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Journal of the European Ceramic Society 34 (2014) 107–118

Robocasting of 45S5 bioactive glass scaffolds for bone tissue engineering

Siamak Eqtesadi a, Azadeh Motealleh a, Pedro Miranda a,∗, Antonia Pajares a, Alexandra Lemos b,José M.F. Ferreira b

a Departamento de Ingeniería Mecánica, Energética y de los Materiales, Universidad de Extremadura, Escuela de Ingenierías Industriales, Avda. de Elvas s/n,06006 Badajoz, Spain

b Department of Materials Engineering and Ceramics, University of Aveiro, CICECO, 3810-193 Aveiro, Portugal

Received 14 May 2013; received in revised form 2 August 2013; accepted 4 August 2013Available online 4 September 2013

bstract

hree dimensional scaffolds with controlled pore architecture were prepared from 45S5 Bioglass® powders by robocasting (direct-writing) usingarboxymethyl cellulose (CMC) as the single processing additive. A comprehensive sintering study of the resulting structures was performed withinhe temperature range 500–1050 ◦C. Robocast scaffolds with interconnected porosities ranging from 60 to 80% were obtained for a fixed scaffoldesign. All scaffolds exhibited compressive strengths comparable to that of cancellous bone (2–13 MPa), including those sintered at temperatureselow the crystallization temperature of 45S5 bioactive glass. These strength values are substantially higher than any previously reported data for

®

5S5 Bioglass scaffolds and imply that robocasting is the first technique which can be considered suitable for producing vitreous 45S5 scaffoldsith a sufficient mechanical integrity for any practical application. Moreover, this process will enable the development of 45S5 Bioglass® scaffoldsith customized external geometry, and optimized pore architecture.

2013 Elsevier Ltd. All rights reserved.

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eywords: 45S5 bioactive glass; Robocasting; Direct-writing; Scaffold; Sinteri

. Introduction

The 45S5 Bioglass® developed by Larry Hench and col-eagues at the University of Florida,1 with a composition of 45%iO2, 24.5% CaO, 24.5% Na2O and 6% P2O5 (in weight %,t.%), is a commercially available inorganic biomaterial with

long history of success as bone replacement.2,3 This glasss claimed to be highly bioactive and both osteoinductive andsteoconductive, enabling its application in bone regenerationnd tissue engineering.4–7 However, despite the fact that afterore than 40 years of research no other calcium phosphate

r bioactive glass composition has surpassed 45S5 Bioglass®

iological performance, its commercial success is somewhatimited and calcium phosphates remain market leaders as bone

3

ubstitute materials.The main reason for this limited success lays in the difficul-

ies found when processing 45S5 bioactive glass powders into

∗ Corresponding author. Tel.: +34 924 289 300x86735; fax: +34 924 289 601.E-mail addresses: pmiranda@unex.es (P. Miranda), jmf@ua.pt

J.M.F. Ferreira).

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955-2219/$ – see front matter © 2013 Elsevier Ltd. All rights reserved.ttp://dx.doi.org/10.1016/j.jeurceramsoc.2013.08.003

trength

orous structures suitable for use as scaffolds for bone regen-ration: that is, with a porosity appropriate for inducing bonengrowth and the sufficient mechanical integrity to be handleduring implantation and, if possible, to replace bone struc-ural function during the whole tissue regeneration process. Theroblem lies in the small sintering window—temperature differ-nce between glass transition temperature (∼550 ◦C) and thatf the onset of crystallization (∼610 ◦C)—associated to 45S5lass particles, which prevents any significant degree of den-ification to occur before crystallization.8 However, extensiveensification is required to obtain a mechanically compe-ent 45S5 Bioglass®-based scaffold in a foam-like structure,hich would otherwise be made of loosely bonded parti-

les and too fragile to handle. Accordingly, fabrication oforous scaffolds from 45S5 bioactive glass powder throughraditional processing methods, such as foam replication, isade by using high sintering temperatures (∼1000 ◦C), where

he material is almost fully crystalline.9,10 Unfortunately, the

ioactivity of such crystallized structures is retarded—thenset of apatite formation in simulated body fluid (SBF)ncreases from 10 h to 22–25 h—over that of vitreous 45S5

aterial.11

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Solid freeform fabrication (SFF) methods12 provide a meanso improve the mechanical performance of scaffolds13,14 by pro-iding a much greater level of control over pore architecturend more regular strut morphologies than in conventional tech-iques. In fact the regularly interconnected pores with largenterconnection sizes of scaffolds fabricated by these additive

anufacturing techniques allows one to significantly reducehe overall porosity of the structure required to allow coloniza-ion by bone tissue.15–18 However, despite the aforementioneddvantages there are still few studies reporting the successfulpplication of SFF techniques to fabrication of 45S5 Bioglass®-erived scaffolds.19,20 We recently added to this exiguousiterature by reporting the first successful attempt to fabricate5S5 Bioglass®-derived scaffolds via a colloidal processingoute such as robocasting.21

Robocasting is an extrusion-based, direct writing technique,hat utilizes high solids loaded paste-like suspensions (inks) touild 3D structures through a layer-wise deposition of extrudedylinders.22–25 Critical rheological parameters for the printingnk include its apparent viscosity, yield stress and dynamic elas-ic modulus. An ink with appropriate rheological properties muste able to flow through the nozzle (i.e. low viscosity under hightress) and possess excellent shape retention upon depositioni.e. high yield stress and elastic modulus at rest).25,26 To achievehis goal, the starting powders are usually dispersed in water withhe addition of an anionic dispersant to obtain a concentrateduspension. The obtained flowable slurry is then coagulated bydding suitable concentrations of salts or cationic species, or byhanging the pH of the suspension to turn it into an extrudableaste.

In the aforementioned preceding paper we reported how theifficulties concomitant to the wet processing of 45S5 Bioglass®

owders, which are associated to its relatively high sodiumontent (24.5 wt.%) and the consequent pH increase resultingrom leaching of Na+ ions into the dispersing water, can bevercome by using carboxymethyl cellulose as a single but mul-ifunctional (dispersant, binder) additive.21 The current workxtends the information on the robocasting ink optimizationnd rheological properties and reports the results obtained on

comprehensive sintering study, including mechanical charac-erization, performed on 45S5 Bioglass® robocast scaffolds with

fixed three-dimensional design.

. Experimental procedure

.1. Materials and ink preparation

The bioactive glass with the 45S5 composition, containing5% SiO2, 24.5% CaO, 24.5% Na2O and 6% P2O5 (in weight, wt.%), was supplied by Mo-SCI Corporation (USA). Car-

oxymethyl cellulose (CMC, Mw = 250,000, Lamberti Iberia.a.u., Castellón, Spain) with a viscosity of a 2 wt.% aqueous

olution in the range of η = 1–3 Pa s, as indicated by the supplier,as used as the sole processing additive.The as-received powder was milled for 4 h in an attrition mill

Model 01-HD, Union Process, Akron, OH), using high-purity

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Ceramic Society 34 (2014) 107–118

irconia container and balls as milling media and ethanol as dis-ersing medium, to provide particles in the 1–10 �m size range,hich was found optimal for the preparation of the concentrated

uspensions required for robocasting—larger sizes seriouslyamper colloidal stabilization, while smaller particles tend toeduce the maximum achievable solid loading. The slurries wereried at 60 ◦C and the powder was sieved through 106 �m and5 �m stainless steel sieves to eliminate large particle agglom-rates. The size distribution of the resulting particulates waseasured using a laser-scattering particle size analyser (Mas-

ersizer 2000MU, Malvern, UK).Based on the ink optimization process described in a previ-

us work21 and some additional details provided in Appendix,he optimal recipe for the 45S5 Bioglass® robocasting inkses carboxymethyl cellulose (CMC) as a single multifunc-ional additive, acting both as dispersant and binder to providehe required rheological behaviour. According to these results,

wt.% (relative to solid content) of CMC was first dissolved ineionized water and then 45 vol.% of 45S5 glass powder wasdded to the suspension in small batches between successiveixing procedures performed in a centrifugal planetary mixer

ARE-250, THINKY, JAPAN), for a total mixing time of around0 min.

.2. Robocasting of bioglass scaffolds

A robocasting device (A3200, 3D inks, Stillwater, OK, USA)as employed to fabricate three-dimensional structures of 45S5ioglass® with the optimized inks, containing 1 wt.% of CMCs single processing additive and a solids loading of 45 vol.%.he ink was housed in a syringe and extruded through a conicalozzle (inner diameter, d = 410 �m) by the computer-controlledobotic system. The position of the nozzle moved followinghe CAD model designed previously in the control softwareRobocad 3.0, 3D inks, Stillwater, OK, USA). The scaffoldesign consisted of tetragonal mesh with a centre-to-centre spac-ng between adjacent rods within a layer of s = 2d = 820 �mnd a layer height h = 287 �m; with external dimensions of2 mm × 12 mm × 12 mm. The ink flowed through the nozzlet the volumetric flow rate required to maintain a constant depo-ition speed of 10 mm s−1. To prevent a non-uniform dryingf the structure during assembly, the deposition process wasarried out within a paraffin oil reservoir at room temperature.fter deposition, the samples were removed from the bath andried in ambient conditions for at least one day before thermalrocessing.

.3. Thermal processing of assembled structures

Thermal debinding kinetics of 3D constructs was evalu-ted using thermo gravimetric analysis, TGA, and differentialhermal analysis, DTA, (SETSYS Evolution-16 SETARAM).amples of as-received 45S5 bioactive glass powder and of dried

◦

obocasting inks (for 1 day at 100 C), were heated in air at◦C min−1 up to 1200 ◦C. Based on these kinetic data, the fol-

owing thermal treatment was selected for the removal of theolymeric additive and sintering of the as-printed structures.

pean Ceramic Society 34 (2014) 107–118 109

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Thermal burning out of CMC was performed in air at 400 ◦Cor 1 h, with a heating rate of 1 ◦C min−1. Following the binderemoval, the constructs were sintered in air for 1 h at differentemperatures up to 1200 ◦C, with a heating rate of 2 ◦C min−1.he sintered samples were cooled in furnace at a rate of approx-

mately 10 ◦C min−1.

.4. Microstructural and mechanical characterization of5S5 bioactive glass scaffolds

The density of the scaffolds, ρs, was determined from theass and external dimensions of the structures. The porosity, p,as then calculated as:

= 1 − ρs

ρth

= 1 − ρrel (1)

here ρth = 2.7 g cm−3 is the theoretical density of bulk 45S5ioglass®.3 The linear shrinkage, �L, on the samples was calcu-

ated at specific temperatures from the variation of their externalimensions as:

L(%) = Li − Lf

Li

× 100 (2)

here Li is the initial dimension of the specimen at room tem-erature, and Lf the dimension after sintering at the selectedemperature. Both shrinkage and porosity were evaluated as theverage of measurement performed on four different specimens.

The microstructure of the sintered scaffold was examinedsing a scanning electron microscope, SEM (SEM, S-3600N,itachi, Japan). Internal dimensions (rod diameter, spacing, etc.)f scaffolds fabricated at different temperatures were determineds the average of a minimum of 100 direct measurements per-ormed on SEM images. These dimensions were then used tostimate the macroscopic, pre-designed porosity from geometri-al considerations assuming an ideal network of interpenetratingylinders.

X-ray diffractometry (XRD, PW-1800, Phillips Research,he Netherlands) using CuK� radiation (step width 0.03◦ 2θ,ngular interval 10–65◦ 2θ, and count time 4 s per step) was per-ormed on 45S5 Bioglass® grounded samples after sintering atelected temperatures for identification of the phases present.

The mechanical response of the scaffolds fabricated at eachintering temperature was evaluated by uniaxial compressionn a universal testing machine (AG-IS 10KN, Shimadzu Corp.,yoto, Japan). The compression force was applied in a directionerpendicular to the printing plane (i.e. orthogonal to the scaf-old rods) by using a constant crosshead speed of 0.6 mm min−1.

ine parallelepipedic samples (3 mm × 3 mm × 6 mm) were

ested for each condition. Prior to mechanical testing, the contacturfaces of each sample were ground to produce smooth parallelurfaces.

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ig. 1. Particle size distribution of the as-milled 45S5 Bioglass® powder.

. Results and discussion

.1. Scaffold preparation

The particle size distribution of the as-milled 45S5 Bioglass®

articles used in preparing the inks for robocasting is shown inig. 1. The particles had a wide distribution of sizes, with aedian diameter of 4.3 �m. Fig. 2 shows SEM micrographs of

n as-cut scaffold created using the optimized ink recipe andeometric design described in Section 2, followed by sinteringt 1000 ◦C for 1 h. Note the regularity of the structure and thetraight geometry on the rods, attesting the quality of the inksed for its fabrication.

.2. Sintering and microstructural characterization

Fig. 3 presents TGA and DTA results of 45S5 Bioglass®

s-received powder and of the as-dried ink. The as-receivedowder underwent a slight weight loss up to ∼900 ◦C, with theore marked changes in the range of 200–400 ◦C (Fig. 3a). Thiseight loss is associated to the evaporation of free water and the

elease of surface OH groups, a phenomenon reflected as anndothermic effect, especially observed between room temper-ture and 200 ◦C in the DTA plot (Fig. 3b). After recoveringhe base line, a second endothermic dip centred at ∼560 ◦C isbserved, which is attributed to the first transition temperature ofhe glass (Tg1). This is followed by an exothermic effect startinground ∼ 600 ◦C and with a maximum at about 685 ◦C (Tp1),hich is attributed to crystallization, in agreement with previ-usly reported values.8,27,28 A second exothermic peak centredt ∼850 ◦C (Tp2), attributed to the formation of a secondary crys-alline phase, is followed by a new endothermic valley associatedo a second glass transition (Tg2 ∼ 910 ◦C) of the remaining glass

atrix.21 Melting clearly occurred above 1100 ◦C, as deducedrom the associated endothermic valley.

The weight changes (Fig. 3a) and the corresponding thermalffects (Fig. 3b) were greatly enhanced in the dried ink,robably as a result of a more extended surface hydration

110 S. Eqtesadi et al. / Journal of the European Ceramic Society 34 (2014) 107–118

Fig. 2. SEM micrographs of a 45S5 Bioglass® tetragonal mesh scaffold fabri-cated by robocasting and sintered at 1000 ◦C for 1 h: (a) corner view of an as-cutss

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Fig. 3. Results of the simultaneous (a) thermal gravimetric analysis (TGA,weight loss) and (b) differential thermal analysis (DTA, heat flow) of 45S5Bc

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ample for use in uniaxial compressive tests and (b and c) side views of the sameample.

f the powder undergone while in contact with dispersion

edia (ball milling and ink preparation), and of the burnout ofMC additive. Indeed, the weight loss centred in the range of00–800 ◦C, but extending beyond these limits, is attributed to

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ioglass® as-received powder (light colour) and as-dried robocasting inks (darkolour).

he burnout of carbonaceous remnants from CMC degradation.his degradation occurred in two successive steps, originating

wo exothermic peaks at ∼225 ◦C and ∼350 ◦C (Fig. 3b), asonfirmed by a TG/DTA test performed on pure CMC (resultsot shown). This suggested the use of a debinding treatment at00 ◦C for 1 h prior to sintering.

Interestingly, the thermal events Tg1 and Tp1 occurred in thenk, respectively, at about 60 ◦C and 30 ◦C below the correspond-ng temperatures in the as-received 45S5 Bioglass® powder,mplying a concomitant shift of the intermediate exothermicffect associated to the nucleation step. This shift is attributedo a change in the composition at/near the particles’ surface as

result of leaching of modifier cations (Na+, Ca2+) during wet-illing, and also to the reduction in the particle size.29 Besides,

p2 and Tg2 cannot be distinguished for the dried ink, possi-ly because they are also shifted to lower values and becameverlapped with the peak for first crystallization. The apparentncrease of melting temperature of the ink is also consistent withhe leaching of modifier cations into the milling media.

These results are in good agreement with the XRD results

hown in Fig. 4. The 45S5 Bioglass remained amorphous upo above 550 ◦C but there is evidence of crystallization alreadyt 600 ◦C. The angular location and intensity of the new peaks

S. Eqtesadi et al. / Journal of the European Ceramic Society 34 (2014) 107–118 111

Fig. 4. X-Ray diffractograms of 45S5 Bioglass® as-received and as-milledpowders, and ground scaffolds sintered at the indicated temperatures. Sodiumcarbonate (Na2CO3, JCPDF database 19-1130) peaks are clearly visible on theas-milled powders, and barely on the sample sintered at 550 ◦C. Peaks at 600 ◦C,correspond to Na2CaSi2O6 (JCPDF database 77-2189) crystallized from theglass. At 1000 ◦C, Na2CaSi2O6 remains the main crystalline phase with someadditional peaks, noted by � and ©, corresponding to Na2Ca4(PO4)2Si2O4

(7

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dv50% for all sintering temperatures. Consequently, all the densi-fication produced during sintering, from an as-dried green bodyporosity of 80 ± 2% to a minimum of 60 ± 2% after sintering

JCPDF database 32-1053) and tetragonal zirconia (ZrO2, JCPDF database2-7115), respectively.

ppearing upon crystallization of 45S5 Bioglass® closely matchhe standard JCPDF database 77-2189, confirming that the majorrystalline phase is Na2CaSi2O6. Minor peaks of a secondaryhase, attributed to Na2Ca4(PO4)2Si2O4 (JCPDF database 32-053), were also identified on samples sintered at 1000 ◦C, all inood accordance with previous reports.8,30–32 There are, how-ver, some peaks already observable at 550 ◦C and below, whichorrespond to sodium carbonate (Na2CO3, JCPDF database 19-130) formed during the milling process in ethanol, with somemall contribution of tetragonal zirconia from the milling mediand container. The sodium carbonate, deposited onto the sur-ace of the particles during milling, partially dissolves in wateruring ink preparation, which produces a substantial decreasen the intensities of the peaks corresponding to this extraneoushase. Actually, the peaks more clearly observable in the driednk or in samples sintered at 550 ◦C, located at 30.2◦ and 35.3◦θ, can also be appreciated even at 1000 ◦C, and are thereforettributed to the overlapping tetragonal zirconia phase ratherhan the sodium carbonate. Nonetheless, the presence of evenmall amounts of sodium carbonate in the XRD results supporthe explanation for the systematic down shift of thermal eventsexcept melting) in Fig. 3 based on a changed composition, richn Na+ cations, at/near the particles’ surface.

Fig. 5 shows the evolution of scaffolds’ macroscopic linearhrinkage and internal dimensions with the sintering tempera-ure, with standard deviations as error bars. The data at roomemperature correspond to as-dried green structures. The struc-ure shrinkage during sintering was essentially isotropic, and the

caffolds kept the designed parallelepipedical shape upon sinter-ng for 1 h up to 1050 ◦C. Above 1050 ◦C or for longer sinteringimes at this temperature, indications of melting/softening of

FD

s function of sintering temperature. Data represent mean values with standardeviations as error bars.

he material became evident to the naked eye: the scaffold spec-men deformed under its own weight and the shape was distorted,ith some material accumulation at the bottom. When treated

t 1200 ◦C for 1 h the samples completely melted and collapsedith almost no remaining macroporosity. Due to the severe dis-

ortions they presented, samples sintered above 1050 ◦C wereot characterized. In good agreement with previous reports, thecaffold contraction progressed steadily with increasing sin-ering temperature, although at a somewhat faster rate aroundhe two glass transition temperatures (500–550 ◦C, Tg1, and00–850 ◦C, Tg2)8,28), with the same trend being observed forll the internal dimensions of the scaffolds.

As shown in Fig. 6, the reduction in the different dimensionsid not translate into any appreciable change in the pre-designedolumetric macroporosity, which remained constant at around

ig. 6. Evolution of scaffold porosities as function of sintering temperature.ata represent mean values with standard deviations as error bars.

112 S. Eqtesadi et al. / Journal of the European Ceramic Society 34 (2014) 107–118

erived

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Fig. 7. SEM micrographs of the rod surfaces of 45S5 Bioglass®-d

t 1050 ◦C, was associated to a reduction of the microporosityithin the glass-ceramic rods. Meaning that the rods reduced

heir initial ∼60% porosity (which roughly agrees with the 45%olid loading of the inks) to around a final 20% porosity at theighest sintering temperature. Failure to achieve a high densifi-ation of the rods even at temperatures so close to the meltingoint is not an unexpected result, given the poor sintering abil-ty of 45S5 Bioglass® even under pressure-assisted sintering.31

gain, the evolution of densification with sintering tempera-ure found in this work (Fig. 6) is in agreement with previousesults:8,33 there is a steady densification of the scaffolds withncreasing sintering temperature, but densification is signifi-antly enhanced around the two glass transition temperatures,nd sluggish in-between.

The evolution of in-rod microporosity described in Fig. 6 cane correlated with the microstructural analysis of the scaffolds.ig. 7 shows SEM micrographs of rod surfaces of scaffold sin-

ered at the indicated temperatures. No sintering occurred upo the glass transition temperature at 550 ◦C, when the glassoftened enough to allow some viscous flow, and even at that

emperature the microstructure does not differ all that muchrom that of green samples (compare top images in Fig. 7),xcept for the evidences of inter-particle neck formation (see

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scaffold before and after sintering at the indicated temperatures.

nset in top-right image of Fig. 7). At around 600 ◦C, just afterhe glass transition, the glass started to crystallize (Fig. 4),hich severely slows down the densification process by reduc-

ng the species diffusivity and increasing the system viscosity.t this point, it is worth mentioning that the maximum per-

entage of crystallinity achievable in 45S5 Bioglass® is about0%,9 so that even at the highest sintering temperatures theaterial remains a glass–ceramic. Densification regains pace at

round 850 ◦C, when the residual phosphate-rich glassy phaseoftens and progresses appreciably with further increasing sin-ering temperature. Indeed, porosity at the rod surfaces is almostone at 1000 ◦C (bottom-left image in Fig. 7), although thereemains a significant level of internal (closed) porosity (Fig. 2).p to this sintering temperature, the rod surface microstruc-

ure evidences the presence of crystalline grains homogenouslymbedded in a glass matrix. However, the intergranular glassyhase seemingly disappeared from the surface of the rods uponncreasing the sintering temperature to 1050 ◦C, or the hold-ng time at 1000 ◦C to ≥2 h and the naked grains becamepparent (bottom-right image in Fig. 7). These changes were

ccompanied by an apparent increase in surface porosity, whilehe material’s microporosity continued diminishing steadilyith the chronothermic increments, as shown in Fig. 6. This

S. Eqtesadi et al. / Journal of the European Ceramic Society 34 (2014) 107–118 113

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Fig. 9. Evolution of the compressive strength of 45S5 Bioglass®-derived scaf-folds with sintering temperature. Data represent mean values with standarddeviations as error bars. Cancellous bone strength (shaded band)36,37 and resultsf 10,19

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ig. 8. Representative compressive stress–strain curves of 45S5 Bioglass®-erived scaffolds sintered at the indicated temperatures for 1 h.

uggests that the glassy phase—with a very low viscosity solose to the melting temperature—might be draining from theurface towards the interior of the rods and aiding in the reduc-ion of the aforementioned closed porosity.

.3. Mechanical response of sintered scaffolds

Fig. 8 shows representative stress–strain curves of scaffoldsintered at the indicated temperatures for 1 h that were testednder uniaxial compression in the direction perpendicular tohe deposition plane. The scaffolds show an elastic responselmost until the compressive strength of the structure (i.e. theaximum stress) is reached. At this point, the first longitu-

inal cracks pops-in, breaking the unsupported rod segmentslose to the joints with adjacent layers, where maximum tensiletresses are located,34,35 and the applied stress drops. As crack-ng multiplies and propagates all over the scaffold, the stressontinues decreasing progressively down to zero with someventual temporary increases associated to a certain densifica-ion of the fractured structure under the applied compression,hich is typical on cellular structures.The evolution of the compressive strength of robocast 45S5

ioglass®-derived scaffolds with the sintering temperature ishown in Fig. 9. It is evident, that strength increases steadily andtrongly with sintering temperature, but there is no clear corre-ation of strength data with rod density (porosity is indicated as

label beside each data point in Fig. 9) or rod diameter, and theifferent stages observed in the data for the latter (Figs. 5 and 6)re not so clearly observed in Fig. 9. Especially indicative of thisack of complete correlation is the fact that while densificationas improved in scaffolds sintered at 1050 ◦C, there is a statis-

ically significant (p < 0.05) reduction in strength compared totructures sintered at 1000 ◦C (from 13 ± 1 MPa to 11 ± 1 MPa).

his reduction is attributed to the deleterious change in thetarting flaw population at the rod surfaces, where crackingnitiates,34,35 upon the aforementioned disappearance of the

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or scaffolds from other authors are included for comparison. The totalorosity of each structure is indicated as a label beside each data point.

lass matrix (bottom-right image in Fig. 7). This result empha-ises that surface defects are to be minimized whenever possibleince their negative influence on scaffold strength can counteractnd even supersede the positive effect of a reduction in the totalorosity of the structure. Thus, surface flaw population effects,nd maybe some toughness variations associated to the differenticrostructural and phase-compositional changes occurring in

he material at different sintering temperatures, might be respon-ible for blurring the correlation between strength data andorosity/shrinkage results. In any case, it is evident from theseesults that, at least in terms of mechanical performance, the opti-al condition for sintering 45S5 Bioglass®-derived scaffolds

abricated by robocasting is 1000 ◦C for 1 h.On the other hand, as shown in Fig. 9, when comparing

hese strength results with existing literature reports10,19 andancellous bone properties it becomes evident that scaffoldsroduced by robocasting exhibit strength values far superioro anything previously reported for 45S5 Bioglass®-derivedcaffolds—under the same sintering conditions, the enhance-ent is well over one order of magnitude (∼4000%). The reason

or the enhancement of mechanical resistance over conventionalcaffolds fabricated by foam replication10 should be found onhe more regular pore architecture and thicker struts provided byobocasting. More surprising is the fact that other solid freeformabrication methods such as stereolithography (STL),19 whichn principle provide the same advantages, have not been ableo duplicate these strength results even in scaffolds with lowerorosities. Nonetheless, as the authors of the latter report pointut, the reason for this may lay in a less than optimal selectionf the sintering process.19 45S5 Bioglass® is difficult to sin-er, especially into complex scaffold geometries, and great careeeds to be taken in selecting the thermal treatment parameters to

void deleterious cracking during sintering. This must be espe-ially critical in the case of STL, where larger amounts of organicinders (compared to robocasting) have to be burnt out in the

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14 S. Eqtesadi et al. / Journal of the Euro

rocess. In any case, all robocast 45S5 Bioglass®-derived scaf-olds sintered between 550 ◦C and 1000 ◦C fall within or evenurpass the range 2–12 MPa frequently quoted for cancellousone.36,37 Consequently, robocasting provides, for the first time,

means to produce mechanically competent Bioglass® (not justioglass®-derived) scaffolds. Indeed, the robocast scaffolds sin-

ered at 550 ◦C remain completely amorphous and exhibit aompressive strength that, while barely improved over that ofs-dried green structures, lies at the lower end of cancellous boneange and is enough for safe handling during implantation. Thisould help overcome the main hurdle for the successful applica-ion of 45S5 Bioglass® as a broad-use bone substitute material:he difficulty to produce scaffold with sufficient mechanicalntegrity without reducing the outstanding bioactivity of amor-hous 45S5 particles.3

At this point it is worth discussing some issues regardingorosity. Although the total porosities reported here (∼60% and80% for scaffolds sintered at 1000 ◦C and 550 ◦C, respec-

ively) are lower than typical values of conventional 45S5ioglass®-derived scaffolds (90–95%),10 conventional (foam

eplication) scaffolds with similar porosities as those reportedere will still fall well short of robocast scaffolds in term oftrength (as can be easily estimated by extrapolating reportedata10). Moreover, there is ample evidence that, while suchxtremely high porosities might be required to attain signifi-ant bone ingrowth in the tortuous architectures of conventionalcaffolds, porosities such as those reported here are enougho allow bone regeneration and blood vessel growth in thease of calcium-phosphate robocast structures.15–17 In any case,obocast scaffold porosities could be easily increased to matcheported values by modifying the initial design, for example byncreasing rod-spacing within each deposition layer. More initro, in vivo and degradation studies are nonetheless warrantedo assess the full potential and possible limitations of these regu-ar pore architectures with thicker struts produced by robocastings compared to more conventional trabecular-like geometries.

Finally, the authors would like to point-out that the mechan-cal performances of the robocast 45S5 Bioglass®-derivedcaffolds both in terms of strength and toughness could be fur-her enhanced by impregnating them with suitable biodegradableolymers,38–41 although this might come at the expense of reduc-ng bioactivity.49 This might prove especially useful in the casef the less resistant but more bioactive glassy scaffolds sinteredt 550 ◦C.

. Conclusions

Robocasting technique provides a means to produce 45S5ioglass®-derived scaffolds with customized external geome-

ry and internal pore architecture with compressive strengthshat are far superior to any previously reported values. This ischieved through the use of carboxymethyl cellulose (CMC) as

single, multifunctional additive that allows one to overcomehe inherent difficulties associated to the colloidal processingf 45S5 Bioglass® and prepare the concentrated suspensionsequired by robocasting.

agfit

Ceramic Society 34 (2014) 107–118

The enhancement in mechanical performance associated tohe novel pore architectures, with thicker struts and wider inter-onnections, produced by robocasting is enormous: more thanne order of magnitude compared to existing scaffold strengthata and within or even surpassing cancellous bone strengthange. This enhancement allows one to produce, by sinter-ng at sufficiently low temperature (550 ◦C), fully amorphousobocast 45S5 Bioglass® scaffolds, which preserve intact allhe outstanding bioactivity of 45S5 Bioglass®, with enough

echanical integrity for practical use. Therefore it could beaid that robocasting, and perhaps also other solid freeformabrication techniques, can overcome the main obstacle for theuccessful application of 45S5 Bioglass® as a broad-use boneubstitute material. The simplicity and versatility of the pro-osed robocasting process will greatly facilitate the developmentf patient-specific 45S5 Bioglass® scaffolds and significantlyxtend the current range of biomedical applications of this excel-ent bioactive material.

cknowledgments

The authors would like to thank Prof. Angel Luis Ortiz foris assistance in the analysis of XRD data. This study wasupported by the Junta de Extremadura and FEDER Fundsnder grant IB10006. The support from Department of Materi-ls Engineering and Ceramics, CICECO, University of Aveiro,s highly appreciated. Siamak Eqtesadi thanks the Universidade Extremadura for his grant (Ref. 1105).

ppendix.

ptimization of robocasting ink composition and rheology

In the optimization of the inks to be used in robocasting,ifferent polyelectrolyte dispersing agents were tested: (i) aub-group of anionic nature: ammonium polymethacrylate (Dar-an C; mol. wt. = 2000; R.T. Vanderbilt Co., Norwalk, CT);oly (methylvinyl ether) (EasySperse; ISP, Wayne, NJ); Targon128 (BK Ladenburg GmbH, Ladenburg, Germany); and (ii) aub-group of cationic ones: tetramethylammonium hydroxideTMAH, Moses Lake Industries, Inc., USA); PolyethyleneiminePEI, Low Mw, Aldrich, USA).

Zeta potential (ζ) measurements as a function of the pHchanged by adding drop by drop appropriate amounts of 0.1 MCl to the suspension) were performed for the 45S5 Bioglassarticles dispersed in water in the absence and in the presencef 1 wt.% (relative to solid content) of each different dispersantsing a Zeta-potentiometer (Zetasizer Nano ZS, Malvern, UK).t each pH value, an excessive number (100) of measurementsas set, and the instrument automatically stopped when a con-

tant zeta-potential value was measured, typically after 20–25cans.

The preparation of colloidal suspensions using different

dded amounts of each of the abovementioned dispersants ran-ing from 0.5 to 2 wt.% was attempted. Anionic dispersants wererstly attempted for ink preparation because, according to the

raditional approach, they are expected to adsorb at the surface of

S. Eqtesadi et al. / Journal of the European Ceramic Society 34 (2014) 107–118 115

Table A1Maximum solid loadings achieved in aqueous 45S5 bioglass suspensions with different added amounts of the various polyelectrolyte dispersants tested.

Dispersant Type Maximum achievable solid loadings (vol.%)

Dispersant concentration (wt.%) 0.5 1.0 1.5 2.0

Darvan C Anionic 29.8 31.9 33.8 35.3EasySperse Anionic 28.1 29.2 31.5 33.1Targon 1128 Anionic 22.5 23.7 25.1 27.8TMAH Cationic 21.2 22.5 23.8 26.5P

tleso(w(

acctstsiCjζ

npsw

Fptc

atto

tcsotepitwTitaas

EI Cationic 19.5

he solid particles and stabilize the suspensions, which can in aater stage be coagulated by adding cationic additives.26,42 How-ver, as shown in Table A1, a maximum of 35 vol.% solids loadeduspensions could be hardly achieved (only in the presencef 2 wt.% Darvan C), which is somewhat low for robocastingoptimal would be around 40–45 vol.%). Even worse resultsere obtained using cationic polyelectrolytes as dispersants

Table A1).The explanation for these difficulties can be found by looking

t the ζ-potential results plotted in Fig. A1. For all dispersiononditions other than in the presence of PEI, the ζ-potentialurves show a continuous decrease from a common isoelec-ric point at around pH 1.5—close to the values reported forilica,18 suggesting that 45S5 glass particles exhibit a silicaype surface—up to a pH of about 5–6 where the trend invertedlightly to again resume its drop after reaching a relative max-ma at about pH 8. The intermediate relative maxima imply thata2+, CaOH+ ions leached from the 45S5 Bioglass, and not

ust H+ and OH− ions, are playing a role in the determination of-potential.21 The fact that the electrophoretic curves are not sig-ificantly shifted (except in the case of PEI) from that of the bare

articles suggests that there is not a significant adsorption of theelected polyanions on the surface of 45S5 Bioglass® particles,hich severely limits their effectiveness as dispersants. This is

ig. A1. Zeta-potential of 45S5 Bioglass® particles versus pH. Data for barearticles (hollow symbols) and after the addition of a fixed amount (1 wt.%) ofhe indicated anionic (dark symbols: EasySperse; Targon 1128; Darvan C) orationic (light symbols: TMAH and PEI) dispersants are included.

p(wap

csrweppamnocowtIbff

22.2 23.4 26.7

ttributed to electrostatic repulsion between anionic polyelec-rolytes and the negatively charged 45S5 bare particles, and tohe fact that TMAH ions are not prone to be specifically adsorbednto silica-like surfaces.43

On the other hand, the strong upwards shift in the elec-rophoretic curve to positive potentials in the presence of theationic PEI proves that PEI species do specifically adsorb at theurface of the 45S5 Bioglass® particles, as previously reportedn similar silica-like surfaces.44 However, the resulting isoelec-ric point (pH ∼ 10) coincides with the pH of the dispersionnvironment created by the 45S5 glass powder in water—indeed,H augments from 6 to 10 within the first 30 s due to Na+/H+

onic exchange.45 This implies the absence electrostatic interac-ions among the particles under these conditions, and the systemill tend to flocculate under the action of Van der Walls forces.his cannot be avoided by lowering the pH of the dispers-

ng medium since this has the counter effect of acceleratinghe 45S5 Bioglass® dissolution, which makes pH to increasegain.46–48 Moreover, considering that anionic/cationic inter-ctions of species added to the system in different successiveteps are the base of the most commonly used mechanism toromote the rheological transition from a fluid-like behaviourstarting suspension) to a paste-like character (usable in direct-rite assembling techniques), it became clear that this traditional

pproach could not offer successful prospects for robocastingrocessing of 45S5 Bioglass®.

Based on all the above considerations, we have adopted aompletely different and simplified approach: searching for aingle additive that could act as both dispersant and binder forobocasting. In this pursue, carboxymethyl cellulose (CMC)as regarded with interest due to its aptitude to play differ-

nt but complementary roles as processing additive.21 The lessolar character of the CMC chains in comparison to the anionicolyelectrolyte molecules facilitates surface adsorption. Suchdsorption is likely to occur, however, through small chain seg-ents, with laces and tails extending to the solution, since it is

ot electrostatically driven (both CMC molecules and the surfacef the particles have the same negative polarity). The predictedonfiguration of the CMC molecules adsorbed at the surfacef the 45S5 Bioglass® particles is likely to play a steric effecthich enabled reaching significantly higher solid loadings, up

o 45 vol.%, compared to anionic polyelectrolytes (Table A1).n particular, high molecular weight CMC (M = 250,000, Lam-

w

erti Iberia s.a.u., Castellón, Spain, with a viscosity η = 1–3 Pa sor a 2 wt.% aqueous solution) was found more appropriateor the preparation of inks for robocasting since longer chains

116 S. Eqtesadi et al. / Journal of the European Ceramic Society 34 (2014) 107–118

Fig. A2. Plots of (a) shear stress and (b) apparent viscosity as a function of shearrate for 45 vol.% 45S5 Bioglass® suspensions with indicated amounts of CMCas binder.

ib

dem1g((vtatdm

Fig. A3. Effect of the amounts of CMC on the yield stress, τy, and the averageelastic modulus at low stress (0.1–1 Pa), G′, of 45 vol.% 45S5 Bioglass® sus-ps

w

τ

wspsmwpasesasrspsaaaattinalatffurther leads to a higher degree of surface coverage and a denseradsorbed layer, which decreases the probability of one poly-meric chain to adsorb in more than one particle, thus hindering

nduced a stronger thickening effect, which was attributed to aridging flocculation mechanism.21

The flow properties of 45 vol.% suspensions prepared usingifferent amounts of this CMC as the sole binding agent werevaluated using a rheometer (Bohlin CVO, Malvern, UK). Floweasurements were conducted at shear rates between 0.2 and

00 s−1 using a cone and plate configuration (4◦, 40 mm, andap of 150 �m). Fig. A2 shows the shear stress versus shear rateFig. A2a) and the apparent viscosity versus shear rate curvesFig. A2b). All systems exhibit an evolution of shear stressersus shear rate that approaches the Bingham plastic behaviour,ypical of suspensions containing polymeric binder/thickeninggents. The polymer molecules help forming a network struc-ure under rest or near rest conditions, which is then graduallyisrupted under the applied shear stress field as the polymericolecules align along the flow direction.

ensions. G′ data were calculated from previous results.21 Error bars representtandard deviations (from fitting in the case of τy).

The initial data plotted in Fig. A2a (up to 3 s−1) were fittedith the Herschel–Bulkley model,29 which is described as:

= τy + K(γ0)n (3)

here γ0 is the shear rate, τ is the shear stress, τy is the yieldtress, n is the shear thinning exponent and K is the viscosityarameter. The yield stress is the shear stress in the absence ofhear motion and represents the minimum stress leading to defor-ation of the static form of the material. The model parametersere estimated as n = 0.66 ± 0.06 and K = 580 ± 50 Pa sn for sus-ensions containing up to 1 wt.% of CMC and n = 0.36 ± 0.03nd K = 2600 ± 200 Pa sn for samples with 2 wt.%. Fig. A3hows the evolution of the estimated yield stress and the averagelastic (storage) modulus, G′, at low shear stress of suspen-ions with added CMC. Modulus values were calculated as theverage of the virtually constant G′ data within the 0.1–1 Patress range from previous work.21 These parameters reflect,espectively, the strength and stiffness of the attractive networktructure resulting from interactions between 45S5 Bioglass®

articles, CMC and water. In the absence of CMC, both the yieldtress and the modulus of the suspension are very low. Upon theddition of CMC, the yield stress increases linearly with themount of binder up to a concentration of 1 wt.%, decreasingt 2 wt.%. This thickening effect promoted by CMC might bettributed to chain entanglement of the molecules adsorbed athe surface of the particles and/or to an increased viscosity ofhe dispersing medium. On the other hand, the elastic modulusncreases first very slowly, then dramatically (by 2 orders of mag-itude) at 1 wt.% CMC additions, and finally decreases slowlyt higher CMC concentrations. The dramatic change in modu-us observed for 1 wt.% CMC indicates a fluid-to-gel transitionttributed to the occurrence of bridging-flocculation, in whichhe longer polymeric chains might adsorb at the surface of dif-erent neighbouring particles. Increasing the amount of CMC

pean

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1

1

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ridging-flocculation and reducing both the modulus and theield stress. Since high values of this two parameters are desiredeatures for robocasting inks, 1 wt.% was considered as theptimal CMC amount. This concentration yielded maximumetwork stiffness (>100 kPa), and yield stress (>1 kPa) at a rel-tively low viscosity (Fig. A2), which facilitates flowabilityhrough the deposition nozzle. Indeed, such suspensions exhib-ted both very good flowability (unlike suspension with higherMC concentrations) and excellent shape retention capacitypon deposition, which made them ideal inks for direct writingechniques.

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