PBS for PLGA/Bioglass1 composite foams with different PLGA/
Bioglass1 weight ratios: 100/0 (), 100/10(*), 100/25 (~), 100/50
(^).
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Fig. 3 shows that the mechanical properties of the
PLGA/Bioglass1 composites in term of compression
modulus are significantly influenced by the incor-
poration of Bioglass1 particles: the compression mod-
ulus of the 100/50 PLGA/Bioglass1 composite foams is
three times higher than that of the pure PLGA foams.
Such improvement of the elastic modulus has been oftenobserved in degradable polymer composites reinforced
by HA and other ceramic particles [64], confirming the
expected effect of adding a rigid, inorganic particulate
phase (Bioglass1) to the polymer matrix. The mechan-
ical properties of similar foams with different polymer
matrices (PDLLA, PLGA) and increasing content of
Bioglass1 have been measured by compression test and
discussed in a previous study [66]. It was confirmed that
the compression modulus of the composite scaffolds can
be significantly enhanced by addition of Bioglass1,
especially for amorphous and though poly(D,L-lactide)
and poly(D,L-lactide-co-glycolide).
For in vitro degradation studies, the PLGA/Bio-
glass1 composites were incubated into phosphate buffer
saline in normal physiological conditions (at pH=7.4
and 37
C). Fig. 4a shows that the PLGA/Bioglass
1
composites absorbed a high amount of water during the
first week of incubation. However WA reached an
equilibrium value and started to decrease after 21 days
in all composites. In the neat PLGA foam, WA was
lower than in the composites and increased slightly over
the entire incubation period.
According to weight loss (WL) vs incubation time
plots (Fig. 4b), WL in the composites increased during
the whole incubation period and proportionally to the
Bioglass1 content. The weight of the composites con-
taining 525 wt.% Bioglass1 decreased largely during
the first week of incubation (WL = 15%) and con-
tinued to decrease proportionally to the Bioglass1 con-tent. At the end of the incubation period (35 days), WL
was around 12, 16, and 25% in the composites contain-
ing 10, 25 and 50 wt.% of Bioglass1, respectively. The
weight of the neat PLGA foams did not change to a
large extent during the whole incubation period
(WL47%).
The pH variation patterns of the media containing the
composite foams are shown in Fig. 4c. The pH of the
incubation medium decreased from the initial value
Fig. 5. SEC chromatograms for 50–50 PLGA75–PLGA50 mixtures:
(a) not incubated, (b) after 35 days of incubation in PBS.
Fig. 6. Changes in molecular weight versus incubation time in PBS for unfilled and Bioglass1 filled PLGA75 and 50 foam samples of the indicated
compositions.
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(7.4) in the unfilled PLGA foams and in the composite
filled with 10 wt.% Bioglass1 after the first week of
incubation. On the contrary, the pH of the medium
slightly increased to values 57.4 for the 25 wt% and 50
wt.% Bioglass1 composites. Such difference may be
correlated to dissolution of alkaline ions from the Bio-
glass
1
particles that locally compensate for the acid-ification of the medium due to acidic products of the
polymer degradation. Such a buffering effect of Bio-
glass1 was reported elsewhere [60], and it has been
considered to be another benefit of using Bioglass1
particles in composite scaffolds with the aim to avoid
possible inflammatory response due to acidic degrada-
tion of the polymers.
The SEC chromatogram for a 50/50 mixture of
PLGA75 and 50 shows the presence of two peaks that
overlap each other, as shown in Fig. 5a. For this reason
the value of Mw at the maximum of the distribution
peak that represents the most significant fraction of the
polymer chains present in the samples has been con-sidered. This Mw value at the top of the peak has then
be used to follow the evolution of the molar mass of the
two copolymers during the incubation period. Fig. 6
shows the evolution of the Mw of both PLGA75 and
50, as measured by SEC, by using Mw at the top of each
peak as the indicative value for polymer degradation. At
the end of the incubation time, the peak for PLGA75
appears like a shoulder in the larger peak for PLGA50
and the SEC program gave only one value of Mw at thetop of the peak which was assigned to PLGA50, as
shown in Fig. 5b.
Fig. 6 shows that the Mw of PLGA75 rapidly
decreased during the first three weeks of incubation for
all Bioglass1 particle contents, but with a slower rate
for the composites containing 50 wt.% Bioglass1 as
compared to the unfilled PLGA foam. At day 21, the
Mw of PLGA75 was higher in the 100/50 PLGA/Bio-
glass1 composite in comparison to the others. The Mw
of PLGA50 did not change during the first three weeks
of incubation. For longer incubation times, the Mw of
PLGA50 started to slightly decrease. The results repor-
ted here seem to indicate that the presence of Bioglass1
particles retards the degradation rate of the polymer, at
least for PLGA75. The phenomenon was enhanced for
Fig. 7. SEM micrographs of PLGA/Bioglass1 (100/50) composite samples after incubation time of three weeks in PBS. View of the surface of the
composite sample at (a) low and (b) high magnification.
Fig. 8. SEM micrographs of PLGA/Bioglass1 (100/25) composite samples after five weeks incubation time in PBS. View of a longitudinal section of
the composite sample at (a) low and (b) high magnification.
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longer degradation times, in agreement with previous
results on poly(D,L-lactide) (PDLLA)/ Bioglass1 foams
[66].
The in vitro bioactivity of the samples was assessed by
investigating the formation of HA on their surfaces
during immersion in PBS under normal physiological
conditions. After three weeks, a large number of micro-particles with diameter 1–2 mm was formed on the
surfaces of the pore walls, as shown in Fig. 7 for a 100/
50 PLGA/Bioglass1 composite. These particles were
assembled to form larger aggregates (5 mm). Their mor-
phology is typical of carbonated HA, which has been
also observed to form on Bioglass1 coated PDLLA
foams after incubating in simulated body fluid (SBF)
[61]. HA crystals were also formed on the interior of
foams as observed on longitudinal sections of composite
samples after five weeks on incubation in PBS, as shownin Fig. 8. The formation of HA on the surfaces of com-
posite foams after immersion in PBS was confirmed by
ESEM observations coupled with EDX analysis. Fig. 9
Fig. 9. (a) ESEM micrograph and (b) EDX analysis of a PLGA/Bioglass1 (100/50) composite sample after incubation time of three weeks in PBS,
confirming HA formation.
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shows the morphology of a 100/50 PLGA/Bioglass1
composite after three weeks of immersion in PBS. HA
particles appear on the surface of the polymer foam,
and high Ca and P peaks belonging to HA were detec-
ted by EDX analysis (Fig. 9b). EDS analyses carried out
on different regions of the same sample showed a var-
iation in the Ca and P peak heights, which are related to
different thickness of the HA clusters formed. However,
the average Ca/P ratio of a given sample was overall
remarkably constant. The Ca/P ratio of the as-receivedBioglass1 particles was around 5. After incubation in
PBS for one week, this ratio decreased to 1.3 to reach a
value of 2 after seven weeks of incubation. This Ca/P
ratio is close to that of carbonated HA which confirms
the transformation of Bioglass1 into HA following dis-
solution mechanisms as reported in the literature [59].
The crystallinity of HA formed on the surface of
PLGA/Bioglass1 composites was confirmed by XRD,
as shown in Fig. 10 for a 50 wt.% Bioglass1 composite
sample incubated for seven and 35 days in PBS. In the
same figure the XRD pattern of the unfilled PLGA
foam is shown for comparison The HA peaks can
be seen at diffraction angle 2-theta=32
. Finally,Raman spectroscopy was used to provide an indepen-
dent confirmation of HA formation. Typical spectra are
shown in Fig. 11. A more detailed study on the structure
and further characteristics of the HA formed and its
effect on composite biocompatibility, e.g. using osteo-
blast cell culture experiments, is the focus of current
investigations.
4. Discussion
As mentioned in the Introduction several methods
have been reported in the literature for production of bioresorbable porous scaffolds for tissue engineering
applications. All the methods exhibit relative advan-
tages and disadvantages, depending on the material
used and porous microstructure required. Thermally
induced phase separation (TIPS) offers, in particular,
many advantages [24]: (i) optimal control over pore
volume fraction, (ii) possibility of designing pore shape,
orientation and size, (iii) amenable to be applied to any
polymer soluble in a suitable solvent, (iii) reproduci-
bility, and (iv) possibility to incorporate bioactive sub-
stances or growth factors into the polymer matrix. In
Fig. 10. XRD diagrams of PLGA/Bioglass1 composite samples (100/50) incubated for (a) 7 and (b) 35 days in PBS showing the development of
crystalline HA. The pattern of the neat PLGA foam (c) is also shown for comparison.
Fig. 11. Raman spectra of PLGA/Bioglass1 composite samples (100/
50) as function of incubation time in PBS for (a) 0 days, (b) 7 days, (c)
35 days. The formation of HA is indicated by the P-O peaks at 962
cm1.
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the present experiments using TIPS, ultrasonication of
the polymer/Bioglass1 mixtures before freezing of the
polymer solution allows for an optimal dispersion of the
Bioglass1 particles throughout the polymer matrix. The
composition of the foams was controlled by changing
the polymer/Bioglass1 weight ratio while foam density
was mainly governed by the concentration of the poly-mer solution. In this study, dimethylcarbonate was
advantageously used as solvent instead of the more fre-
quently used dioxane, which is suspected to be carcino-
gen for humans [67].
Foams produced by the TIPS process are char-
acterised by a high porosity (>90%), comprising two
distinct pores sizes: i) macropores of average diameter
> 100 mm, and ii) micropores with an average dia-
meter of 20–30 mm, which form an interconnected net-
work. The tubular macropores are seen to be highly
oriented and parallel to each other as a result of the
unidirectional cooling process.
It is well known that for bone regeneration the idealpore size of the scaffold should approximate that of
bone [68]. In particular, scaffold macroporosity in the
range 100–250 mm of mean pore diameter is ideal for
bone-cell colonization. This is typically the range of
macroporosity that can be generated using the TIPS
process (freeze-drying). In previous studies the influence
of the processing parameters (cooling rate) and for-
mulation conditions (nature of the solvent, polymer
concentration) on pore size distribution has been pre-
sented [24,30,31]. These parameters can be conveniently
chosen for tailoring the scaffold porosity.
Pore interconnection is another relevant property thatmust be achieved in an optimised scaffold in order to
supply body fluid circulation [68]. In foams produced by
TIPS, the walls of the tubular macropores are porous
and open forming a highly interconnected microporous
network (see Fig. 2a–d). The macropores are thus con-
nected to each other through the microporous network.
This desired high pore interconnectivity achieved by the
TIPS method is not possible to be obtained by other
techniques, for example salt-leaching or gas foaming
methods. In previous studies, the transport properties of
3D porous foams made by TIPS, which are in close
relation to the pore interconnectivity, have been inves-
tigated by impedance spectroscopy, which relies uponthe measurement of the ionic conduction of water-satu-
rated foams [69,70].
Analysis of the available literature has indicated that
the application of bioactive glasses in resorbable bio-
composites for tissue engineering scaffolds is not as
widespread as that involving synthetic HA or related
calcium phosphate ceramics [64]. However, it is well
known that bioactive glasses possess a higher index of
bioactivity than hydroxyapatite [59]. In particular, the
glass used here, 45S5 Bioglass1, exhibits ‘‘Class A’’
bioactivity, i.e. it shows properties of osteoconduction
and osteoproduction, while HA exhibits ‘‘Class B’’
bioactivity, i.e. only osteoconductive behaviour [5,17,
59]. Thus, bioactive glass provides an ideal environment
for colonization, proliferation and differentiation of
osteoblasts to form new bone that has a mechanically
strong bond to the implant surface [5,59]. It has been
recently shown that this response is genetically con-trolled, with seven families of genes up regulated within
48 hours of the exposure of primary human osteoblasts
to the ionic dissolution products of bioactive glasses
[71]. Some compositions of bioactive glasses also show
strong interaction with soft tissues [2]. Pilot studies in
our laboratory (not published) have shown that rat
fibroblasts cultured with 45S5 Bioglass1 secrete
increased amounts of VEGF. In addition, in recent in
vivo studies, loosely woven meshes of polyglycolide acid
fibres (Dexon1 mesh) coated with 45S5 Bioglass1 have
been placed subcutaneously into rats leading to
increased vascularization compared with uncoated
meshes [72, 73].Thus, the incorporation of Bioglass1 into resorbable
polymer scaffolds is seen as a convenient way towards
tissue engineering scaffolds for both hard and soft tissue
applications. The principle of incorporation of Bio-
glass1 in a biodegradable polymer for tissue engineer-
ing scaffolds has been patented [74]. In our own
previous research, Bioglass1 particles have been applied
both as coating and fillers in PDLLA foams for
enhanced bioactivity [61,65,66]. Osteoblast cell culturing
studies on Bioglass1 coated PDLLA foams have
demonstrated the positive effect of the bioactive coating
in promoting cell adhesion and spreading even afteronly few minutes in culture [65]. The degradation beha-
viour of PDLLA foams filled with Bioglass1 particles
has been studied by immersion in PBS [66]. Since the
conditions of the test were the same as those of the
present study on PLGA foams, the results are compar-
able and the relative effect of Bioglass1 additions on
degradation of PDLLA and PLGA can be assessed.
The results of the degradation studies in PBS for the
PLGA/Bioglass1 composites presented here show that
the introduction of a bioactive filler in the polymer
foams increases their capacity to absorb water during
the initial incubation period (Fig. 4a). An equilibrium in
WA was rapidly reached, i.e. for three weeks on incu-bation, after which WA started to decrease. However
this decrease in WA can be correlated to a significant
weight loss, especially for composites with a high Bio-
glass1 content. Comparing with the results collected on
PDLLA/Bioglass1 composites treated in the same way
[66], it becomes evident that weight loss was much
higher in the PLGA composites than in the PDLLA
ones, all the other conditions being the same (Bioglass1
content and time of incubation). This means that
degradation was higher for PLGA composites, which is
also confirmed by the rapid decrease of the PLGA75
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molecular weight during the 35 days of incubation
(Fig. 6). A quantitative study on the mechanisms of
degradation of both polymers and on the effect of Bio-
glass1 on their degradation rate is however beyond the
scope of the present work. Nevertheless the experi-
mental results presented here may serve as a suitable
and reliable data base for verification of models and forcomparison of the behaviour of different polymer/Bio-
glass1 systems.
The presence of Bioglass1 particles on the pore walls
both on the outer and internal surfaces of the foams
(Fig. 2c, d). may encourage both bone and soft tissue in-
growth from the implant/tissue interface to the interior
of the scaffold. The bioactivity of the composite scaf-
folds, determined by the rapid formation of carbonated
HA crystals on the sample surfaces during immersion in
PBS, was confirmed by electron microscopy, XRD
analyses and Raman spectroscopy.
It is envisaged that using Bioglass1 particles both as
filler and coatings in porous resorbable polymer scaf-folds will add to the possibilities of tailoring the
mechanical properties and the rate of in vivo resorption
of the composite scaffolds for the required application.
Further logical steps to the optimisation of scaffolds for
tissue engineering should focus on tailoring the micro-
structure of the composite foams for determined appli-
cations, including the development of graded porosity
and graded bioactive glass coatings, this being the focus
of current developments.
5. Conclusions
The present challenge for the progress of tissue engi-
neering is to design and fabricate reproducible bior-
esorbable 3-dimensional scaffolds, which are able to
function for a certain period of time in the body also
under load-bearing conditions. In particular for bone
and cartilage tissue engineering, numerous porous bior-
esorbable and bioactive composite systems are being
currently considered, mainly based on synthetic bior-
esorbable polymers and bioactive glass or calcium
phosphate ceramic phases. Current research and devel-
opment activities focus on choosing the adequate com-
ponents of the composites and on optimising theprocessing routes to produce the required composite
architecture, including porosity, for the envisaged
application. The composites presented in this paper
exhibit a very attractive combination of bioresorption
and bioactivity attributes in a highly porous PLGA
scaffold. They are therefore potential materials for fab-
rication of hard and soft tissue engineering scaffolds. In
particular, the use Bioglass1 instead of HA or other
calcium phosphates yields composites of class A bioac-
tivity with the possibility to be used also in soft tissue
engineering applications. To the authors‘ knowledge,
this is the first experimental report assessing the effects
of Bioglass1 particles addition on the in vitro degrada-
tion of PLGA. Future research should focus on tailoring
novel microstructures for the envisaged applications,
including the development of composite foams with
graded porosity and ‘‘engineered’’ Bioglass1 coating
microstructures.
Acknowledgements
Helpful discussions with Professor Larry L. Hench
(Imperial College London) are appreciated. The authors
acknowledge experimental assistance of I. Notingher, J.
A. Roether and J. Blaker (Imperial College London)
and of L. Pravata (University of Liege, Belgium). VM is
‘‘Postdoctoral Researcher’’ by the ‘‘Fonds National de
la Recherche Scientifique’’ (F.N.R.S). CERM is
indebted to the ‘‘Services Fe ´ de ´ raux des Affaires Scienti-
fiques, Techniques et Culturelles’’ for financial supportin the frame of the ‘‘Poles d’Attraction Inter-
universitaires: PAI 4/11’’.
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