Rational synthesis of a nanocrystalline calcium phosphate cement ...

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

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Materials Science and Engineering C j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / m s e c

Rational synthesis of a nanocrystalline calcium phosphate cement exhibiting rapid conversion 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, Spain Materials and Structures Laboratory, Tokyo Institute of Technology, 4259 Nagatsuta, Midori-ku, 226-8503 Yokohama, Japan Department of Biomaterials, Max Planck Institute of Colloids and Interfaces, D-14424 Potsdam, Germany d Electron Microscopy for Materials Science, University of Antwerp, B-2020 Antwerp, Belgium b c

a r t i c l e

i n f o

Article history: Received 2 February 2009 Received in revised form 21 March 2009 Accepted 15 April 2009 Available online 19 April 2009 Keywords: Calcium phosphate cement Rapid conversion Nanostructure Mechanical properties Nanoindentation Osteoblast cell culture

a b s t r a c t The rational synthesis, comprehensive characterization, and mechanical and micromechanical properties of a calcium phosphate cement are presented. Hydroxyapatite cement biomaterial was synthesized from reactive sub-micrometer-sized dicalcium phosphate dihydrate and tetracalcium phosphate via a dissolutionprecipitation reaction using water as the liquid phase. As a result nanostructured, Ca-deficient and carbonated B-type hydroxyapatite is formed. The cement shows good processibility, sets in 22 ± 2 min and entirely transforms to the end product after 6 h of setting reaction, one of the highest conversion rates among previously reported for calcium phosphate cements based on dicalcium and tetracalcium phosphates. The combination of all elucidated physical-chemical traits leads to an essential bioactivity and biocompatibility of the 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 elasticity and plasticity at sub-micrometer/micrometer level. The measured elastic modulus and hardness were established to be Es = 23 ± 3.5 and H = 0.7 ± 0.2 GPa, respectively. These values are in close agreement with those reported in literature for trabecular and cortical bones, reflecting good elastic and plastic coherence between synthesized cement biomaterial and human bones. © 2009 Elsevier B.V. All rights reserved.

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

⁎ Corresponding author. Mailing address: Instituto de Cerámica de Galicia, Universidade de Santiago de Compostela, Avda. Mestre Mateo, s/n, E-15782, Santiago de Compostela, Spain. Tel.: +34 981 563100x16885; fax: +34 981 564242. E-mail address: [email protected] (I.S. Neira). 1 Present address: Department of Inorganic Chemistry, Fritz Haber Institute of the Max Planck Society, D-14195 Berlin, Germany. 0928-4931/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msec.2009.04.011

applications [7], or as materials for the development of scaffolds in bone tissue engineering [8]. From a fundamental and an application point of view, it is important to design a rational synthesis route to a cement with setting time not more than thirty minutes (surgeons' requirements) and a rapid conversion rate to the end product, preferably thermodynamically stable hydroxyapatite (HA) — the inorganic component of mineralized bone tissues. It is noteworthy that in many cases the rapid conversion of the CPC precursors to HA not only leads to the rise of the appropriate physiological properties, but also results in the strengthening of the mechanical properties. In particular, rapid conversion accelerates achievement of the final compressive strength value, which is almost linearly dependent on the extent of the CPC setting reaction [9]. However, the preparation of CPCs with the aforementioned characteristics still remains challenging. Bones are hierarchically organized biological systems originated at the nanometre scale [10]. Therefore, it is beneficial to investigate the mechanical behaviour of the cements at the nano- and microstructural levels with regards to its application for bone graft uses. Such information is important for the exploring of mechanistic compatibility between functional cement biomaterials and bones.

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

Recent advances in this area have demonstrated that nanoindentation probing is one of the most promising techniques which allows to estimate such important mechanical properties of the materials as hardness and elastic modulus at unprecedented scale [11]. An elastic modulus similar to that of the human bones is an essential property for bone grafts, providing a good load transfer from the implant into the bone, which is required for the stimulation of new bone formation [12]. In contrast, if biomaterial is much stiffer than bone, it shields the nearby bone from mechanical stress, generating a sort of disuse atrophy — the bone resorbs. Herein, we present the rational synthesis of a cement together with the detailed characterization of the structure at the micrometer and at the nanometer level. Elastic modulus and hardness of derived biomaterial were also investigated and compared to those of bone tissue 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 nonylphenylpolyethylenglycol acetate (for synthesis, Fluka) were used as received. The preparation of acellular simulated body fluid (SBF) solution having pH 7.4 was adopted from that of Kokubo et al. [13]. 2.2. Preparation of the CPC precursors Dicalcium phosphate dihydrate (DCPD) was synthesized via a precipitation reaction between Ca(H2PO4)2·H2O and CaCO3 [14], while tetracalcium phosphate (TTCP) was prepared from stoichiometric amounts CaCO3 and CaHPO4 by solid-state reaction [15]. It is reasonable to expect that the reactivity of the DCPD and TTCP will increase with decreasing of their crystal sizes that, accordingly, will facilitate the formation of respective cement. On the basis of this concept, the as-synthesized raw materials were subjected to a wet rolling ball-milling to reduce the size of the particles. First, one third of a 5% silica-doped alumina pot (V = 300 mL) was filled with silicon nitride balls of 3.6 mm Ø in case of DCPD grinding or with high purity zirconia balls of 4.8 mm Ø in case of TTCP milling. Afterwards, 15 g of the powder was placed inside the pot, along with 60 mL of distilled water (for DCPD milling) or anhydrous isopropanol (for TTCP milling). The ball-milling was performed at 135 rpm for 24 h. The product was collected by vacuum filtration and washed with acetone. DCPD was dried at the ambient conditions, while TTCP was dried at 80 °C on air.

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2.4. Characterization methods The products were characterized by powder X-ray diffraction (XRD) using a Rigaku RINT 2000 diffractometer. The phases were identified by comparison with 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 S4500 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 dynamic light-scattering (DLS) technique on a Malvern Instruments Zetasizer Nano-ZS. For DLS measurements, a small amount of the powder was fully dispersed in water using high-power ultrasonication, and the resultant stable suspension was subjected to the DLS analysis. Final setting time (ST) of the cement was determined according to the American Dental Association Specification No. 9 using the Vicat needle method. ST value is the average of at least six replicates. Energydispersive X-ray spectroscopy (EDX) for semiquantitative Ca and P content determination was performed with a EDAX DX-95 spectrometry system. The room temperature Raman scattering measurements were carried out on a confocal Raman microscope WITEC CRM200 with linear polarized laser light (λ = 532 nm) as the excitation source. The room temperature diffuse reflectance infrared (IR) Fourier-transform spectra were recorded on a Jeol JIR-7000 spectrometer. The fine microstructure and phase purity of the cement were investigated by transmission electron microscopy (TEM), electron diffraction (ED) and high-resolution TEM (HRTEM) using Jeol 4000EX microscope operating at 400 kV. The samples for TEM were crushed, dispersed in methanol and deposited on a holey carbon grid. Computer-simulated HRTEM images were obtained using the Mac Tempas and Crystal Kit software programs. 2.5. In vitro testing In vitro bioactivity investigation was carried out using SBF solution. The CPC test pellets were placed into a polystyrene bottle filled with SBF and incubated at 37 °C. The polystyrene bottle and SBF solution were refreshed every 3 days. Biocompatibility was studied in vitro through cell culture experiment adopted from that of Manjubala et al. [17] using murine pre-osteoblastic cells (MC3T3-E1) from mouse calvarie. 5 × 105 cells were suspended in 100 µL of culture medium and seeded on the CPC test specimens in the form of pellets. The specimens were incubated at 37 °C for 30 min in Petri dishes to allow the cells to attach to the surface and then 2 ml of medium were added. The medium and Petri dish were refreshed twice a week. Alkaline phosphatase (ALP) enzyme activity was quantitatively determined after 1, 2 and 3 weeks. The cell proliferation and distribution over the CPC surface were analyzed using a light microscope employing Giemsa staining histological method.

2.3. Preparation of the calcium phosphate cement 2.6. Mechanical properties measurements CPC was generated by the dissolution-precipitation reaction between DCPD and TTCP [16]. Equimolar amounts of ball-milled powders 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 resultant putty-like material was rod-shaped, allowed to set for 1 h at ambient conditions and incubated in distilled water at human body temperature of 37 °C for the desired period of time. The CPC specimens for compression strength and nanoindentation tests were prepared, in the form of pellets, using S/L ratio of 3 g/mL. The as-derived putty-like material was placed into a cylindrical mold assembly of 6.9 mm Ø. The sample was pressed for 1 min under 5.3 MPa (lowest value of pressure scale of used loading device). The obtained pellets were kept for 1 h at ambient conditions to set and then incubated in a SBF solution at 37 °C for 23 h.

Compressive strength (CS) measurements were performed using a Shimadzu Autograph AG-I universal testing machine. Prior to CS experiment, 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 ambient conditions and then immersed in distilled water for 23 h at 37 °C. CS value is the average of at least six replicates. Elastic modulus and hardness of the cement were measured using nanoindentation technique (NI). For this purpose, the CPC pellet was epoxy embedded in a solution formed by 77% of methylmethacrylate and the polymerization initiators — 1.4% of benzoylperoxid and 21.6% of nonylphenylpolyethylenglycol acetate. To fill up the pores, epoxy embedding was achieved by gradual polymerization of methylmethacrylate to poly(methylmethacrylate) applying a low temperature profile [(39 °C, 12 h) -> (48 °C, 12 h) -> (55 °C, 12 h)]. Then, the section

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for NI was cut in parallel to the pellet base, polished by a set of abrasive papers of decreasing size of grit followed by polishing with diamond paste down to a particle size of 1 µm. Quasi-static nanoindentation tests were performed on a Hysitron TriboScan UBI-1 nanohardness tester in conjunction with a Digital Instruments Nanoscope III atomic force microscope (AFM). A Berkovich diamond indenter tip with regular triangle pyramid geometry and nominal radius of 300 nm was used. A loading function with a peak load of 5 mN, loading rate of 1 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 of 0.5 mN/s, was used. Evaluation of the load-displacement data obtained by NI was performed using the software program implemented in the TriboScan, which applies Oliver and Pharr method [18] to calculate the reduced indentation modulus Er and hardness H. During unloading, the drift effect was observed and corrected by the same software program. Er is related to elastic modulus Es through the Eq. (1):

1 = Er

2 1 − mi Ei

+

2 1 − ms Es

ð1Þ

the nanocrystalline nature and defective HA structure of the cement (vide infra). The average molar ratio of Ca to P elements (Ca/P) in stoichiometric HA is 1.67, while Ca/P ratio in the 24 hour-derived CPC was established to be ~1.49 (EDX), indicating that as-synthesized cement is Ca-deficient HA. The observed non-stoichiometry is mainly a result 2of partial substitution of the PO34 by HPO4 in the HA crystal structure, as established by Raman scattering (not shown). Fig. 4 displays a representative IR spectrum for the 24 hour-derived cement. The set of bands and spectra features agrees fairly well with the reported IR data for HA [24]. The broad peak at ~ 3406 cm- 1 reflects the physisorbed water, peaks observed at ~ 2358 and 2339 cm- 1 correspond to the atmospheric CO2(g), while the appearance of the H-O-H deformation band at ~1643 cm- 1 suggests the existence of free water molecules trapped in the crystal lattice of the HA cement [25]. Additionally, IR spectroscopy confirms the incorporation of CO23 anions into the structure of HA (Fig. 4), indicating that the obtained HA cement is not only Ca-deficient, but also B-type carbonated HA, thus, with the CO23 partially substituting PO3− groups [26]. These carbonate anions are 4 believe to come from the synthesis procedures, namely, from the

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 hundred fifty indents were randomly performed to obtain average values for Er and H. 3. Results According to the XRD analysis, phase-pure DCPD and TTCP compounds were produced by precipitation and high temperature reactions, respectively (Fig. 1A and B). In order to improve the CPC precursors contact and to enhance their reactivity, the primary crystal sizes were reduced by mechanochemical grinding. XRD reveals that the ball-milling does not lead to changes in the phase composition and the 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 TTCP product, a broadening of the XRD peaks is observed (Fig. 1B). This is likely due to the comminution of agglomerates, reduction of crystallites sizes and crystal structure deformation from the ball-milling procedure, overall resulting in partial amorphization of the ball-milled material [21], in good agreement with reported data [22]. SEM and DLS analyses of the ball-milled products clearly reveal almost nonaggregated DCPD and TTCP particles with average particle sizes approximately 1.1 and 1.5 µm, respectively, exhibiting relatively narrow PSD (Fig. 2). The CPC was generated from these reactive sub-micrometer-sized precursors via a dissolution-precipitation reaction using water as the liquid phase. The setting time study of the cement indicates that as-prepared CPC sets within 22 ± 2 min. It is noteworthy that this quick ST was achieved by using exclusively water as the liquid mixing phase, i.e., without any of the setting accelerators widely used in cement formulations [23]. The cement phase composition evolution during the setting reaction at the selected times (t) of 0–8 and 24 h is shown in Fig. 3. According to XRD, the CPC at t = 0 consists of, as expected, pure DCPD and TTCP (ICDD No. 9–77 and 70–1379, respectively). Further phase analysis reveals total consumption of DCPD already at t = 4 h, while TTCP is fully exhausted at t = 6 h. At this time of reaction (t = 6 h) and higher, cement product consists of only phase-pure HA (ICDD No. 72-1243). The powder XRD patterns of the 6–8 and 24 hour-derived HA cements exhibit low intensities and very broad peaks (Fig. 3). This observation would suggest that these samples are essentially amorphous; however, as established by TEM investigation, these CPCs are crystalline. Hence, aforementioned features of the XRD patterns are most likely caused by

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


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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 the most intense diffraction peaks of DCPD, TTCP and HA.

Fig. 4. IR spectrum collected from the 24 hour-derived cement product. The characteristic bands representing apatitic phosphate PO34 and hydroxyl OH groups; lattice (deformation H-O-H band) and physisorbed water; as well as atmospheric CO2(g) are marked by grey regions. Dotted lines are drawn at the positions of CO23 anion vibrations to indicate B-type (carbonate replacing phosphate) of carbonated HA.

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Fig. 5. SEM micrographs 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 the 24 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 the channels.

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 for biomedical application a carbonated HA cement is rather a requirement than a problem. A SEM image of cement product after 1 h of setting is shown in Fig. 5A. It displays a highly aggregated microstructure with visible alteration of the facets of the ball-milled precursors, i.e., owing to the smoothing of edges, the grains become less faceted in comparison to the primary morphologies (Fig. 2, left panels). Since the CPC precursors were used as faceted non-agglomerated particles, the above results show that the partial dissolution of the CPC reagents already occurs at this stage of the reaction (t = 1). The SEM image of 24 hour-derived cement (Fig. 5B) reveals highly associated microstructure of the product with clearly visible macropores (Ø N 50 nm). SEM observations also reveal the lack of the particulate structure regions in this CPC, and that the surface of the cement is free of any crystal faceting. These effects show that the dissolution-precipitation reaction between ball-milled DCPD and TTCP is complete. Low-magnification and high-resolution TEM images of 24 hourderived CPC are shown in Fig. 5C and D, respectively. The sample consists of interlocking nanoscaled platelet crystals. The compact packaging of the nanocrystallites is subsequently confirmed to be a

common feature in the cement. Namely such spontaneous interlocking of the precipitated nanocrystals results in the structural solidity of the cement. The crystallites size lies within the range of 10–15 nm and their surfaces are free of any amorphous or secondary phase (Fig. 5C). These results indicate that the line broadening effect and low intensity of the cement XRD patterns (Fig. 3) are likely related to the nanocrystalline nature of the CPC end product and not associated with 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 nanocrystallite clustering (Fig. 5C, inset). The ED pattern closely resembles the HA 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 ED pattern can be indexed using the characteristic d-spacings of the hexagonal HA phase (d211, d102, and d002) determined from powder XRD analysis of this cement product. The very weak intensity of the d100 and d101 rings representing large d spacing in the ED pattern (Fig. 5C, inset) is related to the platelet shape of the nanocrystals resulting in a preferential orientation of the nanocrystallites. As determined by HRTEM, the nanoparticles in the as-synthesized sample preferentially growth along the c-axis, but there are nanocrystallites grown along uncommon 012 direction (Fig. 5D). The corresponding Fourier transform (FT) of the nanocrystallite is


shown as the top inset in Fig. 5D. This pattern can be indexed according to the HA structure. The computer-simulated HRTEM image, based on the hexagonal P63/m HA structure, is in good agreement with the experimental image (Fig. 5D, inset). Furthermore, there are several interesting aspects concerning the phase composition of 24 hour-derived CPC product resulting from the TEM investigation. A direct measurement of the d100 spacing on the 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 fit perfectly with the FT pattern (given as inset in Fig. 5D). These experimental data suggest that the cement does not crystallize in the ideal HA structure; the observed discrepancy can be explained assuming a non-stoichiometry of the as-synthesized CPC, as determined 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-bonding capacity was evaluated by analyzing the formation of an apatite layer on the surface of the test pellet, revealing the cement apatite-inducing ability [27]. Fig. 6A shows low- and high-magnification SEM images of the CPC surface prior soaking in SBF. The initial test sample exhibits solid microstructure with macroporous surface. According to the SEM

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analysis, the first isolated morphologies of precipitated apatite crystals occur already after 3 days of incubation in SBF (not shown), whereas after 10 days of incubation, a widespread formation of submicrometer-sized scaly-like apatite crystals is observed (Fig. 6B). The morphology of the deposits is very similar to that reported for the apatite formed in SBF [13]. In addition, the phase purity of in vitro tested CPCs was confirmed by XRD. Furthermore, the crystals tend to generate a relatively thick and persistent apatitic layer on the specimen surface, as it can be seen in Fig. 6B–E, thus proving the potentially high bone-bonding ability of the synthesized cement. Cell culture experiment was used to elucidate the biocompatibility of the derived cement biomaterial. For this purpose, the alkaline phosphatase (ALP) enzyme activity measurements were carried out as a function of time to probe the differentiation of MC3T3-E1 osteoblastic cells, and the results are represented in Fig. 7. As the duration of cell culture experiment increases, the ALP activity systematically increases as well, evidencing that the pre-osteoblasts undergo a differentiation toward mature phenotype. Light microscopy images for the Giemsa stained cement at various culture times are provided in Fig. 8. After four days of MC3T3-E1 seeding, the fibroblastic-shaped osteoblasts are adhered to the CPC surface forming island-like appearances (Fig. 8A). In contrast, after twenty one days, the cells are found to intensively proliferate, leading to the

Fig. 6. Comparison of the low-magnification (main panel) and high-magnification (right panel) SEM images from the surface of the CPCs soaked in distilled water (A) and in SBF solution (B). A low-magnification (C) and high-magnification from selected regions (D, E) cross section SEM images are provided for clarification of the HA crystal growth on the surface of the CPC immersed in SBF.

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

smooth covering of the whole exposed cement surface (Fig. 8B). Thus, from cell culture experiment, one can deduce that the present CPC gives raise to the osteoblastic differentiation and promotes the formation of extracellular matrix at later stage; hence, it is biocompatible. The ultimate compressive strength value of the CPC was determined to be 25 ± 3 MPa. The elastic modulus Es and hardness H of 24 hour-derived CPC were estimated using nanoindentation (NI), since recent development of this method has enabled researchers to probe these physical properties of the material at the nano- and micro-structural levels. A representative NI load versus displacement profile is shown in Fig. 9A. This load-displacement curve smoothly follows the loading function without any discontinuities, steps or popin marks, confirming that no cracks arise during nanoindentation [28]. According to the displacement profile, the peak load of 5 mN results in residual indentation depths ranging from approximately 250 to 300 nm. The average elastic modulus and hardness of the cement were calculated to be E s = 23 ± 3.5 and H = 0.7 ± 0.2 GPa, respectively. The inset in Fig. 9A shows a typical FE-ESEM image of the CPC surface after nanoindentation. According to the FE-ESEM, the NI probing generates equilateral triangle-shaped microindent impressions with typical sides' lengths of about 4.5 µm. No cracks are observed by FEESEM, which is consistent with the smooth behaviour load-displacement curve (Fig. 9A). A view of the microindents can be observed more closely from the AFM topography imaging owing to the higher spatial resolution of this technique. A representative 3D AFM image of the regular triangle pyramid faceted impression is shown in Fig. 9B. The AFM thickness profile of the microindent (Fig. 9C) reveals that the surface exhibits a residual indentation depth relatively consistent with the one observed in the load-displacement curve (Fig. 9A). According to the AFM observations, a flow of CPC material above the edges of the microindents during probing (so-called pile-up) is not detected, reflecting high capacity of this biomaterial for work hardening.

be connected to the formation of HA layer on the surface of the reactants [33]. To the best of our knowledge, only Greish and Brown have reported a cement formulation, which is totally converted to HA after approximately 24 h of setting reaction [34]. This was achieved by using mixing liquid containing phosphate anions, which significantly increased the saturation of the liquid phase with PO34 anions compared to traditional water mixing liquid [31]. However, asprepared CPC exhibits a high pH for at least 7 days, prone to generating cell damage when employed in vivo [34]. The mechanochemically-derived sub-micrometer DCPD and TTCP precursors are capable to store a certain amount of the received mechanical energy, therefore, they exhibit enhanced reactivity owing to 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) as well as after 1 and 3 weeks (not shown) are alike. This indicates that the conversion rate of this CPC to HA is very high and the dissolution– precipitation reaction is fully complete already after 6 h of the cement setting. To our knowledge, this is the highest reaction rate observed so far for CPCs based on dicalcium and tetracalcium phosphates, an important step toward enhanced biological and physiological characteristics. Interestingly, complete hydrolysis of dicalcium phosphate to HA can only be achieved in very dilute suspensions (S/L ratio of about 0.01 g/mL) [35], whereas TTCP with particles size of about 1.4 µm (analogous to current study) is only partially hydrolysed to HA after 1 day (S/L ratio of 2 g/mL) [22]. Therefore, it is believed that the observed rapid formation of hydroxyapatite during CPC setting is a

4. Discussion In reported literature of HA CPCs syntheses, the TTCP precursors are not entirely converted to the apatitic cements even after 1 day of setting and in all cases the produced biomaterials still contained unreacted TTCP [16,29–32]. It is believed that unconsumed TTCP is related to the greater size of TTCP crystals cf. DCPD ones, and also can

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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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 cement surface (lateral size: 10 µm by 10 µm; height: from 76.36 nm to 266.50 nm).

result of a straightforward reaction between ball-milled dicalcium and tetracalcium phosphates and not product of their hydrolysis. The potential bioactivity of the as-synthesized cement biomaterial is revealed applying in vitro SBF testing, while the cell culture experiment shows that the CPC is a suitable substrate for attachment, proliferation and differentiation of the pre-osteoblastic cell line MC3T3-E1. This observation clearly confirms that CPC is biocompatible being a good substrate for osteoblastic cell growth. The lack of literature related to NI probing of CPCs on the basis of the dicalcium and tetracalcium phosphates, did not allow us to compare the observed results with reported data. In this context, few reports have been published where tested biocomposite materials include HA in a complex composition [36–38]. Es and H values of 15.2 and 0.59 GPa, respectively, were achieved when a hydroxyapatite cement was used as a filler in a resin matrix, along with SiC whiskers as 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, along with 5% SiO2 as reinforcement phase, Es values ranging from approximately 3.3 to 5.2 GPa were obtained [37,38]. The low values of Es and H reported for the aforementioned biomaterials [36–38] in comparison to those obtained in the present work seem likely to be associated with the presence of resins, since they are less stiff than crystalline solids and, accordingly, exhibit lower elastic modules. Ideally, Es and H of the biomaterials, that are meant to be used as bone grafts, should match that of human bones. The synthesized cement exhibits an elastic modulus (23 ± 3.5 GPa) concordant with those of average human cortical bone (20–25.8 GPa) [39] and slightly higher than 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 human trabecular bones (0.52–0.62 GPa) and in consistency with H of human cortical bones (0.62–0.74 GPa) [40]. Overall, the observed results clearly demonstrate a high elastic and plastic similarity at the nano-/microlevel between the synthesized cement biomaterial and human bone, thus matching an essential requirement of biomaterials for bone substitution. 5. Conclusions Advanced calcium phosphate cement was effectively synthesized from ball-milled DCPD and TTCP powders. This approach yields a CPC product that sets in ~22 min and entirely converts to the end product already after 6 h of setting reaction. Detailed experimental studies

conclusively support that synthesized cement biomaterial is nanostructured, potentially bioactive and biocompatible Ca-deficient Btype carbonated hydroxyapatite. This work further highlight that the local elastic modulus and hardness of the cement are very similar to those of human bones, suggesting its essential isoelastic performance in bone repair and substitution uses. Acknowledgment The authors thank C. Pilz and K. P. Kommareddy for help with cell culture experiment. We also thank Dr. A. V. Olenev, Dr. K. A. Kovnir, and T. P. B. Cotter for helpful discussions. This work was partially supported by the Spanish Ministry of Education and Science (MAT2002-03857), the European Union Marie Curie EST Fellowship on Biomimetic Systems (MEST-CT-2004-504465) and the European Union Framework 6 Program under a contract for an Integrated Infrastructure 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


[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 bioinorganic chemistry, 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] [34] [35] [36] [37]

E.F. Burguera, F. Guitián, L.C. Chow, Key Eng. Mat. 284–286 (2005) 15. Y.E. Greish, P.W. Brown, J. Biomed. Mater. Res., B 67B (2003) 632. A. Perloff, A.S. Posner, Science 124 (1956) 583. H. Xu, D. Smith, C. Simon, Biomaterials 25 (2004) 4615. 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.