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Processing strategies for smart electroconductive carbon nanotube-based bioceramic bone

grafts

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2014 Nanotechnology 25 145602

(http://iopscience.iop.org/0957-4484/25/14/145602)

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Nanotechnology

Nanotechnology 25 (2014) 145602 (14pp) doi:10.1088/0957-4484/25/14/145602

Processing strategies for smartelectroconductive carbon nanotube-basedbioceramic bone graftsD Mata1, F J Oliveira2, N M Ferreira1,2, R F Araújo3, A J S Fernandes1,M A Lopes4, P S Gomes5, M H Fernandes5 and R F Silva2

1 I3N, Physics Department, University of Aveiro, 3810-193 Aveiro, Portugal2 CICECO, Materials and Ceramic Engineering Department, University of Aveiro, 3810-193 Aveiro, Portugal3 Centre of Chemistry, University of Minho, 4710-057 Braga, Portugal4 CEMUC, Metallurgical and Materials Engineering Department, Faculty of Engineering,University of Porto, 4200-465 Porto, Portugal5 Laboratory for Bone Metabolism and Regeneration, Faculty of Dental Medicine, University of Porto,4200-465 Porto, Portugal

E-mail: diogomata@ua.pt

Received 26 September 2013, revised 23 January 2014Accepted for publication 5 February 2014Published 12 March 2014

AbstractElectroconductive bone grafts have been designed to control bone regeneration. Contrary topolymeric matrices, the translation of the carbon nanotube (CNT) electroconductivity into oxideceramics is challenging due to the CNT oxidation during sintering.

Sintering strategies involving reactive-bed pressureless sintering (RB + P) and hot-pressing(HP) were optimized towards prevention of CNT oxidation in glass/hydroxyapatite (HA) matrices.Both showed CNT retentions up to 80%, even at 1300 ◦C, yielding an increase of theelectroconductivity in ten orders of magnitude relative to the matrix. The RB + P CNT compactsshowed higher electroconductivity by ∼170% than the HP ones due to the lower damage to CNTsof the former route. Even so, highly reproducible conductivities with statistical variation below 5%and dense compacts up to 96% were only obtained by HP.

The hot-pressed CNT compacts possessed no acute toxicity in a human osteoblastic cell line.A normal cellular adhesion and a marked orientation of the cell growth were observed over theCNT composites, with a proliferation/differentiation relationship favouring osteoblastic functionalactivity.

These sintering strategies offer new insights into the sintering of electroconductive CNTcontaining bioactive ceramics with unlimited geometries for electrotherapy of the bone tissue.

Keywords: carbon nanotubes, ceramic matrix composite, reactive sintering, hot-pressing,electroconductive bone graft, in vitro osteocompatibility

S Online supplementary data available from stacks.iop.org/Nano/25/145602/mmedia

(Some figures may appear in colour only in the online journal)

1. Introduction

Calcium phosphate (CaP) bioceramics have been widelyused in bone surgery as synthetic grafts for more than

three decades [1–3]. Their success is owed to the chemicalresemblance with the mineral component of the naturalbone [3], although all of these family-related CaPs presentpoorer mechanical properties than natural bone [3]. At load

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Nanotechnology 25 (2014) 145602 D Mata et al

bearing sites, this represents a concern given the loss of thematerial structural integrity at the first stage of the new bonegrowth. To improve this, CaPs are usually reinforced with otherphases, for example: bioglass reinforced hydroxyapatite—Glass/HA. Besides the improved mechanical properties [4],this composite shows higher bioactivity [5] than single phaseHA. Other attractive phases, such as carbon nanotubes (CNTs),have been recognized to strengthen CaP ceramics [6] due totheir high biocompatibility with bone tissue [7], bioactivity [8],low cost and outstanding mechanical properties [9]. Despitegreat efforts, the mechanical properties of the composites areusually only marginally better than those of the ceramic alone.

Apart from the prospective mechanical and biological en-hancement, CNTs can also be included within the compositionof scaffolds to confer on them new properties, such as electricalconductivity. CNTs present ultimate electrical conductivity [9]allowing electroconductive biomaterials at low percolationthresholds (Pc) to be obtained without damaging the biologicalprofile of the CaP matrix. Moreover, non-metallic reinforcingphases like CNTs establish a breakthrough in the production ofhigh electroconductive biomaterials to work under electricalstimulation routines, free of corrosion related toxicologicalrisks.

Electrical stimulation, broadly attained by capacitive cou-pled electrical field, direct current field and electromagneticfield, has been used to enhance the bone healing process[10, 11]. In vitro, the direct application of electrical currenthas been found to induce cell elongation, modulate cell align-ment, and to favour the migration and invasion of humanmesenchymal stem cells (hMSCs) into injured sites [12].Electrical stimulation also seems to enhance the cell pro-liferation and the osteogenic differentiation of osteoblasticprecursor cells [13, 14]. These findings seem to substantiatethe improved musculoskeletal healing attained with electricalstimulation in clinical trials aimed at different applications ofbone repair/regeneration, as reviewed in [15].

These data have inspired the development of conductivesynthetic matrices to assist in the bone regeneration processunder electrical stimulation. This is particularly useful sinceelectric/electromagnetic strategies are spatially non-delimited,indiscriminately affecting both targeting and non-targetinganatomical locations. Ideally for clinical application, the elec-trical stimuli should be confined within the tissue volumeexpected to be regenerated. As such, CNT-loaded bioactivecomposites may be considered as a new generation of ‘smart’materials, which are able to deliver confined electrical stim-uli to the damaged bone site in a spatially and temporallycontrolled manner and to tune particular responses of theelectrically excitable new bone tissue [16]. In vitro, the elec-trical stimulation of bone cells on CNT–polymeric substrateshas been found to accelerate cell proliferation and osteogenicdifferentiation [17, 18]. Yet, polymers present much lowerbone grafting qualities than CaPs. These ceramics are alsodielectric [19], but studies investigating the use of CNTs toimprove their electrical properties are not found. By combiningthe biological profile and mechanical strength of the Glass/HAceramic, with the morphology and electrical conductivity of

the CNTs, a superior multi-functional bone graft is proposed—CNT/Glass/HA, inspired by the respective apatite-like phaseand the collagen type I fibres of the natural bone.

Contrary to polymeric composites, the translation of theCNT properties into bulk ceramics is not straightforward.Particularly for oxide matrices, a major concern deals with thesintering step at high temperatures. In H2O-free atmospheres,at 680–1080 ◦C, the HA phase begins to lose the structuralOH groups, i.e. the dehydroxylation occurs [20]. The releaseof oxygen-containing groups from the matrix, above theoxidizing temperature of the CNTs, at 500–600 ◦C, may fullyoxidize or severely damage them. This becomes a real concernwhen the maximization of the CNT electrical properties isenvisaged.

Concerning this, the selection of a suitable sintering routeis the first challenge to overcome in the processing of CNT–HA formulations. Such composites have been consolidatedby pressureless or pressure-assisted sintering techniques [6].When compared to the latter ones, the former has the advantageof being less destructive to CNTs [21] and of not being geo-metrically limited or having an expensive technique. The maindisadvantages of the absence of pressure are the lack of fulldensification and total prevention of the dehydroxylation andsubsequent oxidizing reactions in most CNT–HA systems [22].Even so, a recent work reports a pressureless sintering with agood balance of CNT loading and densification [23]. Therewas however no data on the electrical characterization of thecomposites or on the crystallinity of the CNT structure in theprocessed materials.

This work focuses on the optimization of the sinteringroute of highly electroconductive CNT/Glass/HA syntheticbone grafts. The powders were consolidated under reductiveatmospheres by different routes: simple pressureless sinter-ing (P), reactive powder bed under pressureless sintering(RB + P) and hot-pressing (HP). To assess the efficiency ofthe sintering routes, the loading and structural crystallinityof CNTs, after sintering alongside the compressive strengthand electrical conductivity of the composites, were thoroughlyinvestigated. Also, the balance of CNT loading and densi-fication were maximized to preserve the mechanical prop-erties of the matrix and achieve a 3D electron transportingnetwork. The in vitro biological response of the attainedcomposites was further assessed with human-osteoblastic-likecells in terms of cell viability, proliferation and functionalactivity.

2. Experimental details

2.1. Preparation of CNT/Glass/HA powders

Bioglass (65P2O5, 15CaO, 10CaF2, 10Na2O mol%) and HApowders were lab-prepared from high purity (>98%) gradereagents: the glass was formed at 1450 ◦C for 90 min in aplatinum crucible then rapidly cooled to room temperature;HA was synthesized by a conventional precipitation method.A full description of these processes was reported previously[4, 5]. The Glass and HA materials were then individuallyplanetary milled in ethanol (≥99.9%, Sigma) with agate jar

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Nanotechnology 25 (2014) 145602 D Mata et al

and balls until reaching a particle diameter size of D0.5 = 3.1±1.6 µm and D0.5 = 2.2± 1.1 µm, respectively. Subsequently,Glass/HA powders with 2.5 wt% of glass were planetaryco-milled/mixed in the same conditions for 24 h, giving afinal particle diameter size of D0.5 = 1.8± 1.4 µm.

MWCNTs (NC7000) were provided by Nanocyl with thecharacteristics presented in table S1 (see supporting mate-rial available at stacks.iop.org/Nano/25/145602/mmedia). TheCNTs were functionalized by a Diels–Alder cycloadditionreaction of 1,3-butadiene upon heating at 150 ◦C for 7 days.Further details of the process are presented elsewhere [24].

Isopropyl alcohol (≥99.8%, Sigma), used previously[4, 5], was also applied for thorough mixing the compositepowder with a fixed CNT loading of 2.5 wt% (4.4 vol%)—CNT/Glass/HA composites. The powders were processed bya two-stepped method combining high-speed shearing for30 min (IKA T25-Ultra-Turrax, working at 20 500 rpm) andsonication during 60 min (Selecta, working at 60 kHz, 200 W).After fast solvent evaporation under vacuum at 80 ◦C, to avoidphase separation, the dried powders were crushed in an agatemortar and sieved to less than 75 µm.

2.2. Sintering routes

For pressureless sintered (P) powders, disc-shaped sampleswith a diameter of 5.5 ± 0.1 mm diameter and thickness of1.2± 0.1 mm were prepared by isostatic pressing at 300 MPa.Green compacts were further fired in a open Al2O3 crucibleinside a graphite furnace under vacuum (3 Pa) or Ar (70 kPa,purity> 99.9999%), as follows: heating rate of 4 ◦C min−1

with a fixed dwelling time of 60 min at 1300 ◦C, followedby cooling at rate of 10 ◦C min−1 to room temperature. Whenusing a reactive-bed pressureless sintering (RB+ P) approach,only Ar atmospheres were applied, and the open crucible wasreplaced by a closed graphitic crucible filled with a bed ofHA+ carbon black (≥99.5%, Degussa) powder mixture. Thedegree of HA dehydroxylation was controlled by varying thesintering temperature (1200, 1300 ◦C), the total amount of thepowder (5–30 g) and the carbon loading (10–50 wt%). Thediameter and thickness of P and RB + P sintered compactswere in the range of 4.5–5 mm and 1–1.1 mm.

For hot-pressing (HP), the powder mixture was placedin a 2 cm diameter graphite die and preliminarily pressed atfixed pressure of 30 MPa. Then, the chamber was pumpeddown to ∼3 Pa, and subsequently, the samples were firedby the following HP cycle: heating rate of 10 ◦C min−1

with a fixed dwelling time of 60 min at temperature rang-ing from 1000 to 1300 ◦C, followed by cooling at rate of10 ◦C min−1 to room temperature. The final compacts aredisc-shaped samples 20 mm in diameter and 3 ± 1 mm thick.Finally, the discs were cut in two geometries: parallelepipedof 2.5× 2.5× 4.5 mm3 for compression testing and electricalmeasurements and square slices of 5× 5× 1 mm3 (surfacefinishing: P4000) for cell cultures. For in vitro testing, thematerials were ultrasonically cleaned in ethanol and sterilizedby autoclaving.

2.3. Materials characterization techniques

Thermogravimetric analyses (TGA, Setsys Setaram) were per-formed under a constant O2 flow of 200 sccm at 10 ◦C min−1.Advantageously, the TGO2 analysis allowed the precise de-termination of the CNT wt% in the composite formulations.These values are derived from the step parallel tangents (Tang1 and Tang 2) to the TGA curve near the known CNT oxidationtemperature as follows:

CNT%=WeightTang1%−WeightTang2%. (1)

The CNT loss% for a sintered CNT composite wasdetermined regarding the CNT% before sintering (CNTBS)

and after sintering (CNTAS) according to,

CNT loss%=CNTBS%−CNTAS%

CNTBS%× 100. (2)

The appearance of the sintered samples was also recordedto further confirm the CNT loading.

Micro-Raman spectroscopy (Jobin Yvon T64000) wasconducted at 532 nm in the wavenumber range of1000–2000 cm−1 to assess the structural crystallinity of CNTsbefore and after sintering. The CNT crystallinity was estimatedby calculating the graphite in-plane crystallite size (La) usingthe Cancado equation [25].

The phase composition of sintered composites was as-sessed by x-ray diffraction (X’PERT-MPD Philips) workingwith a Cu Kα1 radiation (λ = 0.154 056 nm), from 5◦ to80◦ with a stepped size of 0.02◦. The Gazzara and Messiermethod [26] was applied for phase quantification, weightpercentage (wt%), by considering the integrated intensities,I (hkl), of the most representative planes of each phase:HA (211); HA(112); HA(300); α-TCP(150); α-TCP(132);α-TCP(170); β-TCP(214); β-TCP(300); β-TCP(0210).

The weight percentage of phase A is given by

phase A%=ρA

∑3i=1 IA(hkl)i

3∑n phasesj=A (ρ j

∑3i=1 I j (hkl)i

3 )

× 100 (3)

where ρ is the density (kg m−3).The experimental (ρexp) and theoretical (ρtheo) densities

of sintered composites were respectively obtained by the prin-ciple of Archimedes in ethylene glycol (three measurementswere conducted on each sample) and by the rule of mixtures byconsidering the theoretical densities of the following phases:HA (d = 3.16 g cm−3 [27]), β-TCP (d = 3.07 g cm−3 [27]),TTCP (d = 3.05 g cm−3 [27]), α-TCP (d = 2.86 g cm−3 [27]),Bioglass (d = 2.7 g cm−3 [28]) and MWCNT-NC7000 (d =1.66 g cm−3 [29]).

The porosity% was calculated using the ρexp, the ρtheoand the phase percentages obtained by the XRD analysesaccording to

porosity%=

1−ρexp∑n phases

j=1 (phase j %ρtheo j )

× 100. (4)

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Nanotechnology 25 (2014) 145602 D Mata et al

The microstructure and composition of sintered samples,polished and chemically etched, were characterized by highresolution scanning microscopy (HR-SEM, Hitachi SU70working at 15 keV with a resolution of 1 nm) assisted byenergy dispersive spectroscopy (EDS, Bruker Quantax 400).The samples were firstly polished with SiC papers (gradesP1000, P1200, P2500, P4000) followed by a 50 nm sizedsilica finishing. The chemical etching was accomplished with0.01 M HNO3 for 10–60 s, with two etching strengths, to showthe CNT locations in the microstructure and reveal the grains.The grain size was determined by Heyn’s method, accordingto the ASTM E112-196.

DC electrical conductivity measurements at room tem-perature of ceramic samples were performed in two differentapparatus: a Keithley 617 programmable electrometer with astepped applied voltage of 0.5 V in the range of 0–100 V fordielectric samples; ISO-TECH IPS-603 programmable powersupply with a stepped applied voltage of 0.1 V in the rangeof 0–1 V for conductive samples. For the parallelepiped HPsamples the copper wires were fixed directly to both ends ofeach sample. For the disc-shaped P and RB + P samples,the contact diameter was reduced to 2.2 mm by masking thesample and evaporating a Au/Pd film on each side of thesample.

Compressive tests were carried out in a Zwick/Roell Z020equipment with a load cell of 2 kN and a constant displacementrate of 0.3 mm min−1. All mechanical and electrical tests wereperformed with five specimens for each condition.

2.4. Cell culture studies

Human-osteoblastic-like cells (MG63 cell line, ATCC numberCRL-1427TM) were cultured in α-Minimal Essential Medium(α-MEM), supplemented with 10% foetal bovine serum,50 µg ml−1 ascorbic acid, 50 µg ml−1 gentamicin and2.5 µg ml−1 fungizone, at 37 ◦C, in a humidified atmosphereof 5% CO2 in air. At around 70-80% confluence, the cell layerwas detached with trypsin—EDTA solution (0.05% trypsin,0.25% EDTA; 5 min, 37 ◦C). The cell suspension was seeded(5× 104 cells cm−2) over standard cell culture cover slips(control), Glass/HA composite and CNT/Glass/HA composite,and cultures were incubated for 5 days in the experimentalconditions described above. Cultures were characterized atdays 1, 3 and 5.

2.5. Characterization of the cell response

DNA content was analysed by the PicoGreen DNA quantifi-cation assay (Quant-iTTM PicoGreen R© dsDNA Assay Kit,Molecular Probes Inc., Eugene), according to the manufac-turer’s instructions. Cultures were treated with Triton X-100(0.1%) (Sigma) and fluorescence was measured on a Elisareader (Synergy HT, Biotek) at wavelengths of 480 and520 nm, excitation and emission respectively, and correctedfor fluorescence of reagent blanks. The amount of DNAwas calculated by extrapolating a standard curve obtainedby running the assay with the given DNA standard. Resultsare expressed as a percentage of variation from control.

The MTT assay was used to assess the cell viabil-ity/proliferation. This is based on the reduction of the MTT(3-(4,5-Dimethylthiazol-2-yl)-2,5-diphenyltetrazolium) by vi-able cells to a dark blue formazan product. MTT (0.5 mg ml−1)

was added to each well, and cultures were incubated for 3 hat 37 ◦C. Subsequently, the formazan salts were dissolvedin dimethylsulfoxide (DMSO) and the absorbance (A) wasdetermined at λ= 600 nm on an Elisa reader (Synergy HT,Biotek). Results are expressed as a percentage of variationfrom control.

The lactate dehydrogenase (LDH) assay was used toanalyse cell viability by assessing the cell membrane integrity.This test is based on the reduction of NAD by the action ofLDH released to the medium by damaged cells. The result-ing reduced NAD (NADH) is utilized in the stoichiometricconversion of a tetrazolium dye. Determination of the totalLDH was performed using the in vitro toxicology assay kitlactate dehydrogenase based (Sigma-Aldrich; St. Louis, MO),according to the manufacturers’ instructions. The amount ofLDH leakage to the medium was normalized by total LDH,and calculated as follows:

LDH leakage%=LDH medium

Total LDH× 100. (5)

Results are expressed as a percentage of variation from control.Alkaline phosphatase (ALP) activity was evaluated in

cell lysates (0.1% Triton X-100, 5 min) by the hydrolysisof p-nitrophenyl phosphate in alkaline buffer solution (pH∼10.3; 30 min, 37 ◦C) and colourimetric determination ofthe product (p-nitrophenol) at 400 nm in an ELISA platereader (Synergy HT, Biotek). ALP activity was normalizedto total protein content (quantified by Bradford’s method).Results are expressed as a percentage of variation fromcontrol.

For SEM observation, samples were fixed (1.5% glu-taraldehyde in 0.14 M sodium cacodylate buffer, pH= 7.3,10 min), dehydrated in graded alcohols, critical-point dried,sputter-coated with an Au/Pd thin film (SPI Module Sput-ter Coater equipment), and observed under a high resolu-tion (Schottky) environmental scanning electron microscope(Quanta 400 FEG ESEM).

For confocal laser scanning microscopy (CLSM) assess-ment, samples were fixed (3.7% paraformaldehyde, 15 min).Cell cytoskeleton filamentous actin (F-actin) was visualizedby treating the cells with Alexa Fluor 488 Phalloidin (1:20dilution in PBS, 1 h) and counterstaining with propidiumiodide (1 µg ml−1, 10 min) for cell nuclei labelling. Labelledcultures were mounted in Vectashield R© and examined under aLeica SP2 AOBS (Leica Microsystems) microscopy.

Three independent experiments were performed; in eachexperiment, three replicas were accomplished for the biochem-ical assays and two replicas for the qualitative assays. Resultsare presented as mean ± standard deviation (SD). Two-wayanalysis of variance (ANOVA) was used in combination withBonferroni’s post-hoc test to data evaluation. Values of p ≤0.05 were considered statistically significant.

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Nanotechnology 25 (2014) 145602 D Mata et al

3. Results and discussion

3.1. Reactive-bed pressureless sintering (RB + P) route

As demonstrated by White et al [23], the retention of HAand CNT in the final composite can be achieved by simplyshifting the equilibrium of reaction H2O+C (equation (S2),see supporting material available at stacks.iop.org/Nano/25/145602/mmedia), according to Chatelier’s principle. Gas flowscontaining CO+H2+H2O mixtures running in a tube furnacewere found to be the most successful [23]. A simpler approachconcerns the formation of a CO+H2+H2O atmosphere in-side a sealed crucible by using a high reactive HA+ C powderbed, as follows: (1) free powder of HA of the bed havinghigher reactivity than the compact powder (see supportingmaterial available at stacks.iop.org/Nano/25/145602/mmedia)starts by decomposing to form H2O vapour (equation (S1),see supporting material available at stacks.iop.org/Nano/25/145602/mmedia); (2) H2O reacts with solid C powder of thebed to produce CO+H2 gas mixtures (equation (S2) availableat stacks.iop.org/Nano/25/145602/mmedia); (3) CO+H2 gasintimately contacts the compact pellets and locally retainsCNTs (by forming carbon, equations (S2) and (S3), see sup-porting material available at stacks.iop.org/Nano/25/145602/mmedia) and the H2O (equation (S2), see supporting materialavailable at stacks.iop.org/Nano/25/145602/mmedia) neededto prevent HA dehydroxylation. RB+ P sintering studies wereperformed at 1200 and 1300 ◦C. The lower temperature wasalso investigated in an attempt to achieve a better balancebetween CNT loading and densification.

Following the thermodynamics of the equation (S2) (seesupporting material available at stacks.iop.org/Nano/25/145602/mmedia), the respective equilibrium constants (K ) at1200 ◦C and 1300 ◦C are 500 and 1000, respectively. Thesevalues were calculated using the Factsage 6.1 software. Con-cerning Chatelier’s rule, the backward reaction of equation(S2) (see supporting material available at stacks.iop.org/Nano/25/145602/mmedia) is promoted when the reaction quotient(Q) is higher than the equilibrium one, Q > K . This can beachieved by increasing either CO, H2 or both species. Bearingthis in mind, HA + C reactive-bed compositions were strate-gically selected to obtain a narrow range of partial pressuresof CO (PCO, where PCO = PH2 ) that follows the condition ofQ > K . As shown in figure 1(a), for each temperature, fourPCO conditions were used where three of them respect thecondition Q > K . For these calculations it was assumed thatthe crucible was completely sealed and the powders were fullyreacted. Also, it is relevant to highlight that these values aretheoretical being relative to equilibrium conditions, yet theygive the closest perspective to the experimental conditions.

Figure 1(b) shows the variation of porosity and HA contentof Glass/HA and CNT/Glass/HA samples with the sinteringtemperature and sintering atmosphere composition (PCO).With respect to the HA fraction, the values for both samplecompositions are almost independent of the temperature usedfalling in the narrow range of 80–100%. Overall, the porosityof the CNT composite decreases with the increasing ofthe sintering temperature from 25–30% to 20–25%, but for

Figure 1. (a) Plot showing the influence of the RB + P atmospherecomposition (PCO) on the direction tendency of the C+H2Oreaction (equation (S2) available at stacks.iop.org/Nano/25/145602/mmedia) by comparing the respective equilibrium constant (K ) andthe reaction quotient (Q) at 1200 ◦C and 1300 ◦C. (b) Plot showingthe dependence of the porosity and HA phase of the Glass/HA andCNT/Glass/HA sintered by P + RB at 1200 ◦C and 1300 ◦C on thePCO of the sintering atmosphere.

the matrix alone the porosity was not significantly varied(10–15%).

Concerning the influence of the atmosphere composition,it can be seen that the HA fraction variation is close for bothsample types, although in Glass/HA samples the HA phasewas not totally retained due to two competitive reactions.Firstly, the matrices alone have a larger amount of HA thanthe CNT/Glass/HA samples, thus, the produced H2O fromthe CO+H2 mixture becomes insufficient to fully retain theHA phase (increase of HA retention for higher PCO values).Secondly, since there are no CNTs retained in the Glass/HAsamples, the solid carbon formed locally by the high reductiveCO (equations (S2) and (S3), see supporting material availableat stacks.iop.org/Nano/25/145602/mmedia) may acceleratethe HA decomposition, as reported elsewhere [30] (increase ofHA retention for lower PCO values). Considering this balance,the Glass/HA samples present a maximum HA retention at asimilar PCO value than the CNT/Glass/HA samples for bothsintering temperatures (figure 1(b)).

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Nanotechnology 25 (2014) 145602 D Mata et al

For the CNT/Glass/HA samples, the HA content increaseswith PCO up to a maximum value of 100% for both sinteringtemperatures. Interestingly, the full HA dehydroxylation sup-pression was achieved when the Q values were approximatelytwo and three times the K values at 1200 ◦C and 1300 ◦C,respectively (figure 1(a)). For higher values of Q the atmo-sphere becomes probably too reductive, accelerating the HAdecomposition.

Overall, the porosity of the Glass/HA samples varieswith PCO in the same trend as the HA fraction. The higherHA decomposition results from a higher carbon impuritycontent which promotes a subsequent decrease in the poros-ity due to the carbon incorporation in the microstructure.For CNT/Glass/HA samples the porosity varies with PCOwith an opposite trend of the HA fraction. Particularly inCNT/Glass/HA samples, this occurs because the higher HAcontent promotes higher CNT retention and subsequent lowerporosity. This can be easily evaluated by visual observa-tion of the appearance of the samples. Figure 2(a) shows acolour mapped ternary graph of the phase composition ofCNT/Glass/HA samples (highlighted with red dots), sintered at1200 and 1300 ◦C and plotted as a function of their appearance,following a three colour grade: white outer layer and light greyinner; uniformly dark grey; uniformly black. Representativephotograph images of the three situations are shown in figures2(b)–(d). It is evident in figure 2(a) that the darkening ofthe samples increases with the increasing HA fraction dueto the large amount of unreacted CNT loading. Concerningthe extreme conditions, the two points in the light grey areacorrespond to compositions of the lower HA fraction obtainedfor sintering atmospheres where PCO = 0, using a 100% HAreactive bed. In this condition, only H2O is produced whichinduces the CNT oxidation across the sample, especially at thesurface, where a white layer can be seen (figure 2(b)). At theother extreme, for samples with HA content higher than 91%,the four points in the black area, the CNT loading is enoughto give uniformly black samples, figure 2(d).

3.2. Sintering efficiency of P, RB + P and hot-pressing (HP)routes

Regarding the purpose of the present work, the balance of CNTloading and densification should be maximized to preservethe mechanical properties of the matrix and to have a highelectroconductive composite. Despite the successful retentionof HA and CNT, the porosity of RB + P sintered samples isstill quite high (figure 1(b)), negatively affecting their physicalproperties. To overcome this, pressure-assisted routes areusually applied with high success, achieving porosity valuesbelow 4% for CNT/HA composites [6]. Also, it was foundthat the applied pressure has a similar role to the reactivebed in suppressing HA dehydroxylation [22, 31–33], althoughthe mechanisms involved in the phase retention are differentfor the two sintering processes. In RB + P sintering themechanism is guided by a local chemical equilibrium while apurely kinetic mechanism is concerned in HP sintering [22].Considering these potentialities, the hot-pressing route wasalso investigated.

Figure 2. (a) Colour mapped ternary graphic illustrating theinfluence of the phase composition (represented with red dots) onthe appearance of CNT/Glass/HA samples fired by P + RB.((b)–(d)) representative photographs of (a).

Figures 3 and 4 summarize data regarding the mor-phology, CNT structural crystallinity and physical propertiesof Glass/HA and CNT/Glass/HA samples consolidated at1300 ◦C by three different routes: pressureless sintering in Ar(P); reactive-bed pressureless sintering in a PCO = 1.77 MPaatmosphere (RB + P); hot-pressing at 30 MPa (HP).

Figure 3 denotes that the reactive-bed sintered CNT/Glass/HA samples have improved their HA retention by ∼16% andsignificantly decreased the CNT loss by ∼75% relative to theP sintering route. The results for the P sintering approach arein accordance with those reported by others [34]. Yet, porosityof RB+ P samples only improved∼7%. The final porosity ofthe composites is not only dependent of the CNT eliminationby H2O oxidation as is shown in section 3.1, but in fact resultsfrom the balance between CNT elimination and CNT retention.The low densification of RB + P CNT/Glass/HA samplesof ∼76% is caused by the high level of CNT in the final

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Nanotechnology 25 (2014) 145602 D Mata et al

Figure 3. Plot showing the dependence of the porosity, HA phaseand CNT loss of Glass/HA and CNT/Glass/HA consolidatedsamples at 1300 ◦C on the applied sintering route (P, RB + Pand HP).

composite (CNT loss of ∼20%) which is seen as a physicalobstacle to the sintering and densification process [29]. Thiscorroborates sintering studies shown by White et al underdifferent CO+H2-based atmospheres [23]. Densifications of82, 86 and 95% were shown with a respective CNT loss of54, 62, 88% [23]. These results also indicate that the presentRB+ P route is more efficient at preserving CNTs than the gasmixture flow route used by White et al [23], probably due to thehigher values of PCO that allowed a higher penetration depthof CO gas in the samples. Additionally, Glass/HA samplesrepresented an improvement of the porosity of 18%, probablydue to the local formation of carbon impurities, accordingto equation (S3) (see supporting material available at stacks.iop.org/Nano/25/145602/mmedia). The low porosity ofpressureless sintered Glass/HA samples is due to the presenceof the glass phase [4]. It forms a liquid phase that promotesthe particle rearrangement during the sintering process andaccelerates densification (classic model of Kingery [35]).Figures 4(a)–(f) show representative low magnified SEMimages of the microstructures to illustrate the porosity of thesamples. For the HP Glass/HA and CNT/Glass/HA samplesthe porosity decreased abruptly by ∼88% and ∼84% whencompared with RB+ P sintering, as shown in figures 4(c)–(f).However, the HA and CNT retention by the applied pressure

Figure 4. ((a)–(f)) SEM images of P, RB + P and HP consolidated Glass/HA and CNT/Glass/HA samples.

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was less efficient than the reactive-bed approach which resultsin a decrease of HA of ∼10% and an increase of CNT lossof 7%.

Micro-Raman spectra of figure 5(a) compare the structuralcrystallinity of post-sintered CNTs and the starting functional-ized materials. Data was processed as described in section 2.3.In the wavenumber range of 1200–1800 cm−1 all spectra showtypical signatures of MWCNT phases: one asymmetric strongpeak at∼1580 cm−1 corresponding to the G-band (E2g stretch-ing mode of the graphitic crystalline structure, square symbol)with a shoulder at higher wavenumbers ∼1618 cm−1, theD′-band (graphite in-plane disorder, circle symbol); D-band(arising from lattice disorder, triangle symbol) at lowerwavenumbers ∼1350 cm−1. Other peaks related to theGlass/HA matrix are also seen in the range of 300–1000 cm−1

wavenumbers corresponding to the internal vibration modesof the PO3−

4 tetrahedral [36].Figure 5(a) shows that all processed CNTs have higher

structural crystallinity than the starting material. This mightbe explained by the combined action of the functional groupsloosening at high temperature and the oxidation of amorphouscarbon by H2O molecules. As shown in figure S5 (seesupporting material available at stacks.iop.org/Nano/25/145602/mmedia), all functional groups are fully eliminated at1100 ◦C. By losing these deposits, it is expected that thepost-sintered CNTs will have La values closer to the rawmaterials with higher values than the functionalized ones(table S1, see supporting material available at stacks.iop.org/Nano/25/145602/mmedia). The La values could be evenhigher when the amorphous carbon is eliminated [37]. Thismight occur by oxidation at 350 ◦C (figure S3, see supportingmaterial available at stacks.iop.org/Nano/25/145602/mmedia)promoted by H2O resulting from the HA of the matrix or fromthe reactive bed (see section 3.1).

Evidence for the amorphous carbon elimination can beseen in the spectral signature by the decrease of the D-bandand the narrowing of the G-band. Comparing the spectra of Pand RB + P samples (figure 5(a)) it becomes obvious that thelast samples present less amorphous carbon. This is explainedby the differing amounts of H2O that ultimately dictate theefficiency of the oxidation reaction. The P samples generateless H2O than RB + P because in the first case it results onlyfrom the HA decomposition and in the latter also from thereactive powder bed.

Figure 5(a) shows that HP samples show a minor changein the CNT structural crystallinity, with a spectrum signaturesimilar to the raw (not shown) and functionalized CNTs,although the La value of HP CNTs is closer to the raw CNTs of18.2 ± 0.4 (table S1, see supporting material available at stacks.iop.org/Nano/25/145602/mmedia) than the functionalized,as expected. Surprisingly, despite the HA dehydroxylation(figure 3) there is no improvement of the CNT structuralcrystallinity, in contrast to the conventional sintering. This isexplained by the fast driving out of the entrapped H2O vapourby the combined applied pressure and vacuum, resulting inan incomplete oxidation of the amorphous carbon. Previousreports by others showed a similar gas expulsion effect inducedby vacuum sintering conditions [34]. Moreover, it should be

Figure 5. (a) Micro-Raman spectra of functionalized CNTs andpost-sintered CNTs by P, RB + P and HP (CNTs: • D-band, �G-band, N D′-band; Glass/HA: ? PO3−

4 bands). (b) Plot illustratingthe influence of the sintering route of (a) on the electricalconductivity and compressive strength of Glass/HA andCNT/Glass/HA samples.

highlighted that the CNT structural crystallinity (indexed bythe intensity ratio of D and G peaks (ID/IG) [25]) did notdecrease as a result of the HP processing. In this regard,the proposed HP method becomes more attractive amongother available pressure-assisted sintering approaches such asspark plasma sintering (SPS) with reports giving clues forCNT damaging under sintering conditions in the range of900–1200 ◦C and 7.5–60 MPa [38, 39].

Figure 5(b) shows the respective electrical conductivityand compressive strength of the final composites. The resultsare in accordance with the morphological characterization offigure 4. The CNT/Glass/HA RB + P sintered samples havetheir compressive strength decreased by ∼20% relative tothe P sintered ones due to the high CNT loading. In thesepartially densified composites the load bearing ability of thematrix is much decreased and CNTs may act as microstructural

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Figure 6. ((a)–(d)) SEM images of microetched and macroetched samples of hot-pressed Glass/HA and CNT/Glass/HA samples showing,respectively, the 2D and 3D intergranular disposition of CNTs.

defects, accelerating the propagation of cracks leading to thedeterioration of the mechanical properties. On the other hand,the Glass/HA samples have mechanical strength increased by∼80% due to the lower porosity and higher retention of the HAphase, the highest mechanical performance CaP phase [40].The electrical conductivity of both types of samples, Glass/HAand CNT/Glass/HA, increases when sintered in the presenceof a reactive bed, figure 5(b). The higher conductivity in theGlass/HA samples is caused by the thermal diffusion of thecarbon of the reactive bed giving an increase of two orders ofmagnitude relative to the P route. For the CNT/Glass/HA, anexceptional increase of ten orders of magnitude can be seendue to the high level of CNT retention.

For almost fully dense HP samples the compressivestrengths highly increased with respect to the RB+ P sinteredsamples, as expected. The Glass/HA and CNT/Glass/HA sam-ples exhibited an increase of ∼108% and 293%, respectively.Additionally, it is shown in figure 5(b) that the compres-sive strengths of hot-pressed Glass/HA and CNT/Glass/HAcompacts are close within a 10% window. The electricalconductivity of the Glass/HA samples increased aroundtenfold when compared to the RB + P samples due tothe higher carbon level resulting from a fast C diffusionunder pressure from the graphite die. For the hot-pressedCNT/Glass/HA samples the electrical conductivity decreasesby 56% when compared with the RB + P samples. Ofrelevance, this value is only an trend due to the high statisticalcoefficient of variation of the conductivity values for RB + Psamples of 82% compared to 5% for HP samples. The lowerreproducibility of the RB + P route might be attributed to aninhomogeneous CNT retention due to an extended dispute ofneighbour samples for the local CO+H2 atmosphere. Theconductivity decrease can be explained by the lower structuralcrystallinity of ∼34% and lower CNT loading of ∼1.7%. As

shown in section 2 (see supporting material available at stacks.iop.org/Nano/25/145602/mmedia), the electrical conductivityof CNTs is highly sensitive to the structural crystallinity. Thisindicates that porosity has an inferior influence on the electricalconductivity of the composite than the CNT crystallinity andloading; nevertheless, these two features affect the 3D electrontransporting network.

To understand how this CNT network is organized in themicrostructure a more detailed analysis was accomplished.Figures 6(a)–(d) show the microstructure of polished sur-faces of HP Glass/HA and CNT/Glass/HA samples chemicallywet-etched with two different etching strengths: microetchingand macroetching. The top view of the microetched polishedsamples clearly shows the cross-sectional profile of CNTs(highlighted with blue arrows) placed in the grain boundaries(figures 6(a) and (b)). To further reveal the 3D position ofCNTs, respective 30◦ tilted macroetched surfaces were alsoshown (figures 6(c) and (d)). The composites have predomi-nantly CNTs located at the grain boundaries corroborating the2D observations.

3.3. Optimization of HP temperature

Concerning the influence of the temperature on the CNTelimination at 1300 ◦C (see section 3.2), in the present section,the temperature was varied below the upper limit condition,1300 ◦C, to achieve the best compromise between densificationand CNT loading.

Figure 7 shows that increasing HP temperature decreasesCNT retention but, surprisingly, augments HA fraction. Thistrend is opposite to the one found in the RB + P route whereboth phases are as a general rule preserved together. Thisis evidence that CNT elimination in HP is governed by thecombined destructive action of temperature and pressure [21]

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Figure 7. Plot showing the temperature dependence of the porosity,CNT loss and phase composition of CNT/Glass/HA samplesconsolidated by HP on the temperature.

and not by simple H2O oxidation. Other evidence that H2Odoes not play a relevant role in the CNT elimination is given bythe negligible variation of the CNT crystallinity with the HPtemperature. The respective average value of La = 18.8± 0.5is close to the value of the raw CNTs before sintering of18.2 ± 0.4 (table S1, see supporting material available at stacks.iop.org/Nano/25/145602/mmedia). These observations arejustified by the key role of the applied pressure and vacuum inthe H2O elimination, previously discussed in section 3.2.

The increase of the HA retention with the HP temperatureprobably results from the decrease of the reducing effectof the vacuum atmosphere (figure S4, see supporting ma-terial available at stacks.iop.org/Nano/25/145602/mmedia).For a constant dwell time, the higher temperatures give higherdensification kinetics resulting in shorter periods of time withopen porosity. Consequently, the fast elimination of H2O byvacuum action decreases, allowing more HA retention. Therelease of some H2O from the HA structure at the closeporosity sites may permit the formation of a local H2O-richatmosphere that ends by helping the HA retention [41].Interestingly, the decomposition evolution of the oxyapatitephase, i.e. fully dehydroxylated HA, can also be seen withtemperature. Under 1200 ◦C, only the β-TCP phase is present,at 1200 ◦C, a mixture of β-TCP and α-TCP is formed andabove 1300 ◦C all β-TCP is converted to α-TCP, in accordancewith the literature [20]. A representative XRD pattern of theCNT/Glass/HA composite hot-pressed at 1100 ◦C is shown infigure S6(a) (see supporting material available at stacks.iop.org/Nano/25/145602/mmedia).

The temperature influence on the porosity of the CNT/Glass/HA composites is shown in figure 7. The porosity valuesare in agreement, even slightly better than those found in theliterature. For example, the porosity of 3.6% for materialsprocessed at 1100 ◦C is slightly lower than the 5% obtainedfor CNT/HA composites, containing similar CNT loading andhot-pressed in the same conditions, 1100 ◦C at 30 MPa [32].

Figure 8. (a) Plot depicting the influence of the HP temperature onthe grain size (an Arrhenius plot is shown as an inset).(b) Representative SEM images of (a).

At extremely high temperatures, the porosity of the com-posite varies with temperature following an atypical profile.At 1300 ◦C, the porosity increase resulting from the CNTelimination overcomes the porosity decrease by densification,giving a final increase of porosity. At temperatures lower than1300 ◦C there is an abrupt decrease of CNT loss to below∼8%. The use of temperatures around 1200 ◦C combines highCNT retention in composites with the lowest porosity.

In figure 8 the grain size dependence on the temperatureis depicted, along with the Arrhenius plot (inset). One canobserve the grain increase following an exponential law withan activation energy of ∼0.97 ± 0.004 eV (figure 8(a)). Thisvalue is much higher than those previously found for pureGlass/HA systems (0.01–0.14 eV) [4]. This indicates thatCNTs act as physical barriers to preliminary grain-to-graincontact to form columns and subsequent grain growth. Rep-resentative SEM micrographs of CNT/Glass/HA samples areshown in figure 8(b). A maximum grain size of 1.9± 0.4 µmwas achieved at 1300 ◦C (figure 8(b)) around ∼24% smallerthan the Glass/HA samples of 2.5± 0.8 µm (figure 6(a)). Yet,the regions with high CNT density such as the one shown infigure 6(b) present a much lower grain size of 0.7± 0.4 µmwhich corresponds to a 3.5-fold grain reduction. Similar graingrowth restriction of similar magnitude, tenfold, has beenreported by others [29].

Overall, the property dependence on the HP temperature(figure 9) mirrors the microstructural evolution. The maximumvalue of electrical conductivity of ∼150 S m−1 is obtained at1100 ◦C due to the maximized balance between CNT retentionand densification (figure 7). Comparable conductivities ofaround 200 S m−1 were achieved at similar CNT loadingsfor >95% densified samples using pressure-assisted meth-ods [42, 43]. Here, it is proved that CNT loading has ahigher influence in the electrical conductivity than porositydue to the long-range CNT connecting through intergranularlocations. HP samples at 1000 ◦C presenting higher porosityhave larger electrical conductivity than those processed at1300 ◦C. A similar behaviour was presented by the high porous

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Figure 9. Plot illustrating the effect of the HP temperature on theelectrical conductivity and compressive strength of CNT/Glass/HAsamples.

RB+P samples howsoever having high conductivity (figures 3and 5(b)).

Comparing the average conductivity of these partiallydensified samples (figure 5(b)) with the HP sample at 1100 ◦C(figure 9) it is evident that the structural crystallinity of theCNTs also has an important contribution to the final electricalconductivity. HP samples have ∼25% more CNTs than theRB+ P samples but the respective structure crystallinity of theCNTs is ∼57% lower which resulted in an effective electricalconductivity reduction of ∼60%.

The highest mechanical strength samples were obtainedat 1100 ◦C (figure 9), corroborating the results reported byothers [39]. This performance resulted from the best combina-tion of porosity, CNT loss, grain size and HA retention. Thedetailed mechanical characterization of these composites wasreported previously [44]. Also, by comparing the compressivestrength at 1100 and 1300 ◦C of composites with similarporosity it becomes clear that the fraction of the secondaryphase α-TCP also has a relevant contribution (figure 7). At1300 ◦C, the amount of α-TCP phase becomes significant tostrengthen the HA, by putting the HA lattice under stress dueto the higher occupied volume of α-TCP compared to HA [4],and not due to the intrinsic properties of the α-TCP itself [40].

Envisioning the preparation of biomaterials for the elec-trotherapy of bone and bearing in mind the superior mechanicaland electrical performance of the CNT/Glass/HA compositeconsolidated by hot-pressing at 1100 ◦C, only this material,simply referenced as CNT/Glass/HA, was further evaluatedfor its osteoblastic cytocompatibility.

3.4. Osteoblastic cytocompatibility of CNT/Glass/HAcomposites

The osteoblastic cytocompatibility of the CNT/Glass/HA com-posite was assessed with MG63 osteoblastic cells, cultured for

5 days (figures 10 and 11). Cell response was compared to thaton the Glass/HA composite and the standard cell culture coverslip (control).

DNA content, which reflects the number of cells, in-creased throughout the culture time over the three testedsurfaces, figure 10(a). At day 1, similar values were found,suggesting an identical cell attachment on the coverslip and thecomposites. However, afterwards, cell proliferation was loweron the CNT/Glass/HA composite, shown by the DNA contentat days 3 and 5. MTT assay was also performed in the culturedsubstrates, figure 10(b). This method is based on the reduc-tion ability of mitochondrial dehydrogenases in viable cellsand, thus, both cell metabolic activity and cell proliferationcontribute for the measured MTT reduction values. Resultsshowed also decreased values on the CNT/Glass/HA compos-ite. The percentage decrease was similar for the DNA contentand the MTT reduction values, suggesting that although cellproliferation was lower on the CNT/Glass/HA composite, cellsmaintained a normal metabolic activity. Also, the viability ofthe attached cells, assessed as the membrane integrity, wasclose for both composite surfaces as evidenced by the similarLDH values at day 5 (figure 10(c)).

Results show that CNT composites possessed no acutetoxicity in a human osteoblastic cell line but there are evidentdifferences in the cell proliferation response between bothcomposite surfaces. The inclusion of CNTs in the Glass/HAmatrix resulted in distinct surface properties, namely for thebioactivity level, known to modulate cell behaviour. The sin-tered Glass/HA composites have a Ca/P ratio of 1.68 (figureS6(b), see supporting material available at stacks.iop.org/Nano/25/145602/mmedia), close to 1.67 of highly bioactive stoi-chiometric HA [45]. Besides HA, the bioglass in the Glass/HAcomposites is even more bioactive than HA alone, as reportedby others [5], and is expected to be much more bioactivethan CNTs. As such, the bioactivity level and the surface areapercentage of bioactive phases on the Glass/HA compositeis higher than on the CNT/Glass/HA which may explain thehigher cell proliferation on the former at days 3 and 5. Yet,this does not explain the proliferation differences betweenthe CNT/Glass/HA and the control. Following this, the cellALP activity was further investigated. The ALP enzyme is animportant osteoblastic differentiation marker, having a key rolein the mineralization of the bone collagenous matrix [46]. Atday 5, an increased ALP activity (figure 10(d)) was observedover the CNT composite. Due to the established reciprocalrelationship between proliferation and differentiation duringthe development of the osteoblastic phenotype [46], the lowercell proliferation but higher ALP activity found over theCNT/Glass/HA composite suggest an inductive effect on theosteoblastic differentiation, compared to that observed overthe Glass/HA composite.

The performance of the CNT/Glass/HA composite wasalso assessed by SEM and CLSM observation. Representativeimages are shown in figure 11 for cultures with 1 and 5 days. OnSEM, cells showed an elongated morphology and the cell layerdisplayed an evident oriented cell growth, mostly following thealignment direction of the CNT agglomerates on the Glass/HAmatrix. By contrast, on the cover slip and on the Glass/HA

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Figure 10. Behaviour of osteoblastic cells cultured over standard cell culture plates (control), Glass/HA and CNT/Glass/HA composites.(a) DNA content, (b) cell viability/proliferation, (c) LDH leakage and (d) ALP activity. Results are expressed as a percentage of variationfrom control cultures; significantly different from the control (*) and from the Glass/HA composite (**).

Figure 11. Representative SEM images of the cell growth pattern over standard cell culture plates, Glass/HA and CNT/Glass/HA compositeat days 1 and 5, and CLSM images at day 5, showing the F-actin cytoskeleton (green) and nucleus (red).

composite, cells showed a polygonal/elongated morphologyand a random cell growth pattern. On CLSM observationcells exhibited a well-organized F-actin cytoskeleton and aprominent nucleus on the three surfaces. The intense F-actinstaining at the cell boundaries and the evidence of on-going cell

division suggest a healthy cell behaviour on the three substrates(figure 11). Also, compared to the SEM images, similarinformation was provided regarding the organization of thecell layer, i.e. an aligned cell growth over the CNT/Glass/HAcomposite compared to the other two surfaces.

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As observed in the present work, previous studies alsoshowed that the spatial organization of the CNTs modulatesthe directional growth and the behaviour of osteoblastic cells.Thus, increased cell adhesion and proliferation of osteoblastswas reported over seeded aligned conductive nanofibers ofbiodegradable poly-DL-lactide (PLA) embedded with MWC-NTs [18], and also enhanced proliferation, osteoblastic geneexpression and ALP activity were described for hMSC onaligned CNT structures [47], compared to randomly orientedCNT networks. This behaviour may be mediated by the highactin cytoskeletal tension accumulated on aligned cells, whichthus leads to the activation of selective mechano-transductionpathways, via the activation of Rho-family GTPase signalling,known to induce osteogenic differentiation [48].

The behaviour observed on the CNT/Glass/HA compositeis in line with the available studies regarding the perfor-mance of CNTs containing materials. Regarding ceramic–matrix composites, the inclusion of CNTs in plasma sprayedhydroxyapatite coatings was found to significantly improvethe physical properties of the coating, maintaining its adequatebiological interaction with human osteoblasts [49]. Also, thedevelopment of hydroxyapatite–MWCNT composite coatingssignificantly improved the biological performance of the con-struct, by increasing cell proliferation and alkaline phosphataseactivity of MC3T3-E1 pre-osteoblastic cells, compared to thebare substrate (titanium alloy) and pure HA coating [50]. Inbioactive glass-containing materials, inclusion of CNTs alsoresulted in positive outcomes. Thus, the addition of MWCNTs(2 wt%) to poly(3-hydroxybutyrate) composites with bioactiveglass particles enhanced the adhesion and proliferation ofMG-63 osteoblastic cells [51], and on 45S5 Bioglass R©-derivedglass–ceramic scaffolds coated with CNTs, hMSCs showedlower cell adhesion and viability but increased ALP, as com-pared to uncoated scaffolds [52].

Overall, the inclusion of CNTs in ceramic–matrix com-posites seems to prompt and improve the proliferation and/orthe activity of osteoblast lineage cells and the whole biologicalresponse related to the bone tissue metabolism and regenera-tion.

4. Conclusions

High electroconductive CNT/Glass/HA bone grafts with con-ductivities in the range of 140–640 S m−1 were successfullyproduced with just 2.5 wt% of CNTs by reactive-bed pressure-less sintering (RB + P) and hot-pressing (HP). Both sinteringroutes were optimized towards CNT preservation through twomechanisms: by preventing the HA dehydroxylation in RB+Psamples to ultimately avoid CNT oxidation. Following this,reactive powder beds of HA + C were used to form CO+H2atmospheres at positive pressure to shift the equilibrium of thereaction of H2O+C, according to the Chatelier’s principle.Phase retentions of 100% and 80% were achieved for HAand CNT by using an atmosphere of PCO = 1.77 MPa bysimply adjusting the HP temperature. For HP samples, theentrapped H2O from the HA dehydroxylation did not oxidizethe CNTs due to the fast dragging out by the combined applied

pressure and vacuum. The optimal balance of CNT loss anddensification of ∼5% and ∼96% was found at 1100 ◦C.

The CNT/Glass/HA composites allowed cellular adhesionand proliferation with favoured improvement of the functionalactivity. These materials offered further modulation of the cellgrowth orientation according to the alignment direction of theCNT agglomerates.

Non-metallic electroconductive bone grafts offer new ex-citing possibilities in bone regeneration strategies by deliveringin situ electrical stimulus to cells and controlling the boneformation rate. Hopefully, conductive smart materials mightturn the in situ bone electrotherapy of clinical relevance bydecreasing the postoperative healing times.

Acknowledgments

The work was carried out with the financial support from theprojects Biomaterials for Regenerative Medicine (CENTRO-07-ST24-FEDER-002030), co-financed by QREN, programmeMais Centro-Programa Operacional Regional do Centro andUniao Europeia/ Fundo Europeu de Desenvolvimento Re-gional, and PEst-C/CTM/LA0011/2013. Rui F Araujo isgrateful to the Foundation for Science and Technology (FCT,Portugal) for funding the project SFRH/BPD/88920/2012.

References

[1] Jarcho M, Bolen C H, Thomas M B, Bobick J, Kay J F andDoremus R H 1976 J. Mater. Sci. 11 2027–35

[2] Akao M, Aoki H and Kato K 1981 J. Mater. Sci. 16 809–12[3] LeGeros R Z and LeGeros J P 1993 An Introduction to

Bioceramics ed L L Hench and J Wilson (Singapore: WorldScientific) p 145

[4] Lopes M A, Monteiro J F and Santos J D 1999J. Biomed. Mater. Res. 48 734–40

[5] Lopes M A, Knowles J C, Santos J D, Monteiro F J andOlsen I 2000 Biomaterials 21 1165–72

[6] Lahiri D, Ghosh S and Agarwal A 2012 Mater. Sci. Eng. C32 1727–58

[7] Usui Y et al 2008 Small 4 240–6[8] Akasaka T, Watari F, Sato Y and Tohji K 2006 Mater. Sci.

Eng. C 26 675–8[9] Harris P J F 2009 Carbon Nanotube Science—Synthesis,

Properties and Applications 2nd edn (Cambridge:Cambridge University Press)

[10] Griffin M and Bayat A 2011 Eplasty 11 e34[11] Goldstein C, Sprague S and Petrisor B 2010 J. Orthop.

Trauma 24 S62–5[12] Griffin M, Iqbal S, Sebastian A, Colthurst J and Bayat A 2012

J. Bone Joint Surg. Br. 94 33[13] Woo D G, Shim M-S, Park J S, Yang H N, Lee D-R and

Park K-H 2009 Biomaterials 30 5631–8[14] Schwartz Z, Simon B, Duran M, Barabino G, Chaudhri R and

Boyan B 2008 J. Orthop. Res. 26 1250–5[15] Akai M and Hayashi K 2002 Bioelectromagnetics 23 132–43[16] Navarro M, Michiardi A, Castano O and Planell J A 2008 J. R.

Soc. Interface 5 1137–58[17] Supronowicz P R, Ajayan P M, Ullmann K R,

Arulanandam B P, Metzger D W and Bizios R 2002J. Biomed. Mater. Res. 59 499–506

13

Nanotechnology 25 (2014) 145602 D Mata et al

[18] Shao S, Zhou S, Li L, Li J, Luo C, Wang J, Li X and Weng J2011 Biomaterials 32 2821–33

[19] Gittings J P, Bowen C R, Dent A C E, Turner I G, Baxter F Rand Chaudhuri J B 2009 Acta Biomater. 5 743–54

[20] Wang T, Dorner-Reisel A and Muller E 2004 J. Eur. Ceram.Soc. 24 693–8

[21] Peigney A, Rul S, Lefevre-Schlick F and Laurent C 2007J. Eur. Ceram. Soc. 27 2183–93

[22] Rapacz-Kmita A, Paluszkiewicz C, Slosarczyka A andPaszkiewicz Z 2005 J. Mol. Struct. 744 653–6

[23] White A A, Kinloch I A, Windle A H and Best S M 2010 J. R.Soc. Interface 7 S529–39

[24] Proenca M F, Araujo R F, Paiva M C and Silva C J R 2009J. Nanosci. Nanotechnol. 9 6234–8

[25] Cancado L G, Takai K, Enoki T, Endo M, Kim Y A,Mizusaki H, Jorio A, Coelho L N, Magalhaes-Paniago Rand Pimenta M A 2006 Appl. Phys. Lett. 88 163106

[26] Gazzara C P and Messier D R 1977 J. Am. Ceram. Soc. Bull.56 777–80

[27] Kalita S J, Bhardwaj A and Bhatt H A 2007 Mater. Sci. Eng. C27 441–9

[28] Hench L L 1998 Biomaterials 19 1419–23[29] Inam F, Yan H, Peijs T and Reece M J 2010 Compos. Sci.

Technol. 70 947–52[30] Jacob K D, Reynolds D S and Hill W L 1962 Indust. Eng.

Chem. 21 1126–32[31] Li H, Zhao N, Liu Y, Liang C, Shi C, Dua X and Li J 2008

Composites A 39 1128–32[32] Lei T, Wang L, Ouyang C, Li N-F and Zhou L-S 2011 Int. J.

Appl. Ceram. Technol. 8 532–9[33] Meng Y H, Tang C Y, Tsui C P and Chen D Z 2008 J. Mater.

Sci.: Mater. Med. 19 75–81[34] Li A, Sun K, Dong W and Zhao D 2007 Mater. Lett.

61 1839–44[35] Kingery W D 1959 J. Appl. Phys. 30 301–6[36] Cusco R, Guitian F, Aza S and Artus L 1998 J. Eur. Ceram.

Soc. 18 1301–5

[37] Mata D, Silva R M, Fernandes A J S, Oliveira F J,Costa P M F J and Silva R F 2012 Carbon 50 3585–606

[38] Lahiri D, Singh V, Keshri A K, Seal S and Agarwal A 2010Carbon 48 3103–20

[39] Xu J L, Khor K A, Sui J J and Chen W N 2009 Mater. Sci.Eng. 29 44–9

[40] Liang L, Rulis P and Ching W Y 2010 Acta Biomater.6 3763–71

[41] Wang T, Dorner-Reisel A and Muller E 2004 J. Eur. Ceram.Soc. 24 693–8

[42] Inam F, Yan H, Reece M J and Peijs T 2008 Nanotechnology19 195710

[43] Tatami J, Katashima T, Komeya K, Meguro T and Wakihara T2005 J. Am. Ceram. Soc. 88 2889–93

[44] Mata D, Oliveira F J, Ferro M, Gomes P S, Fernandes M H,Lopes M A and Silva R F 2013 J. Biomed. Nanotechnol.9 1–19

[45] Hong L, Xu H C and Groot K D 1992 J. Biomed. Mater. Res.26 7–18

[46] Aubin J 2008 Principles of Bone Biology ed J P Bilezikian,L G Raisz and T J Martin (New York: Academic)

[47] Namgung S, Baik K Y, Park J and Hong S 2011 ACS Nano5 7383–90

[48] Arnsdorf E J, Tummala P, Kwon R Y and Jacobs C R 2009J. Cell Sci. 122 546–53

[49] Balani K, Anderson R, Laha T, Andara M, Tercero J,Crumpler E and Agarwal A 2007 Biomaterials 28 618–24

[50] Hahn B-D, Lee J-M, Park D-S, Choi J-J, Ryu J, Yoon W-E,Lee B-K, Shin D-S and Kim H-E 2009 Acta Biomater.5 3205–14

[51] Misra S K, Ohashi F, Valappil S P, Knowles J C, Roy I,Silva S R P, Salih V and Boccaccini A R 2010 ActaBiomater. 6 735–42

[52] Meng D, Rath S N, Mordan N, Salih V, Kneser U andBoccaccini A R 2011 J. Biomed. Mater. Res. A 99 435–44

14