Materials Science and Engineering C 29 (2009) 2124–2132

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Materials Science and Engineering C

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Rational synthesis of a nanocrystalline calcium phosphate cement exhibiting rapidconversion to hydroxyapatite

Inés S. Neira a,b,c,⁎, Yury V. Kolen'ko b,1, Oleg I. Lebedev d, Gustaaf Van Tendeloo d, Himadri S. Gupta c,Nobuhiro Matsushita b, Masahiro Yoshimura b, Francisco Guitián a

a Galician Institute of Ceramics, University of Santiago de Compostela, E-15782 Santiago de Compostela, Spainb Materials and Structures Laboratory, Tokyo Institute of Technology, 4259 Nagatsuta, Midori-ku, 226-8503 Yokohama, Japanc Department of Biomaterials, Max Planck Institute of Colloids and Interfaces, D-14424 Potsdam, Germanyd Electron Microscopy for Materials Science, University of Antwerp, B-2020 Antwerp, Belgium

⁎ Corresponding author.Mailing address: Instituto de Cede Santiago de Compostela, Avda. Mestre Mateo, s/n, E-1Spain. Tel.: +34 981 563100x16885; fax: +34 981 56424

E-mail address: ines.sanchez@usc.es (I.S. Neira).1 Present address: Department of Inorganic Chemist

Max Planck Society, D-14195 Berlin, Germany.

0928-4931/$ – see front matter © 2009 Elsevier B.V. Adoi:10.1016/j.msec.2009.04.011

a b s t r a c t

a r t i c l e i n f o

Article history:Received 2 February 2009Received in revised form 21 March 2009Accepted 15 April 2009Available online 19 April 2009

Keywords:Calcium phosphate cementRapid conversionNanostructureMechanical propertiesNanoindentationOsteoblast cell culture

The rational synthesis, comprehensive characterization, and mechanical and micromechanical properties of acalcium phosphate cement are presented. Hydroxyapatite cement biomaterial was synthesized from reactivesub-micrometer-sized dicalcium phosphate dihydrate and tetracalcium phosphate via a dissolution-precipitation reaction using water as the liquid phase. As a result nanostructured, Ca-deficient andcarbonated B-type hydroxyapatite is formed. The cement shows good processibility, sets in 22±2 min andentirely transforms to the end product after 6 h of setting reaction, one of the highest conversion rates amongpreviously reported for calcium phosphate cements based on dicalcium and tetracalcium phosphates. Thecombination of all elucidated physical-chemical traits leads to an essential bioactivity and biocompatibility ofthe cement, as revealed by in vitro acellular simulated body fluid and cell culture studies.The compressive strength of the produced cement biomaterial was established to be 25±3 MPa.Furthermore, nanoindentation tests were performed directly on the cement to probe its local elasticityand plasticity at sub-micrometer/micrometer level. The measured elastic modulus and hardness wereestablished to be Es=23±3.5 and H=0.7±0.2 GPa, respectively. These values are in close agreement withthose reported in literature for trabecular and cortical bones, reflecting good elastic and plastic coherencebetween synthesized cement biomaterial and human bones.

© 2009 Elsevier B.V. All rights reserved.

1. Introduction

Calcium phosphate cements (CPCs) belong to a very importantclass of biomaterials because of their potential to be utilized in thehuman body for bone repair and substitution [1,2]. Essentially, theseapplications stem from their associated biological and physiologicalproperties — CPCs are highly bioactive [3], biocompatible [4] andosteoconductive [5], i.e., cements are able to be resorbed by biologicalserum, allowing its progressive substitution by newly formed bone.Nevertheless, owing to the lower fracture toughness parameters ofCPCs in comparison to human bones [6], the clinical applications ofthe cements are limited to areas where bones are free of dynamic load,i.e., for non—load bearing as well as craniofacial and periodontal

rámica deGalicia, Universidade5782, Santiago de Compostela,2.

ry, Fritz Haber Institute of the

ll rights reserved.

applications [7], or as materials for the development of scaffolds inbone tissue engineering [8].

From a fundamental and an application point of view, it isimportant to design a rational synthesis route to a cement withsetting time not more than thirty minutes (surgeons' requirements)and a rapid conversion rate to the end product, preferably thermo-dynamically stable hydroxyapatite (HA) — the inorganic componentof mineralized bone tissues. It is noteworthy that in many cases therapid conversion of the CPC precursors to HA not only leads to the riseof the appropriate physiological properties, but also results in thestrengthening of the mechanical properties. In particular, rapidconversion accelerates achievement of the final compressive strengthvalue, which is almost linearly dependent on the extent of the CPCsetting reaction [9]. However, the preparation of CPCs with theaforementioned characteristics still remains challenging.

Bones are hierarchically organized biological systems originated atthe nanometre scale [10]. Therefore, it is beneficial to investigate themechanical behaviour of the cements at the nano- and micro-structural levels with regards to its application for bone graft uses.Such information is important for the exploring of mechanisticcompatibility between functional cement biomaterials and bones.

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Recent advances in this area have demonstrated that nanoindentationprobing is one of the most promising techniques which allows toestimate such important mechanical properties of the materials ashardness and elastic modulus at unprecedented scale [11]. An elasticmodulus similar to that of the human bones is an essential propertyfor bone grafts, providing a good load transfer from the implant intothe bone, which is required for the stimulation of new bone formation[12]. In contrast, if biomaterial is much stiffer than bone, it shields thenearby bone from mechanical stress, generating a sort of disuseatrophy — the bone resorbs.

Herein, we present the rational synthesis of a cement togetherwith the detailed characterization of the structure at the micrometerand at the nanometer level. Elastic modulus and hardness of derivedbiomaterial were also investigated and compared to those of bonetissue and composite biomaterials.

2. Materials and methods

2.1. Starting materials

Ca(H2PO4)2·H2O (90.0% Kanto), CaCO3 (98.0% Wako), CaHPO4

(98.0% Wako), isopropanol (99.5% anhydrous, Sigma-Aldrich), acetone(SC grade, Wako), methylmethacrylate (99% Merck), benzoylperoxid(for synthesis, 25% H2O, Merck) and nonylphenylpolyethylenglycolacetate (for synthesis, Fluka) were used as received. The preparation ofacellular simulated bodyfluid (SBF) solution having pH 7.4was adoptedfrom that of Kokubo et al. [13].

2.2. Preparation of the CPC precursors

Dicalcium phosphate dihydrate (DCPD) was synthesized via aprecipitation reaction between Ca(H2PO4)2·H2O and CaCO3 [14],while tetracalcium phosphate (TTCP) was prepared from stoichio-metric amounts CaCO3 and CaHPO4 by solid-state reaction [15]. It isreasonable to expect that the reactivity of the DCPD and TTCP willincrease with decreasing of their crystal sizes that, accordingly, willfacilitate the formation of respective cement. On the basis of thisconcept, the as-synthesized raw materials were subjected to a wetrolling ball-milling to reduce the size of the particles. First, one third ofa 5% silica-doped alumina pot (V=300 mL) was filled with siliconnitride balls of 3.6 mm Ø in case of DCPD grinding or with high purityzirconia balls of 4.8 mm Ø in case of TTCP milling. Afterwards, 15 g ofthe powder was placed inside the pot, along with 60 mL of distilledwater (for DCPDmilling) or anhydrous isopropanol (for TTCPmilling).The ball-milling was performed at 135 rpm for 24 h. The product wascollected by vacuum filtration and washed with acetone. DCPD wasdried at the ambient conditions, while TTCP was dried at 80 °C on air.

2.3. Preparation of the calcium phosphate cement

CPC was generated by the dissolution-precipitation reactionbetween DCPD and TTCP [16]. Equimolar amounts of ball-milledpowders were thoroughly mixed and placed inside a latex finger cot,along with an appropriate amount of distilled water (solid to liquid(S/L) ratio: 2 g/mL), and then hand kneaded for 1 min. The resultantputty-like material was rod-shaped, allowed to set for 1 h at ambientconditions and incubated in distilled water at human body tempera-ture of 37 °C for the desired period of time. The CPC specimens forcompression strength and nanoindentation tests were prepared, inthe form of pellets, using S/L ratio of 3 g/mL. The as-derived putty-likematerial was placed into a cylindrical mold assembly of 6.9 mm Ø. Thesample was pressed for 1 min under 5.3 MPa (lowest value of pressurescale of used loading device). The obtained pellets were kept for 1 h atambient conditions to set and then incubated in a SBF solution at 37 °Cfor 23 h.

2.4. Characterization methods

The products were characterized by powder X-ray diffraction(XRD) using a Rigaku RINT 2000 diffractometer. The phases wereidentified by comparisonwith the data reported in the ICDD database.The morphology was studied by scanning electron microscopy (SEM)and field-emission environmental SEM (FE-ESEM) using Hitachi S-4500 and FEI Quanta 600F microscopes operating at 15 and 5 kV,respectively. Average particle sizes and particle size distributions(PSD) of the ball-milled powders were determined using dynamiclight-scattering (DLS) technique on a Malvern Instruments ZetasizerNano-ZS. For DLS measurements, a small amount of the powder wasfully dispersed in water using high-power ultrasonication, and theresultant stable suspension was subjected to the DLS analysis. Finalsetting time (ST) of the cement was determined according to theAmerican Dental Association Specification No. 9 using the Vicat needlemethod. ST value is the average of at least six replicates. Energy-dispersive X-ray spectroscopy (EDX) for semiquantitative Ca and Pcontent determination was performed with a EDAX DX-95 spectro-metry system. The room temperature Raman scattering measure-ments were carried out on a confocal Raman microscope WITECCRM200 with linear polarized laser light (λ=532 nm) as theexcitation source. The room temperature diffuse reflectance infrared(IR) Fourier-transform spectra were recorded on a Jeol JIR-7000spectrometer. The fine microstructure and phase purity of the cementwere investigated by transmission electron microscopy (TEM),electron diffraction (ED) and high-resolution TEM (HRTEM) usingJeol 4000EX microscope operating at 400 kV. The samples for TEMwere crushed, dispersed in methanol and deposited on a holey carbongrid. Computer-simulated HRTEM images were obtained using theMac Tempas and Crystal Kit software programs.

2.5. In vitro testing

In vitro bioactivity investigationwas carried out using SBF solution.The CPC test pellets were placed into a polystyrene bottle filled withSBF and incubated at 37 °C. The polystyrene bottle and SBF solutionwere refreshed every 3 days.

Biocompatibilitywas studied invitro through cell culture experimentadopted from that ofManjubala et al. [17] usingmurine pre-osteoblasticcells (MC3T3-E1) from mouse calvarie. 5×105 cells were suspended in100 µL of culture medium and seeded on the CPC test specimens in theform of pellets. The specimens were incubated at 37 °C for 30 min inPetri dishes to allow the cells to attach to the surface and then 2 ml ofmediumwere added. Themediumand Petri dishwere refreshed twice aweek. Alkaline phosphatase (ALP) enzyme activity was quantitativelydetermined after 1, 2 and 3weeks. The cell proliferation and distributionover the CPC surfacewere analyzed using a light microscope employingGiemsa staining histological method.

2.6. Mechanical properties measurements

Compressive strength (CS) measurements were performed using aShimadzu Autograph AG-I universal testing machine. Prior to CSexperiment, the CPC test pellets with an aspect ratio of 2 (diameter:6.9 mm; length: 13.8 mm) were allowed to set for 1 h at ambientconditions and then immersed in distilled water for 23 h at 37 °C. CSvalue is the average of at least six replicates.

Elastic modulus and hardness of the cement were measured usingnanoindentation technique (NI). For this purpose, the CPC pellet wasepoxy embedded in a solution formed by 77% of methylmethacrylateand the polymerization initiators — 1.4% of benzoylperoxid and 21.6%of nonylphenylpolyethylenglycol acetate. To fill up the pores, epoxyembedding was achieved by gradual polymerization of methylmetha-crylate to poly(methylmethacrylate) applying a low temperatureprofile [(39 °C,12 h)->(48 °C,12 h)->(55 °C, 12 h)]. Then, the section

Fig. 1. Powder XRD patterns are shown for: (A) — DCPD product prepared byprecipitation reaction (1) following by ball-milling (2) and (B) — TTCP productprepared by solid-state reaction (1) following by ball-milling (2). Tick marks below thepatterns correspond to the positions of the Bragg reflections expected for themonoclinic DCPD and TTCP (ICDD no. 9-77 and no. 70-1379, respectively).

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for NI was cut in parallel to the pellet base, polished by a set of abrasivepapers of decreasing size of grit followed by polishing with diamondpaste down to a particle size of 1 µm. Quasi-static nanoindentationtests were performed on a Hysitron TriboScan UBI-1 nanohardnesstester in conjunction with a Digital Instruments Nanoscope III atomicforce microscope (AFM). A Berkovich diamond indenter tip withregular triangle pyramid geometry and nominal radius of 300 nmwasused. A loading function with a peak load of 5 mN, loading rate of1 mN/s, holding time 60 s, unloading to 1 mN with rate of 0.4 mN/s,holding at 1 mN for 20 s and unloading back to 0 mN with rate of0.5 mN/s, was used.

Evaluation of the load-displacement data obtained by NI wasperformed using the software program implemented in the TriboScan,which applies Oliver and Pharr method [18] to calculate the reducedindentation modulus Er and hardness H. During unloading, the drifteffect was observed and corrected by the same software program. Er isrelated to elastic modulus Es through the Eq. (1):

1Er

=1− m2i

� �

Ei+

1− m2s� �

Esð1Þ

wherein νi=0.07 is the Poisson's ratio for the diamond tip [19],νs=0.28 has been chosen as the Poisson's ratio for synthetic HA [20]and Ei=1141 GPa is the diamond elastic modulus [19]. Two hundredfifty indents were randomly performed to obtain average values for Erand H.

3. Results

According to the XRD analysis, phase-pure DCPD and TTCPcompounds were produced by precipitation and high temperaturereactions, respectively (Fig. 1A and B). In order to improve the CPCprecursors contact and to enhance their reactivity, the primary crystalsizes were reduced by mechanochemical grinding. XRD reveals thatthe ball-milling does not lead to changes in the phase composition andthe as-prepared samples were found to be phase-pure DCPD and TTCP(Fig. 1A and B). It should be noted that for the ball-milled TTCPproduct, a broadening of the XRD peaks is observed (Fig. 1B). This islikely due to the comminution of agglomerates, reduction of crystal-lites sizes and crystal structure deformation from the ball-millingprocedure, overall resulting in partial amorphization of the ball-milledmaterial [21], in good agreement with reported data [22]. SEM andDLS analyses of the ball-milled products clearly reveal almost non-aggregated DCPD and TTCP particles with average particle sizesapproximately 1.1 and 1.5 µm, respectively, exhibiting relativelynarrow PSD (Fig. 2). The CPC was generated from these reactivesub-micrometer-sized precursors via a dissolution-precipitation reac-tion using water as the liquid phase. The setting time study of thecement indicates that as-prepared CPC sets within 22±2 min. It isnoteworthy that this quick ST was achieved by using exclusively wateras the liquid mixing phase, i.e., without any of the setting acceleratorswidely used in cement formulations [23]. The cement phasecomposition evolution during the setting reaction at the selectedtimes (t) of 0–8 and 24 h is shown in Fig. 3. According to XRD, the CPCat t=0 consists of, as expected, pure DCPD and TTCP (ICDD No. 9–77and 70–1379, respectively). Further phase analysis reveals totalconsumption of DCPD already at t=4 h, while TTCP is fully exhaustedat t=6 h. At this time of reaction (t=6 h) and higher, cement productconsists of only phase-pure HA (ICDD No. 72-1243). The powder XRDpatterns of the 6–8 and 24 hour-derived HA cements exhibit lowintensities and very broad peaks (Fig. 3). This observation wouldsuggest that these samples are essentially amorphous; however, asestablished by TEM investigation, these CPCs are crystalline. Hence,aforementioned features of the XRD patterns aremost likely caused by

the nanocrystalline nature and defective HA structure of the cement(vide infra).

The average molar ratio of Ca to P elements (Ca/P) in stoichio-metric HA is 1.67, while Ca/P ratio in the 24 hour-derived CPC wasestablished to be ~1.49 (EDX), indicating that as-synthesized cementis Ca-deficient HA. The observed non-stoichiometry is mainly a resultof partial substitution of the PO4

3- by HPO42- in the HA crystal structure,

as established by Raman scattering (not shown). Fig. 4 displays arepresentative IR spectrum for the 24 hour-derived cement. The set ofbands and spectra features agrees fairly well with the reported IR datafor HA [24]. The broad peak at ~3406 cm-1 reflects the physisorbedwater, peaks observed at ~2358 and 2339 cm-1 correspond to theatmospheric CO2(g), while the appearance of the H-O-H deformationband at ~1643 cm-1 suggests the existence of free water moleculestrapped in the crystal lattice of the HA cement [25]. Additionally, IRspectroscopy confirms the incorporation of CO3

2- anions into thestructure of HA (Fig. 4), indicating that the obtained HA cement is notonly Ca-deficient, but also B-type carbonated HA, thus, with the CO3

2-

partially substituting PO43− groups [26]. These carbonate anions are

believe to come from the synthesis procedures, namely, from the

Fig. 2. SEM images (left panels) along with the PSD (right panels) from ball-milled DCPD (A) and TTCP (B) powders.

Fig. 3. Comparison of the powder XRD patterns at different times of the cement setting.The dotted lines correspond to themost intense diffraction peaks of DCPD, TTCP and HA.

Fig. 4. IR spectrum collected from the 24 hour-derived cement product. Thecharacteristic bands representing apatitic phosphate PO4

3- and hydroxyl OH- groups;lattice (deformation H-O-H band) and physisorbedwater; as well as atmospheric CO2(g)

are marked by grey regions. Dotted lines are drawn at the positions of CO32- anion

vibrations to indicate B-type (carbonate replacing phosphate) of carbonated HA.

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Fig. 5. SEMmicrographs of the cement products after 1 h (A) and 24 h (B) of setting. A low-magnification TEM image (C) and corresponding ED pattern (given as insert in C) from the24 hour-derived cement product as well as [012] HRTEM image of single cement nanocrystal (D). The FT pattern is shown as a top insert in (D). A simulated (112) HRTEM image(t=5 nm, Δf=−30 nm) is given as a main panel inset in (D). The intense black dots represent the columns of the calcium atoms and the intense bright dots correspond to thechannels.

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atmospheric/water CO2 and/or from tiny impurity of the CaCO3 phase(if any) utilized for raw DCPD and TTCP synthesises. Remarkably,biological apatite contains 3–5% carbonate groups, and therefore forbiomedical application a carbonated HA cement is rather a require-ment than a problem.

A SEM image of cement product after 1 h of setting is shown inFig. 5A. It displays a highly aggregated microstructure with visiblealteration of the facets of the ball-milled precursors, i.e., owing to thesmoothing of edges, the grains become less faceted in comparison tothe primary morphologies (Fig. 2, left panels). Since the CPCprecursors were used as faceted non-agglomerated particles, theabove results show that the partial dissolution of the CPC reagentsalready occurs at this stage of the reaction (t=1). The SEM image of24 hour-derived cement (Fig. 5B) reveals highly associated micro-structure of the product with clearly visible macropores (ØN50 nm).SEM observations also reveal the lack of the particulate structureregions in this CPC, and that the surface of the cement is free of anycrystal faceting. These effects show that the dissolution-precipitationreaction between ball-milled DCPD and TTCP is complete.

Low-magnification and high-resolution TEM images of 24 hour-derived CPC are shown in Fig. 5C and D, respectively. The sampleconsists of interlocking nanoscaled platelet crystals. The compactpackaging of the nanocrystallites is subsequently confirmed to be a

common feature in the cement. Namely such spontaneous interlock-ing of the precipitated nanocrystals results in the structural solidity ofthe cement. The crystallites size lies within the range of 10–15 nm andtheir surfaces are free of any amorphous or secondary phase (Fig. 5C).These results indicate that the line broadening effect and low intensityof the cement XRD patterns (Fig. 3) are likely related to thenanocrystalline nature of the CPC end product and not associatedwith the amorphous character of the cement.

Electron diffraction of the CPC shows a very dense ring pattern,exhibiting a lack of distinct diffraction spots, typical for a nanocrys-tallite clustering (Fig. 5C, inset). The ED pattern closely resembles theHA phase of ICDD No. 21-1272 (hexagonal, P63/m, a=9.432 Å,c=6.881 Å), and devoid any impurity phases. The rings of the EDpattern can be indexed using the characteristic d-spacings of thehexagonal HA phase (d211, d102, and d002) determined from powderXRD analysis of this cement product. The very weak intensity of thed100 and d101 rings representing large d spacing in the ED pattern(Fig. 5C, inset) is related to the platelet shape of the nanocrystalsresulting in a preferential orientation of the nanocrystallites. Asdetermined by HRTEM, the nanoparticles in the as-synthesizedsample preferentially growth along the c-axis, but there arenanocrystallites grown along uncommon 012 direction (Fig. 5D).The corresponding Fourier transform (FT) of the nanocrystallite is

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shown as the top inset in Fig. 5D. This pattern can be indexedaccording to the HA structure. The computer-simulated HRTEM image,based on the hexagonal P63/m HA structure, is in good agreementwith the experimental image (Fig. 5D, inset).

Furthermore, there are several interesting aspects concerning thephase composition of 24 hour-derived CPC product resulting fromthe TEM investigation. A direct measurement of the d100 spacing onthe HRTEM image (Fig. 5D), gives a slightly smaller distance (~7.4 Å)than a similar measurement for the ideal hexagonal HA structure(~8.2 Å). The calculated ED pattern for hexagonal HA does not fitperfectly with the FT pattern (given as inset in Fig. 5D). Theseexperimental data suggest that the cement does not crystallize in theideal HA structure; the observed discrepancy can be explainedassuming a non-stoichiometry of the as-synthesized CPC, as deter-mined by EDX, Raman scattering and IR analyses.

To the first approximation, the potential bioactivity of the CPC –

ability to form a direct chemical bond with surrounding bone tissue –

was explored throughout soaking in SBF solution. This bone-bondingcapacity was evaluated by analyzing the formation of an apatite layeron the surface of the test pellet, revealing the cement apatite-inducingability [27]. Fig. 6A shows low- and high-magnification SEM images ofthe CPC surface prior soaking in SBF. The initial test sample exhibitssolid microstructure with macroporous surface. According to the SEM

Fig. 6. Comparison of the low-magnification (main panel) and high-magnification (right pasolution (B). A low-magnification (C) and high-magnification from selected regions (D, E) csurface of the CPC immersed in SBF.

analysis, the first isolatedmorphologies of precipitated apatite crystalsoccur already after 3 days of incubation in SBF (not shown), whereasafter 10 days of incubation, a widespread formation of sub-micrometer-sized scaly-like apatite crystals is observed (Fig. 6B).The morphology of the deposits is very similar to that reported for theapatite formed in SBF [13]. In addition, the phase purity of in vitrotested CPCs was confirmed by XRD. Furthermore, the crystals tend togenerate a relatively thick and persistent apatitic layer on thespecimen surface, as it can be seen in Fig. 6B–E, thus proving thepotentially high bone-bonding ability of the synthesized cement.

Cell culture experiment was used to elucidate the biocompatibilityof the derived cement biomaterial. For this purpose, the alkalinephosphatase (ALP) enzyme activitymeasurementswere carried out asa function of time to probe the differentiation of MC3T3-E1osteoblastic cells, and the results are represented in Fig. 7. As theduration of cell culture experiment increases, the ALP activitysystematically increases as well, evidencing that the pre-osteoblastsundergo a differentiation towardmature phenotype. Light microscopyimages for the Giemsa stained cement at various culture times areprovided in Fig. 8. After four days of MC3T3-E1 seeding, thefibroblastic-shaped osteoblasts are adhered to the CPC surfaceforming island-like appearances (Fig. 8A). In contrast, after twentyone days, the cells are found to intensively proliferate, leading to the

nel) SEM images from the surface of the CPCs soaked in distilled water (A) and in SBFross section SEM images are provided for clarification of the HA crystal growth on the

Fig. 7. Alkaline phosphatase (ALP) activity of MC3T3-E1 osteoblast cells cultured on theCPC specimen as a function of time.

Fig. 8. Light microscopy images from the Giemsa stained CPC surface after 4 (A) and 21(B) days of cell culture experiments.

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smooth covering of the whole exposed cement surface (Fig. 8B).Thus, from cell culture experiment, one can deduce that the presentCPC gives raise to the osteoblastic differentiation and promotesthe formation of extracellular matrix at later stage; hence, it isbiocompatible.

The ultimate compressive strength value of the CPC was deter-mined to be 25±3 MPa. The elastic modulus Es and hardness H of24 hour-derived CPC were estimated using nanoindentation (NI),since recent development of this method has enabled researchers toprobe these physical properties of the material at the nano- andmicro-structural levels. A representative NI load versus displacementprofile is shown in Fig. 9A. This load-displacement curve smoothlyfollows the loading functionwithout any discontinuities, steps or pop-inmarks, confirming that no cracks arise during nanoindentation [28].According to the displacement profile, the peak load of 5 mN results inresidual indentation depths ranging from approximately 250 to300 nm. The average elastic modulus and hardness of the cementwere calculated to be Es=23±3.5 and H=0.7±0.2 GPa,respectively.

The inset in Fig. 9A shows a typical FE-ESEM image of theCPC surfaceafter nanoindentation. According to the FE-ESEM, the NI probinggenerates equilateral triangle-shaped microindent impressions withtypical sides' lengths of about 4.5 µm. No cracks are observed by FE-ESEM, which is consistent with the smooth behaviour load-displace-ment curve (Fig. 9A). A view of the microindents can be observed moreclosely from the AFM topography imaging owing to the higher spatialresolution of this technique. A representative 3D AFM image of theregular triangle pyramid faceted impression is shown inFig. 9B. TheAFMthickness profile of the microindent (Fig. 9C) reveals that the surfaceexhibits a residual indentation depth relatively consistent with the oneobserved in the load-displacement curve (Fig. 9A). According to theAFMobservations, aflowof CPCmaterial above the edges of themicroindentsduring probing (so-called pile-up) is not detected, reflecting highcapacity of this biomaterial for work hardening.

4. Discussion

In reported literature of HA CPCs syntheses, the TTCP precursorsare not entirely converted to the apatitic cements even after 1 day ofsetting and in all cases the produced biomaterials still containedunreacted TTCP [16,29–32]. It is believed that unconsumed TTCP isrelated to the greater size of TTCP crystals cf. DCPD ones, and also can

be connected to the formation of HA layer on the surface of thereactants [33]. To the best of our knowledge, only Greish and Brownhave reported a cement formulation, which is totally converted to HAafter approximately 24 h of setting reaction [34]. This was achieved byusing mixing liquid containing phosphate anions, which significantlyincreased the saturation of the liquid phase with PO4

3- anionscompared to traditional water mixing liquid [31]. However, as-prepared CPC exhibits a high pH for at least 7 days, prone togenerating cell damage when employed in vivo [34].

The mechanochemically-derived sub-micrometer DCPD and TTCPprecursors are capable to store a certain amount of the receivedmechanical energy, therefore, they exhibit enhanced reactivity owingto their capability to consume this extra energy in different ways [21].The powder XRD patterns of the synthesized CPC at t≥6 h (Fig. 3) aswell as after 1 and 3 weeks (not shown) are alike. This indicates thatthe conversion rate of this CPC to HA is very high and the dissolution–precipitation reaction is fully complete already after 6 h of the cementsetting. To our knowledge, this is the highest reaction rate observed sofar for CPCs based on dicalcium and tetracalcium phosphates, animportant step toward enhanced biological and physiological char-acteristics. Interestingly, complete hydrolysis of dicalcium phosphateto HA can only be achieved in very dilute suspensions (S/L ratio ofabout 0.01 g/mL) [35], whereas TTCP with particles size of about1.4 µm (analogous to current study) is only partially hydrolysed to HAafter 1 day (S/L ratio of 2 g/mL) [22]. Therefore, it is believed that theobserved rapid formation of hydroxyapatite during CPC setting is a

Fig. 9. Representative load versus displacement curve (A) acquired by NI probing on the cement surface. The inset is an FE-ESEM image of a specimen surface after nanoindentation,showing the microindent impression. AFM three-dimensional topography view (B), along with the corresponding height profile (C), from the microindent impression on the cementsurface (lateral size: 10 µm by 10 µm; height: from 76.36 nm to 266.50 nm).

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result of a straightforward reaction between ball-milled dicalcium andtetracalcium phosphates and not product of their hydrolysis.

The potential bioactivity of the as-synthesized cement biomaterialis revealed applying in vitro SBF testing, while the cell cultureexperiment shows that the CPC is a suitable substrate for attachment,proliferation and differentiation of the pre-osteoblastic cell lineMC3T3-E1. This observation clearly confirms that CPC is biocompa-tible being a good substrate for osteoblastic cell growth.

The lack of literature related to NI probing of CPCs on the basis ofthe dicalcium and tetracalcium phosphates, did not allow us tocompare the observed results with reported data. In this context, fewreports have been published where tested biocomposite materialsinclude HA in a complex composition [36–38]. Es and H values of 15.2and 0.59 GPa, respectively, were achieved when a hydroxyapatitecement was used as a filler in a resin matrix, along with SiC whiskersas reinforcement phase [36]. Without the silicon carbide whiskers,lower Es and H values (11.8 and 0.41 GPa, respectively) were observed[36]. When a Sr-HA cement was used as a filler in a resin matrix, alongwith 5% SiO2 as reinforcement phase, Es values ranging fromapproximately 3.3 to 5.2 GPa were obtained [37,38]. The low valuesof Es and H reported for the aforementioned biomaterials [36–38] incomparison to those obtained in the present work seem likely to beassociated with the presence of resins, since they are less stiff thancrystalline solids and, accordingly, exhibit lower elastic modules.

Ideally, Es and H of the biomaterials, that are meant to be used asbone grafts, shouldmatch that of humanbones. The synthesized cementexhibits an elastic modulus (23 ± 3.5 GPa) concordant with those ofaverage human cortical bone (20–25.8 GPa) [39] and slightly higherthan in human trabecular bones (15–19.4 GPa) [40]. Following Es trend,the hardness of this CPC (0.7±0.2 GPa) is also higher than in humantrabecular bones (0.52–0.62 GPa) and in consistency with H of humancortical bones (0.62–0.74 GPa) [40]. Overall, the observed results clearlydemonstrate a high elastic and plastic similarity at the nano-/micro-level between the synthesized cement biomaterial and human bone,thus matching an essential requirement of biomaterials for bonesubstitution.

5. Conclusions

Advanced calcium phosphate cement was effectively synthesizedfrom ball-milled DCPD and TTCP powders. This approach yields a CPCproduct that sets in ~22 min and entirely converts to the end productalready after 6 h of setting reaction. Detailed experimental studies

conclusively support that synthesized cement biomaterial is nano-structured, potentially bioactive and biocompatible Ca-deficient B-type carbonated hydroxyapatite. This work further highlight that thelocal elastic modulus and hardness of the cement are very similar tothose of human bones, suggesting its essential isoelastic performancein bone repair and substitution uses.

Acknowledgment

The authors thank C. Pilz and K. P. Kommareddy for help with cellculture experiment. We also thank Dr. A. V. Olenev, Dr. K. A. Kovnir,and T. P. B. Cotter for helpful discussions. This work was partiallysupported by the Spanish Ministry of Education and Science(MAT2002-03857), the European Union Marie Curie EST Fellowshipon Biomimetic Systems (MEST-CT-2004-504465) and the EuropeanUnion Framework 6 Program under a contract for an IntegratedInfrastructure Initiative (Reference 026019 ESTEEM).

References

[1] J.P. Schmit, J.O. Hollinger, S.B. Milan, J. Oral Maxillofac. Surg. 57 (1999) 1122.[2] S. Larsson, T.W. Bauer, Clin. Orthop. Relat. Res. 395 (2002) 23.[3] K. Kurashina, H. Kurita, M. Hirano, J.M.A. Blieck, C.P.A.T. Klein, K. de Groot, J. Mater.

Sci. Mater. Med. 6 (1995) 340.[4] K. Kurashina, H. Kurita, M. Hirano, A. Kotani, C.P.A.T. Klein, K. de Groot, Biomaterials

18 (1997) 539.[5] F.C.M. Driessens, J.A. Planell, M.G. Boltong, I. Khairoun, M.P. Ginebra, Proc. Inst.

Mech. Eng. H. 212 (1998) 427.[6] L.L. Hench, J. Am. Ceram. Soc. 74 (1991) 1487.[7] S.V. Dorozhkin, M. Epple, Angew. Chem. Int. Ed. 41 (2002) 2130.[8] J.E. Barralet, L. Grover, T. Gaunt, A.J. Wright, I.R. Gibson, Biomaterials 23 (2002) 3063.[9] Y. Miyamoto, K. Ishikawa, H. Fukao, M. Sawada, M. Nagayama, M. Kon, K. Asaoka,

Biomaterials 16 (1995) 855.[10] J.-Y. Rho, L. Kuhn-Spearing, P. Zioupos, Med. Eng. Phys. 20 (1998) 92.[11] G. Lewis, J.S. Nyman, J. Biomed. Mater. Res., B 87B (2008) 286.[12] V. Biehl, J. Breme, Mat.-wiss. u. Werkstofftech. 32 (2001) 137.[13] T. Kokubo, H. Kushitani, S. Sakka, T. Kitsugi, T. Yamamuro, J. Biomed. Mater. Res. 24

(1990) 721.[14] E.C. Moreno, W.E. Brown, G. Osborn, Soil Sci. Soc. Am. J. 24 (1960) 94.[15] W.E. Brown, E.F. Epstein, J. Res. Natl. Bur. Stand. 69A (1965) 547.[16] W.E. Brown, L.C. Chow, US Patent No. 4518430, 1985.[17] I. Manjubala, A.Woesz, C. Pilz, M. Rumpler, N. Fratzl-Zelman, P. Roschger, J. Stampfl,

P. Fratzl, J. Mater. Sci. Mater. Med. 16 (2005) 1111.[18] W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564.[19] C.H. Turner, D.B. Burr, in: S.C. Cowin (Ed.), Bone mechanics handbook, CPC Press,

Boca Raton, FL, 2001, p. 7.[20] D.E. Grenoble, J.L. Katz, K.L. Dunn, R.S. Gilmore, K.L. Murty, J. Biomed. Mater. Res.

(1972) 3.[21] M. Senna, Part. Part. Syst. Charact. 6 (1989) 163.[22] U. Gbureck, J.E. Barralet, M. Hofmann, R. Thull, J. Am. Ceram. Soc. 87 (2004) 311.[23] M. Bohner, J. Mater. Chem. 17 (2007) 3980.

2132 I.S. Neira et al. / Materials Science and Engineering C 29 (2009) 2124–2132

[24] S. Koutsopoulos, J. Biomed. Mater. Res. 62 (2002) 600.[25] K. Nakamoto, Infrared and Raman spectra of inorganic and coordination

compounds, Part B: applications in coordination, organometallic and bioinorganicchemistry, John Wiley & Sons Inc, New York, 1997, p. 54.

[26] H. Monma, T. Takahashi, Gypsum & Lime 210 (1987) 287–291.[27] L.L. Hench, A.E. Clark, in: D.F. Williams (Ed.), Biocompatibility of orthopedic

implants, vol. II, CRC Press, Boca Raton, FL, 1982, p. 129.[28] X. Li, B. Bhushan, Mater. Charact. 48 (2002) 11.[29] E.F. Burguera, F. Guitián, L.C. Chow, J. Biomed. Mater. Res., A 71A (2004) 275.[30] Y. Fukase, E.D. Eanes, S. Takagi, L.C. Chow,W.E. Brown, J. Dent. Res. 69 (1990) 1852.[31] K. Ishikawa, S. Takagi, L.C. Chow, Y. Ishikawa, J. Mater. Sci. Mater. Med. 6 (1995) 528.[32] L. Sun, L.C. Chow, T.W. Bauer, E.F. Burguera, S.A. Frukhtbeyn, Key Eng. Mat. 284–286

(2005) 23.

[33] E.F. Burguera, F. Guitián, L.C. Chow, Key Eng. Mat. 284–286 (2005) 15.[34] Y.E. Greish, P.W. Brown, J. Biomed. Mater. Res., B 67B (2003) 632.[35] A. Perloff, A.S. Posner, Science 124 (1956) 583.[36] H. Xu, D. Smith, C. Simon, Biomaterials 25 (2004) 4615.[37] G.X. Ni, Y.S. Choy, W.W. Lu, A.H.W. Ngan, K.Y. Chiu, Z.Y. Li, B. Tang, K.D.K. Luk,

Biomaterials 27 (2006) 1963.[38] G.X. Ni, W.W. Lu, B. Tang, A.H.W. Ngan, K.Y. Chiu, K.M.C. Cheung, Z.Y. Li, K.D.K. Luk,

J. Biomed. Mater. Res, A 83A (2007) 570.[39] X.E. Guo, in: S.C. Cowin (Ed.), Bone mechanics handbook, CPC Press, Boca Raton,

FL, 2001, p. 10.[40] J.-Y. Rho, M.E. Roy, T.Y. Tsui, G.M. Pharr, J. Biomed. Mater. Res., A 45 (1999) 48.