45s5 bioglasss-derived glass ceramic scaffolds

8/13/2019 45S5 Bioglasss-Derived Glass Ceramic Scaffolds

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Biomaterials 27 (2006) 24142425

45S5 Bioglasss-derived glassceramic scaffolds for

bone tissue engineering

Qizhi Z. Chena, Ian D. Thompsonb, Aldo R. Boccaccinia,

aDepartment of Materials and Centre for Tissue Engineering and Regenerative Medicine, Imperial College London,

Prince Consort Road, London SW7 2BP, UKbOral & Maxillofacial Surgery, GKT Dental Institute, Kings College London, London SE1 9RT, UK

Received 12 August 2005; accepted 9 November 2005

Available online 5 December 2005

Abstract

Three-dimensional (3D), highly porous, mechanically competent, bioactive and biodegradable scaffolds have been fabricated for the

first time by the replication technique using 45S5 Bioglasss powder. Under an optimum sintering condition (1000 1C/1h), nearly full

densification of the foam struts occurred and fine crystals of Na2Ca2Si3O9 formed, which conferred the scaffolds the highest possible

compressive and flexural strength for this foam structure. Important findings are that the mechanically strong crystalline phase

Na2Ca2Si3O9can transform into an amorphous calcium phosphate phase after immersion in simulated body fluid for 28 days, and that

the transformation kinetics can be tailored through controlling the crystallinity of the sintered 45S5 Bioglasss. Therefore, the goal of an

ideal scaffold that provides good mechanical support temporarily while maintaining bioactivity, and that can biodegrade at later stages

at a tailorable rate is achievable with the developed Bioglasss-based scaffolds.

r 2005 Elsevier Ltd. All rights reserved.

Keywords: Scaffolds; Bone tissue engineering; Mechanical properties; Bioactivity; Biodegradation; Replication technique

1. Introduction

Tissue engineering seeks to promote the regeneration

ability of host tissue through a designed scaffold that is

populated with cells and signalling molecules. The specific

criteria for ideal scaffolds used in bone tissue engineering

are summarised as follows[13]: (1) ability to deliver cells,

(2) excellent osteoconductivity, (3) good biodegradability,

(4) appropriate mechanical properties, (5) highly porous

structure: porosity490%[4]and pore sizes 4400500mm

[5], (6) irregular shape fabrication ability, and (7)commercialisation potential.

Bioactive glasses meet the first three criteria: excellent

osteoconductivity and bioactivity [610], ability to deliver

cells [11], and controllable biodegradability[1214]. These

advantages make bioactive glasses promising scaffold

materials for tissue engineering [1517]. Among a variety

of processes for fabrication of porous materials [5,1821],

the replication technique [22] (also called the polymer-

sponge method) produces porous ceramic structures that

are most similar to those of spongy bone [23,24]. This

technique also satisfies scaffolds criteria (5)(7) mentioned

above. Thus, all criteria for an ideal tissue engineering

scaffold, except that related to mechanical competence,

could be satisfied by 45S5 Bioglasss foams fabricated by

the replication method. The replication method has been

applied to produce scaffolds of hydroxyapatite (HA)

[2527]. Surprisingly, this technique, however, has never

been considered before to produce scaffolds from bioactiveglasses. Bioactive glass scaffolds have only been fabricated

by dry-powder processing with porogen additions [2830]

and by solgel and gel-casting techniques[3,31].

The major hurdle in the production of highly porous

Bioglasss-based foam-like scaffolds has been caused by the

following apparently irreconcilable issues of this glass: (a)

it has been reported that crystallisation of 45S5 Bioglasss

turns a bioactive glass into an inert material [32]; (b) full

crystallisation of the glass occurs prior to significant

densification [33]; (c) extensive densification is required to

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E-mail address: [email protected] (A.R. Boccaccini).
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strengthen the struts of a foam, which would otherwise be

made of loosely bonded particles and thus be too fragile to

handle. According to these three factors, to maintain the

bioactivity of 45S5 Bioglasss, one should sinter the foam

at a relatively low temperature at which crystallisation does

not take place or does not occur to a great extent.

However, sufficient densification by sintering will not occurat low temperatures, and therefore a very fragile scaffold

made of loosely packed 45S5 Bioglasss particles is

produced.

The above dilemma might be solved in light of the recent

work of Clupper and Hench [3437], who carried out

quantitative investigations on the effect of crystallinity on

the apatite formation on Bioglasss surfaces in vitro. Their

findings revealed that the crystal phase Na2Ca2Si3O9slightly decreased the formation kinetics of an apatite

layer on the Bioglasss sample surface but it did not totally

suppress the formation of such layer [34]. Moreover, it is

recognised that the bioreaction kinetics of a highly porous

network can be very different from that of a dense product

of the same chemical composition due to a high surface

area in the foams. Hence, it might be possible to find a new

sintering protocol leading to mechanically competent

foams through extensive densification of the struts, while

inducing the formation of a bioactive and biodegradable

crystalline phase. The objectives of this work, therefore,

were to synthesize 45S5 Bioglasss scaffolds using the

replication technique, to achieve mechanically stable 3D

scaffolds through a tailored sintering schedule, and to

assess the bioactivity and biodegradability of the scaffolds.

The final goal is to create an ideal scaffold for bone tissue

engineering.

2. Materials and experiments

2.1. Materials

The starting material was melt-derived 45S5 Bioglasss powder (particle

size 5mm). A fully reticulated polyester-based polyurethane foam with

60 ppi (pores per inch) from Recticel UK (Corby) was used as sacrificial

template for the replication method. The details of the polyurethane foam

used have been reported by other authors [38]. The foam was supplied in

large samples of 20mm in thickness and was cut to size 10mm

10mm 20mm for compression strength tests and 10mm 10mm

60 mm for bending strength tests.

2.2. Scaffold fabrication

The replication method involves preparation of green bodies of ceramic

(or glass) foams by coating a polymer (e.g. polyurethane) foam with a

ceramic (or glass) slurry. The polymer, having the desired pore structure,

simply serves as a sacrificial template for the ceramic coating. The polymer

template is immersed in the slurry, which subsequently infiltrates the

structure and ceramic (glass) particles adhere to the surfaces of the

polymer. Excess slurry is squeezed out leaving a more or less homogeneous

coating on the foam struts. After drying, the polymer is slowly burned out

in order to minimise damage to the ceramic (glass) coating. Once the

polymer has been removed, the ceramic (or glass) network is sintered to a

desired density. The process replicates the macrostructure of the starting

sacrificial polymer foam, and results in a rather distinctive and well-

defined microstructure within the struts. A flowchart of the process is

given inFig. 1.

In our experiments, the slurry for the impregnation of the polyurethane

foam was prepared using the following recipe. Polyvinyl alcohol (PVA)

was dissolved in water, the ratio being 0.01 mol/L. Then 45S5 Bioglasss

powder was added to 100 ml PVA-water solution up to concentration of

40 wt%. Each procedure was carried out under vigorous stirring using a

magnetic stirrer for 1 h.

The polyurethane foams cut to shape were immersed in the above-

prepared slurry and remained in it for 15 min. The foams were manually

retrieved from the suspension as quickly as possible, and the extra slurry

was completely squeezed out. The samples (called green bodies) were then

placed on a smooth surface and dried at ambient temperature for at least

12h. The coating thickness of a green body could be increased byrepeating the above coating procedure. In this work most green bodies

were prepared by single coating, but few were made by double coating.

The double-coated green bodies will be mentioned where they are used in

this paper.

Post-forming heat treatments for the burnout of the polymer template

and sintering for the 45S5 Bioglasss structure were programmed, as

shown inFig. 2. The burning condition of the polymer templates was the

same for all samples: 400 1C/1 h. Sintering conditions were designed to be

900 1C/5h; 950 1C/05 h; and 10001C/02 h. The heating and cooling rates

were 2 and 5 1C/min, respectively.

2.3. Characterisation

The density r foam of the scaffolds was determined from the mass and

dimensions of the sintered bodies. The porosity p was then calculated by

p 1 rfoamrsolid

1 rrelative , (1)

where rsolid 2:7 g=cm3 is the density of solid 45S5 Bioglasss [14].

The microstructure of the foams was characterised in a JEOL 5610LV

scanning electron microscope (SEM), before and after immersion in

simulated body fluid (SBF). Samples were gold- or carbon-coated and

observed at an accelerating voltage of 15 kV.

Selected foams were also characterised using X-ray diffraction (XRD)

analysis with the aim to assess the crystallinity after sintering and

formation of HA crystals on strut surfaces after different times of

immersion in SBF. The foams were first ground into a powder. Then 0.1 g

of the powder was collected for XRD analysis. A Philips PW 1700 Series

automated powder diffractometer was used, employing Cu ka radiation

(at 40 kV and 40 mA) with a secondary crystal monochromator. Data were

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Prepare slurry from the powder

Prepare a green body by dipping a

polymer foam in the slurry

Ceramic (or glass) powder

Dry, burn out sacrificial polymer

foam, and sinter the green body

Ceramic (or glass) foam

AddBinder

Fig. 1. Flowchart of the polymer-sponge method for fabrication of glass

or ceramic foams.

Q.Z. Chen et al. / Biomaterials 27 (2006) 24142425 2415


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collected over the range of 2y 51001 using a step size of 0.041 and a

counting time of 25 s per step.

2.4. Mechanical testing

The compression strength of foams was measured using a Zwick/RoellZ010 mechanical tester at a crosshead speed of 0.5 mm/min. The samples

were rectangular in shape, with dimensions: 10mm in height and

5 mm 5 mm in cross-section. During compression test, the load was

applied until densification of the porous samples started to occur.

Three-point bending strength tests were carried out using a Hounsfield

testing machine. The size of the specimens was 3 mm 4 mm 40 mm.

The load was applied over a 30 mm span and at the mid-point of the

4 mm 40 mm surface. All tests were performed using a cross-head speed

of 0.5 mm/min. The bending strength was calculated according to [39]:

sf3PfL

2bh2 , (2)

where Pf is the load at fracture, L 30 mm is the sample length over

which the load is applied, bE4 mm is the sample width, and hE3 mm is

the sample height.

2.5. Assessment of bioactivity in simulated body fluid

This part of the study was carried out using the standard in vitro

procedure described by Kokubo et al. [40]. The foams were immersed in

75 ml of acellular SBF in clean conical flasks, which had previously been

washed using HCl and deionised water. The conical flasks were placed

inside an incubator at controlled temperature of 37 1C. The pH of the

solution was maintained constant at 7.25. The size of all samples for these

tests was 10 mm 10mm 10 mm. Two samples were extracted from the

SBF solution after given times of 3, 7, 14, and 28 days. The SBF was

replaced twice a week because the cation concentration decreased during

the course of the experiments, as a result of the changes in the chemistry of

the samples. Once removed from the incubation, the samples were rinsedgently, firstly in pure ethanol and then using deionised water, and left to

dry at ambient temperature in a desiccator.

3. Results

3.1. Porous structure of foams

All foams exhibited porosity of90%, as determined by

measurement of their mass and dimensions and applying Eq.

(1). The cell size of sintered scaffolds was estimated as follows.

The cell size of the as-received polymer foam was

7401040mm. The volume shrinkage from a polymer template

to a sintered 45S5 Bioglasss-based scaffold was defined as

VBGfoam=VPU-foam, and it was determined, through measur-ing the volumes of the starting polymer and sintered 45S5

Bioglasss-based foams, to be 33% on average for the sintering

condition of 10001C/1 h. Therefore, the linear shrinkage

VBGfoam=VPUfoam1=3 would be 70%. Finally, the range

of cell sizes of the foams sintered at 10001C for 1h was

calculated to be 0.70 (7401040)mm 510720mm.The macroporous network and the strut microstructure of

typical foams are illustrated inFig. 3. Highly porous scaffolds

were produced at all sintering conditions. A comparison of

Figs. 3a, c and d shows that the cell struts are considerable

thicker when sintered at 10001C for up to 1h than at

9009501C for 25 h. It was observed at high magnification

that extensive sintering of 45S5 Bioglasss particles did not

occur at 900 1C even after 5 h sintering (Fig. 3b), but

densification, which occurs by a viscous flow sintering

mechanism in glass, increased significantly when the foams

were heated up to 950 and 1000 1C (Figs. 3d and 3f). Fine

crystalline grains of0.5mm in diameter could be detected by

SEM observation in foams sintered at 1000 1C for 1 h(Fig. 3f).

The combination of extensive densification and the presence of

a crystalline phase in the struts of scaffolds sintered at 1000 1C

for 1 h are expected to lead to improved mechanical properties

of these foams. Hence mechanical tests and assessment of

bioactivity in SBF were carried out on foams sintered at

1000 1C, as described in Sections 3.3 and 3.4.

The hollow nature of a strut and its wall microstructure

are shown in Fig. 4. Similar morphologies have been

reported for a variety of sintered ceramic foams synthesised

by the polymer-sponge method[41]. It can be seen that the

wall of the strut has been nearly fully densified after

sintering at 1000 1C for 1 h.

3.2. Crystallisation

The XRD investigation revealed that crystallisation had

occurred extensively in all samples sintered at 900, 950 and

1000 1C for 5 h (Fig. 5). However, the bonding of particles

was not obvious at the sintering condition of 9001C/5 h

(Fig. 3a). This observation confirmed the finding of

Clupper and Hench [33] that extensive crystallisation

occurs prior to significant viscous flow sintering in 45S5

Bioglasss and related bioactive glasses. In Fig. 5, both

angular location and intensity of the peaks match the

standard PDF #22.1455, which indicates that the crystal-

line phase is Na2Ca2Si3O9. The same crystalline phase has

been formed and identified in previous studies on sintered

bioactive glasses[36,37].

The present 45S5 Bioglasss-based foams are in fact made

of a glassceramic, as the crystallinity of the sintered 45S5

Bioglasss material cannot be 100%. From the components

of 45S5 Bioglasss and Na2Ca2Si3O9(Table 1), one can find

that the Na2Ca2Si3O9phase would demand too much CaO

to fully crystallise from Bioglasss. Eventually CaO is

depleted when the crystallinity reaches 80.7 mol% (i.e.

77.4 wt%), which is thus the maximum crystallinity achiev-

able by the 45S5 Bioglasss composition.

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900-1000C/0-5hr

2C/min

2C/min

5C/min

400C/1hr

R.T.

Time

Temperature

Fig. 2. Heat treatment program designed for burning-out the polyur-

ethane templates and sintering the 45S5 Bioglasss green bodies.

Q.Z. Chen et al. / Biomaterials 27 (2006) 241424252416


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3.3. Mechanical properties

Compressive and bending strength tests were carried out

on foams prepared by single or double coating and sintered

at 1000 1C for 1 h. A typical compressive stressstrain curve

is shown in Fig. 6, which is jagged and has three distinct

regimes. The foams tend to crack first in thin struts at

stress-concentrating sites, causing the apparent stress to

drop temporarily. But the foam, as a whole, still had the

ability to bear higher loads, causing the stress to rise again.

The repetition of this procedure gave a jagged stressstrain

curve.

In stage I (Fig. 6), the stressstrain curve has a positive

slope until a maximum stress is reached. This maximum

stress causes the thick struts of the foam to fracture and as

a result the stressstrain curve has a negative slope in stage

II. In Stage III, densification of the fractured foams occurs

as stress increases, which is the typical behaviour of foams

under compression[42].

The raw data of compressive strength are plotted against

the foam porosities in Fig. 7. The compressive tests were

frequently accompanied by shearing, which was mainly

caused by the end effects imposed on the specimen during

the test. It has been reported that if the faces of the foam

sample are slightly misaligned with the loading platen,

large stress concentrations can occur causing local buck-

ling, which in turn leads to shearing and thus results in an

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Fig. 4. The hollow centre of a single strut in a Bioglasss derived foam

sintered at 1000 1C for 1h.

Fig. 3. Pore structure and strut microstructure of 45S5 Bioglasss-derived foams sintered at (a)(b) 900 1C for 5 h; (c)(d) 9501C 2 h; and (e)(f) 10001C

for 1h.



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underestimation of both Youngs modulus and strength

[43]. InFig. 7, the apparent strength values of the sheared

foams (marked byD) were much lower than those of purely

compressed samples (marked by solid triangles m). It is

reasonable to consider that the strength values obtained

from pure compression tests represent the compressive

strength of the foams.

Fig. 8illustrates a typical forcedisplacement curve in a

three-point bending strength test. Like in the compressive

strength test, thin struts cracked first at stress-concentrat-

ing sites, giving a typical jagged curve. When a maximum

stress (bending strength) was reached, the sample fractured

into two pieces, causing the stress to drop to zero abruptly.

The raw data of three-point bending strength of as-

sintered and coated foams are given in Fig. 9. Bending

strengths are collectively higher than compressive strengths

at similar porosities. For instance, when porosity is 90%

the highest compressive and bending strengths are in therange of 0.30.4 and 0.40.5 MPa, respectively. This result

is in agreement with the general findings in ceramics and it

is related to the statistic nature of strength value of highly

porous brittle materials[44].

3.4. Bioactivity assessment in SBF

Assessment of bioactivity was carried out on foams

sintered at 1000 1C for 0.5 and 1 h. Similar XRD results

were obtained for both groups of foams.Fig. 10shows the

XRD spectra of the foams sintered at 1000 1C for 1h and

then immersed in SBF for 328 days, together with the

XRD patterns of 45S5 Bioglasss in as-received and as-

sintered conditions.

A significant phenomenon, in addition to the growing

peaks of HA-like phase detected in the spectra of soaked

samples, was that the crystallinity of the sintered foams

decreased with increasing immersion time in SBF. Even-

tually the sharp diffraction peaks of the Na2Ca2Si3O9phase disappeared from the XRD spectrum after soaking

in SBF for 28 days, leaving a typical broad halo (produced

by an amorphous phase) overlapped by the sharp diffrac-

tion peaks of the HA phase. This indicates that at least

under the detection limits of XRD, the sintered 45S5

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0

100

200

100200

300

400

500

600

700

800

100

200

300

400500

600

700

800

900

1000

0 20 40 60 80 100

2 ()

Intensity(a.u.)

900C/5hrs

1000C/1hr

As-Received

Apatite

Apatite

Fig. 5. XRD spectra of 45S5 Bioglasss powder unsintered and sintered at 900 1C for 5 h and 10001C for 1 h. All spectra were obtained using 0.1 g powder.

The major peaks of the phase Na2Ca2Si3O9 [35]are marked by (X).

Table 1Components of 45S5 Bioglasss and crystalline phase Na2Ca2Si3O9(mol.%)

45S5 Bioglasss Na2Ca2Si3O9

SiO2 46.134 50

Na2O 24.35 16.667

CaO 26.912 33.333

P2O5 2.6038 0



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Bioglasss-derived material was mainly composed of an

amorphous phase and crystalline apatite after soaking in

SBF for 28 days.

The microstructural evolution in both groups of foams

(sintered at 1000 1C for 0.5 and 1 h) is summarized in

Table 2. Figs. 11(ad) illustrate typical surface morphol-

ogies of samples sintered at 1000 1C for 0.5 h followed by

immersion in SBF for different time periods.

4. Discussion

In this section, the results of the investigation are

discussed in relation to mechanical properties, microstruc-

ture, and bioactivity.

4.1. Comparison of 45S5 Bioglasss-based foams with

spongy bone

The foams produced by using the polymer-sponge

method are very similar to spongy bone (also called

cancellous bone) in terms of their pore structure. It is thus

of importance to find out whether or not the mechanical

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I

100

Compressive

Stress(MPa)

0 20 40 60 800.0

0.1

0.2

0.3

0.4

0.5

Compressive strain (%)

IIIII

Fig. 6. A typical compressive stress-strain curve of the 45S5 Bioglasss-based foams sintered at 1000 1C for 1 h. The porosity of the foam was 91.0%.

0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

0.9

0.84 0.86 0.88 0.9 0.92 0.94 0.96

Porosity

Co

mpressivestrength(MPa)

Theoretical strength (ti/t=0)

Theoretical strength (ti/ t=0.5)

Experimental strength with shearing involved

Experimental strength without shearing

Reported strength of HA-based foams in literature [25-27]

Fig. 7. Theoretical and experimental compressive strength values of

Bioglasss-based scaffolds in the present work, and those of hydroxyapa-

tite-based foams reported in literature [2527]. The foams with porosity

lower than 89% were prepared by double coating. Theoretical values were

obtained from Eq. (3).

1

2

3

4

00.20

Force(N)

Displacement (mm)

5

Fig. 8. A typical forcedisplacement curve of 45S5 Bioglasss-based foam

in bending test.

0

0.2

0.4

0.6

0.8

1

1.2

1.4

0.76 0.78 0.8 0.82 0.84 0.86 0.88 0.9 0.92 0.94

Porosity

Bendingstrength(MPa)

Fig. 9. Bending strength values of Bioglasss-based scaffolds. The foams

with porosity lower than 89% were prepared by double coating.



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strength of the foams is comparable to that of cancellous

bone.

There have been many reports on the mechanical

properties of cancellous bone, which have been reviewed

in Ref. [23]. It is generally accepted that the mechanical

properties of struts in cancellous bone are close to those of

cortical bone. Typical values are: 12 GPa for Youngs

modulus, 136 MPa for compressive strength, and 105 MPa

for tensile strength [24]. The compressive strength of

spongy bone (not the strut) is in the range of 0.24 MPa,

when the relative density is 0.1[24]. Hence, the measured

compressive strength (0.30.4 MPa) of the present foams

falls in this range, but lies closer to the lower bound. Our

experience indicates that the strength of 0.30.4 MPa is

sufficient for the foam to be handled with, such as

manipulating during SBF tests and cutting of the samples

for mechanical tests.

In addition, it has been reported that the compressive

strength of a HA scaffold significantly increases (e.g. from

10 to 30 MPa[45]) due to tissue ingrowth in vivo. It has

also been speculated that it might not be necessary to

fabricate a scaffold with a mechanical strength equal to

bone because cultured cells on the scaffold and new tissue

formation in vitro will create a biocomposite and will

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0 10 20 30 40 50 60 70 80 90 100

As-received

As-sintered

3 days

7 days

14 days

28 days

2 ()

200

0

1000

800

600

400

0

600

400

200

0

600

400

200

0

400

200

0

400

200

0

Intensity(a.u.)

Fig. 10. XRD spectra of 45S5 Bioglasss-based foams sintered at 1000 1C for 1 h, and immersed in SBF for 3, 7, 14, and 28 days. All spectra were obtained

using 0.1 g powder. The major peaks of Na2Ca2Si3O9 phase and hydroxyapatite are marked by (X) and (K), respectively.

Table 2

Summary of characteristics of 45S5 Bioglasss-derived foams after immersion in SBF

Immersion time in SBF 1000 1C/30 min 1000 1C/1h

3 days Sparsely distributed apatite precipitates Very few apatite precipitates

1 week Strut surface was unevenly covered by aggregated

apatite spheres

Sparsely distributed apatite precipitates

2 weeks Apatite spheres were fused together. Strut surface was fully covered by a large amount of

apatite spheres, size being 1mm

4 weeks The whole foam is made of amorphous calcium

phosphate and crystalline hydroxyapatite

Apatite spheres grew, size being 2.5mm

The apatite could be a mixture of amorphous and crystalline calcium phosphates.



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increase the time-dependent strength of the scaffold

significantly [16]. An ideal scaffold, however, should have

at least a proper strength and fracture toughness to allow it

to be manipulated adequately for tissue engineering

applications. The present 45S5 Bioglasss-based scaffolds

possess such an appropriate mechanical competence.

There is no reliable fracture toughness data available forcancellous bone. But it is predicted that the current foams

will be more brittle than spongy bone. A further study

involving the incorporation of poly(D,L-lactic acid) into the

45S5 Bioglasss-based foams is on-going, aiming at

improving the fracture toughness of these scaffolds.

4.2. Comparison of the present foams with previous

investigations

There are few reports available on porous 45S5

Bioglasss and related bioactive glassceramics with low

porosity (2142%) [2830], but no work has been

published on highly porous (p490%) 45S5 Bioglasss-

based foams, to the best knowledge of the authors.

Therefore, the comparison carried out here is between the

present scaffolds and highly porous (p470%) foams made

from other bioactive ceramics and glasses, including HA,

b-tricalcium phosphate (b-TCP), and 70S30C solgel

derived glass foams.

Table 3 summarises characteristics of highly porous

bioactive ceramic and glass foams developed for bone

engineering, including method of fabrication, pore structure,

and compressive strength data. In general, the compressive

strength varies significantly with foam porosity. For example,

the compressive strength of HA foams, which were

synthesised by the polymer-sponge method, decreased from

0.29 to 0.03MPa when the porosity increased from 69 to 86%

[27]. Some compressive strength data of porous HA-based

foams reported in literature [25,26] have been collected and

are shown in Fig. 7 (marked by ). It is obvious that the

present 45S5 Bioglasss-based foams are in general stronger

than the HA-based foams of similar porosities.It is unwise to directly compare foams exhibiting partially

open pore structure with completely open pore scaffolds

fabricated by the polymer-sponge method. The high mechan-

ical strength of the former[46,47]is obviously achieved at the

cost of a less interconnected pore structure. The windows on

the wall of pores in these foams are mainly in the range of

30120mm, which is considerably smaller than the required

size (400mm) for osteoblast penetration[5].

It is apparent that gel-casting combined with the

polymer-sponge technique produces stronger HA foams

[48,49] than the simple polymer-sponge method [26].

However, a comparison of the present 45S5 Bioglasss-

based foams (0.42 MPa at porosity 89%) with HA foams

produced by Ramay and Zhang[48](0.55 MPa at porosity

77%) indicates that the polymer-sponge method developed

here can produce as strong foams as the gel-casting/

polymer-sponge combined technique, and that the well-

sintered and crystallised 45S5 Bioglasss-based scaffolds

can be as strong as HA foams.

4.3. Comparison of experimental and theoretical strength

data

The modelling of the mechanical behaviour of highly porous

materials has been presented by Gibson and Ashby [24].

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Fig. 11. Hydroxyapatite formed on the surfaces of foam struts after immersion in simulated body fluid (SBF) for (a) 3 days, (b) 7 days, (c) 14 days, and (d)

28 days. The foams were sintered at 1000 1C for 30 min.



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The theoretical compressive collapse stress stheo can be

expressed as a function of the relative density rfoam=rsolidof a cellular structure and the size of the central hollow

struts by Eq. (3):

stheo

sfs 0:2

rfoamrsolid

3=21 ti=t

2ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi1 ti=t

2q , (3)

whereti=tis the ratio of the void and strut sizes on a cross-section of a strut (see Fig. 4) and sfs is the modulus of

rupture of the strut. Theoretical calculations show that themodulus of rupture of a brittle material is typically about

1.1 times larger than the tensile strength [24]. In our

calculation, the tensile strengthsts 42 MPa of bulk 45S5

bioglasss (annealed)[14] was used for the strength of the

partially crystallised material. The ratio t i=t was estimatedto be 0.5 according toFig. 4.

Using Eq. (3), the compressive strength of the present

foams (with ti=t 0:5) and the lower bound of theoreticalstrength (when ti=t 0) were calculated. The results areillustrated inFig. 7. The experimental strengths determined

by compressive strength tests are generally above the lower

bound, and most of them are in good agreement with the

theoretical strengths when ti=t 0:5. This indicates thatthe cell walls have been sintered to be fully dense at 1000 1C

for 1 h.

Two points should be mentioned. (i) The tensile strength

of sintered HA is 40 MPa, which is very close to that of

dense 45S5 Bioglasss (42 MPa). Hence theoretically, the

mechanical strength of a 45S5 Bioglasss-based scaffold

should be similar to, if not higher than, that of a HA

scaffold with a similar porous structure. (ii) As ti=tincreases, the foam becomes stronger. In other words, the

hollow tubular structure is beneficial to the mechanical

performance of the foam, which is a direct result of the

derivation of Eq. 3 [24].

4.4. Possible mechanisms for the transition from

Na2Ca2Si3O9 to an amorphous phase

Since the sintered 45S5 Bioglasss material is in fact a

glassceramic, one might argue that the bioactivity of the

sintered material could be attributed to the residual glass

phase. We suggest that the bioactivity remains also with the

crystalline phase Na2Ca2Si3O9, based on two reasons: (1)

the bioactivity of pure Na2Ca2Si3O9 phase has been

reported [35], and (2) the transition from Na2Ca2Si3O9 to

an amorphous phase provides an explanation for thefinding that the presence of Na2Ca2Si3O9 decreased the

kinetics of apatite formation but did not inhibit the growth

of an apatite layer on the form surfaces, which has been

reported in the literature [34].

The mechanisms behind the transformation of Na2Ca2-Si3O9 to an amorphous phase might be based in the well-

known bone-bonding mechanisms of bioactive glasses,

which were originally proposed by Hench and colleagues

[14]. In the sequence of interfacial reactions on the surface

of Bioglasss in contact with body fluids, the bioactive glass

first dissolves to form a silica-gel layer; then an amorphous

calcium phosphate is formed from the hydrated silica-gel;

and finally apatite crystallites nucleate and grow from the

amorphous calcium phosphate. We suggest that the general

idea of the reaction sequence should be applicable to

Na2Ca2Si3O9 crystallites as well, which however dissolves

at a slower rate than the glass phase. Hence, the

amorphous phase detected by XRD after immersion in

SBF for 28 days (Fig. 10) could be the amorphous calcium

phosphate, according to Hench et al.s theory [14]. This

suggestion has been proved by energy dispersive X-ray

(EDX) analysis, as shown elsewhere [50].

Although the kinetics of the transformation has yet to be

fully understood, it is believed that the high surface area

(including hollow centre of the struts) in the porous

ARTICLE IN PRESS

Table 3

Overview of structural characteristics and mechanical properties of highly porous bioactive ceramic or glass foams for bone tissue engineering

Technique Material Porosity (%) Pore size (mm) Closed (C) or

open (O)

Compressive

strength (MPa)

Ref.

Polymer-sponge 45S5 Bioglasss 8992 510720 O 0.270.42 Present work

Glass-reinforced HA 8597.5 420560 O 0.010.175 [25]

HA 86 420560 O 0.21 [26]

6986 4901130 O 0.030.29 [27]

Gel-casting/

foamed by

vigorous stirring

HA 76.780.2 201000 Partly O/C 4.47.4 [46]

HA Cell: 100500 Partly O/C 1.65.8 [47]

Window: 30120

Polymer-sponge HA 7077 200400 O 0.555 [48]

b-TCP+HA 73 200400 O 9.8 [49]

Solgel/foamed by

vigorous stirring

Bioactive glasses (e.g.

70S30C)

7095 Cell: up to 600 Partly O/C 0.52.5 [31]

Windows: 80120



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network is of relevance in maintaining bioactivity and

biodegradability of the sintered 45S5 Bioglasss-based

foams. The high surface energy should make possible that

the transformation of Na2Ca2Si3O9 to the amorphous

phase of calcium phosphate occurs at a reasonably fast rate

at the body temperature. This assumption is supported by

the fact that the bioactive reactions only occur at thesurface of a bulk solid glass. It is based on this fact that

bioactivity is defined to be the interfacial ability to bond to

bone [14]. Nevertheless, the transformation mechanism

from a crystal to an amorphous phase, as found in this

material system in contact with SBF, remains a subject for

future dedicated research.

4.5. Significance of the transformation of Na2Ca2Si3O9 to

the amorphous phase

An ideal scaffold for bone engineering serves as a

temporary frame to foster new bone growth. It is expected

to provide a temporary mechanical support and later to

degrade at a rate matching the regeneration rate of new

bone tissue. Unfortunately, this is not the manner

conventional bioceramics behave in biological conditions.

Crystalline HA, for instance, can provide reasonably

strong support; but it degrades very slowly in contact

with body fluids, degradation time being of the order of

years. Amorphous HA degrades much faster than crystal-

line HA; but it is too fragile to build highly porous

scaffolds. Bioactive glasses encounter a similar hurdle:

they have excellent bioactivity and tailorable biodegrad-

ability; at the same time they possess poor mechanical

reliability.The above problem could be solved by designing

scaffolds following the discovery of this work, which has

shown that the mechanically strong crystalline phase

Na2Ca2Si3O9 (which is formed in 45S5 Bioglasss upon

sintering) can transform into an amorphous calcium

phosphate (the good resorbability of which has been well

documented [51,52]) in a simulated body fluid environ-

ment. Based on this finding, it is possible to sinter 45S5

Bioglasss green foams at optimised conditions such that

both significant densification and Na2Ca2Si3O9 crystal-

lisation take place. The extensive densification and fine

crystalline grains of Na2

Ca2

Si3

O9

confer the scaffold a

temporary good mechanical performance. The transforma-

tion of Na2Ca2Si3O9to the amorphous calcium phosphate,

which is expected to occur upon exposure to a body fluid

environment, ensures the bioactivity and degradability of

the scaffold.

The transformation of a crystalline phase to a degrad-

able amorphous phase is not an exclusive phenomenon of

45S5 Bioglasss material. HA and related calcium phos-

phates also show a similar transition in an in vivo

environment [53]. The difference with Bioglasss is that

the transition in HA is too slow to match clinical

expectation. It has been shown that only a thin layer of

amorphous phase on the surface of crystalline HA particles

(0.5 mm) is formed after implantation for 3 months, and

that HA particles do not degrade considerably even after

implantation for 6 months [53]. Tissue engineering

applications demand that the degradation kinetics of a

scaffold should match the regeneration kinetics of new

bone in vitro and/or in vivo. In general, the degradation

time should be less than 6 months, depending on theanatomic site for regeneration, the mechanical loads

present at the site, and the desired rate of osseointegra-

tion[1].

Hence, the significance of the Na2Ca2Si3O9 to amor-

phous phase transition in our 45S5 Bioglasss-based foams

lies in its kinetics which seems to be sufficiently fast for

application of the material in bone engineering. More

importantly, the kinetics of the transformation and of the

scaffold degradation can be controlled by factors such as

initial crystallinity, porosity in struts, and grain size of

Na2Ca2Si3O9, all of which can be tailored by the sintering

conditions. Therefore, the goal of an ideal scaffold that

provides good mechanical support temporarily while

maintaining bioactivity, and that can biodegrade later at

a tailorable rate can be achieved with the developed 45S5

Bioglasss-derived scaffolds.

5. Conclusions

This work has successfully synthesized highly porous

(porosity: 90%, cell diameter: 510720mm), mechanically

competent, bioactive and biodegradable 45S5 Bioglasss-

derived glassceramic scaffolds for bone engineering, using

the replication technique. When sintered under an optimal

condition (10001C/1 h), the nearly full densification and

the fine crystals of Na2Ca2Si3O9 confer the scaffolds

competent mechanical strength. A significant finding is

that the mechanically strong crystalline phase can trans-

form into a bioactive and biodegradable amorphous

calcium phosphate upon immersion in SBF. Therefore,

the goal of an ideal scaffold that provides sufficient

mechanical support temporarily, and that can biodegrade

later at a tailorable rate is achievable with the present

Bioglasss-derived glassceramic scaffolds.

Acknowledgements

Helpful discussions on the sintering experiments with

Mr. Jonny Blaker (Imperial College London) are gratefully

acknowledged. Recticel UK at Corby is gratefully

acknowledged for providing the polyurethane foam used

in this research. Helpful discussions with Prof. Larry

Hench are greatly appreciated.

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