Preparation and characterization of bioceramics produced from calcium phosphate cements

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Crystal Researchand Technology

Journal of Experimental and Industrial Crystallography

Zeitschrift für experimentelle und technische Kristallographie

Editorial BoardR. Fornari, BerlinP. Görnert, JenaM. Watanabe, TokyoK. Sangwal, Lublin

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Cryst. Res. Technol. 45, No. 3, 239 – 243 (2010) / DOI 10.1002/crat.200900551

© 2010 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

Preparation and characterization of bioceramics produced from

calcium phosphate cements

O. Andriotis1, O. L. Katsamenis

1, D. E. Mouzakis

2, and N. Bouropoulos*

1,3

1 Department of Materials Science, University of Patras, 26504, Patras, Greece 2 Technological Educational Institute of Larisa, Department of Mechanical Engineering, T.E.I of Larissa, 411

10, Larissa, Greece 3 Foundation for Research and Technology, Hellas, Institute of Chemical Engineering and High Temperature

Chemical Processes, FORTH/ICE-HT, P.O. Box 1414, 26504 Rio Patras, Greece

Received 15 September 2009, revised 12 December 2009, accepted 14 December 2009

Published online 15 January 2010

Key words calcium phosphates, bioceramics, bone substitutes.

The present work reports a method for preparing calcium phosphate ceramics by calcination of calcium

phosphate cements composed mainly of calcium deficient hydroxyapatite (CDHA). It was found that

hardened cements calcinied at temperatures from to 600 °C to 1300 °C were transformed to tricalcium

phosphates. Moreover the compressive strength was determined and porosity was estimated as a function of

the calcination temperature.

© 2010 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

1 Introduction

Bone defects can be caused by injuries, tumor resection or various diseases. Bone grafting is the procedure to

place new bone into defects with the objective of osseous regeneration. Nowadays, different types of synthetic

materials such as polymers, ceramics and glasses are being used in osseous defect sites as bone substitutes [1-

4]. Calcium phosphate ceramics such as hydroxyapatite (HAP, Ca10(PO4)6(OH)2) or tricalcium phosphate (β-

and α-TCP, Ca3(PO4)2) are among the materials that are currently widely used in clinical practice as bone

fillers due to their excellent biocompatibility and osteoconductivity [5-6]. The most common method to

prepare calcium phosphate bioceramics, is by calcination of biological or synthetic apatite at temperatures

above 700 °C [7]. In the present work an alternative approach of preparing calcium phosphate ceramics is

proposed based on the calcination of calcium phosphate cements (CPSs). CPCs are biomaterials that may

directly initiate osteogenesis or promote osteoconduction when placed in direct contact with host bone.

Mechanical properties of CPSs depend on the crystallinity of CDHA and the porosity. They can be used as

bone fillers in craniofacial and maxillofacial surgery [8-9]. Using calcium phosphate cements composed

mainly of CDHA we prepared tricalcium phosphate bioceramics. Next, structural and morphological

characterization of the products was performed using X-ray diffraction and Scanning Electron Microscopy.

Finally mechanical tests were carried out on the specimens. Mechanical properties of the cements are reported

as a function of the calcination temperature.

2 Experimental

Preparation of α-TCP powder, cement setting and calcination For the preparation of α-TCP powder,

equimolar quantities of calcium carbonate and calcium pyrophosphate were mixed under magnetic stirring in

an ethanolic suspension. Next the mixture was dried at 80 °C, placed in a furnace at 1300 °C for 12 h and ____________________

* Corresponding author: e-mail: [email protected]

240 O. Andriotis et al.: Bioceramics produced from calcium phosphate cements

© 2010 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim www.crt-journal.org

rapidly quenched on a metallic surface. The resulted material was crushed in an agate mortar and placed in a

ball mill (Pulverisette 5, Fritsch, Germany). Ball milling was performed at 500 rpm using 1 cm diameter agate

spheres for 5 cycles of 20 min.

The milled powder was mixed with 2.5 % w/v Na2HPO4 solution in a agate mortar at powder/liquid ratio of

0.32 and the product in form of a paste, was moulded in 6 mm diameter by 12 mm height Teflon split moulds.

The specimens were remained in 100% humidity for 12 h and were then removed from the mould and placed at

37 °C in a polyethylene vial containing 60 ml of Ringer solution for 20 days for hardening. Afterwards, the

samples were air dried and calcined at 600, 800, 1000, 1100, 1200 and 1300 °C for 5 h with a heating rate of

5 K/min. Volume shrinkage of the samples was evaluated by measuring dimensional changes before and after

calcination at a certain temperature.

Structural and morphological characterization The products were characterized by X-ray diffraction

(XRD, Siemens/Bruker D5000) with CuKα radiation. Morphological characterization was performed using

Scanning Electron Microscopy (SEM). Gold coated speciments were examined in a LEICA S440 Scanning

Electron Microscope.

Mechanical testing Compressive strength tests were conducted using a Hounsefield HK20-W equipped

with a 20 kN load cell at a crosshead speed of 1.5 mm/min. The diameter and length of each specimen were

measured using a digital caliper and a minimum number of three to five specimens were used to estimate the

mean compressive strength of each temperature group. Mechanical properties such as the compressive strength

and Young Modulus and Work of fracture were determined and averaged using the experimentally obtained σ-

ε curves.

Estimation of porosity Porosity of the specimens was calculated using the following relationship

between porosity P and mechanical strength S as proposed by Rice [10,11]

0exp( )S S b P= − × , (1)

where So is the ideal mechanical strength when porosity is zero and b is a constant related to the particle

stacking and hence pore shape [8,10]. Ishikawa and Asaoka adapted the above equation for calcium phosphate

cements in Diameter-Tensile-Strength (DTS) experimental configuration [13] as follows:

ln 4.367 0.068DTS

S P= − × , (2)

where SDTS is the measured mechanical strength of the porous material, 4.367 is the natural logarithm of the

estimated ideal diameter tensile strength for zero porosity (SDTSo) and the constant b was calculated equal to

0.068. In the present study, the reported values for b and SDTSo=103 MPa were adopted [11]. Transformation of

compressive strength data to DTS data was performed using the Poisson’s ratio for HAP which is equal to 0.28

as reported by Dale [14].

3 Results and discussion

The preparation of the cement is based on the hydrolysis of α-TCP to calcium deficient hydroxyapatite

(CDHA) according to the reaction 3 4 2 2 9 4 4 5

3Ca (PO ) + H O Ca (HPO )(PO ) OH→ [15]. Figure 1 depicts

XRD spectra for cements before and after calcination. XRD pattern recorded from hardened cements without

heat treatment showed broad peaks indicating formation of low crystallinity hydroxyapatite. The presence of β-

TCP which comes from the raw material was also detected. Contamination of α-TCP with β-TCP is common

when quenching is not fast enough or when the components contain impurities which promote the nucleation

of β-TCP [16].

The structural changes that occur during calcination of cements composed of CDHA were studied for

temperatures between 600 and 1300 °C. The XRD spectra are shown in figure 1 and the crystalline phases

identified are listed in table 1. It can be seen that calcination at temperatures higher than 800 °C resulted in the

conversion of CDHA to pure β-TCP which is stable until 1100-1200 °C while at 1200 °C α-TCP is detected.

Further characterization of the solid phases was performed by FTIR spectroscopy. As seen in figure 2 as the

heating temperature increases the O-H stretching band of CDHA at 3570 cm-1 appeared at 600 °C while hardly

detected at 800 °C and at 1000 °C the above band dissapeard. Those results reflect the convertion of CDHA to

β-TCP with increasing temperature, while according to the higly sensitive FT-IR method traces of CDHA

might remain at 800 °C.

Thermal stability of hydroxyapatite depends strongly on its stoichiometry. A systematic work done by

Raynaud et al. demonstrated that thermal stability of synthetic apatitic powders depends on the Ca/P ratio.

Cryst. Res. Technol. 45, No. 3 (2010) 241

www.crt-journal.org © 2010 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

Powders with ratio lower than the stoichiometric value of 1.667 decomposed at 700 °C into a biphasic mixture

of HAP and β-TCP [17]. In another study Gibson et al. illustrated that CDHA upon heating to 700-800 °C

transforms to β-TCP and further heating at about 1120 °C results in the formation of α-TCP [18].

Scanning electron microscopy was carried out in the uncalcinied (hardened) and calcinided cements.

Characteristic micrographs are shown in figure 3. Images obtained from the uncalcinied cements revealed the

presence of entangled needle-like CDHA crystals within the nanoscale (Fig. 3a and b). For the speciments

heated at 600 °C the microstructure was composed of spherical grains and a considerably lower amount of

needle-shaped crystals. Morover, in some cases the formation of necks between particles was determined

which can be considered as the initial stage of sintering. A porous structure was also observed. The diameters

of pores ranged around 1-2 μm (Fig. 3c, d). A solid structure composed of pores was observed at temperatures

higher than 900 °C. This microstructure was the result of sintering in the material. In a recent study it was

shown that sintering calcium phosphate cements at 1100 °C result in a decrease of nanoporosity and the

formation of even more compact conglomerate structures, which rounded up after sintering [19].

Fig. 1 XRD spectra of hardened bone cement (control)

and after heat treatment at 600, 800, 1000, 1100, 1200

and 1300 °C for 5 h. (Online color at www.crt-

journal.org)

Fig. 2 FTIR spectra of hardened bone cement (control) and

after heat treatment at 600, 800, 1000, 1100 , 1200 and

1300 °C for 5 h. (Online color at www.crt-journal.org)

Table 1 Calcination temperature, composition as determined by XRD, compressive mechanical properties and estimated

porosity of calcium phosphate cements. (*present as impurity in the α-TCP source material; **hardly detected by FTIR)

Calcination

Temperature oC

Phases present by

XRD / FTIR

Compressive Strength

MPa

Young’s Modulus

MPa

Estimated

porosity (%)

Non-calcinied HAP, β-TCP* 19,1±1,3 410,0±39,2 43,49±0,07

600 HAP, β-TCP 19.5±5,5 493,1±37,5 43,19±0,28

800 β-TCP, HAP** 25.1± 5,5 1173,7±518,5 39,48±0,22

1000 β-TCP 30,9±6,2 1547,4±258,7 36,42±0,20

1100 β-TCP 29.9±8,4 1111,8±383,8 36,90±0,28

1200 β-TCP, α-TCP 41,5±8.5 1247,5±468,7 32,08±0,20

1300 β-TCP, α-TCP 65.0±13.5 1534,5±334,0 25,48±0,20

The influence of calcination temperature on the compressive strength is shown in table 1 and figure 4,

respectively. It was found that compressive strength of the non-calcinied cement is at 19.1 MPa and by

increasing the heating temperature at 1300 °C, the compressive strength was raised at 65 MPa whereas

Young’s modulus is at 1534,5 MPa.

Figure 4 shows that the temperature-related increase in mechanical strength is accompanied by a parallel

decrease of the estimated material porosity. Furthermore, experimental measurements of volume shrinkage

showed the same trend with the porosity estimation. It is well known that decrease of porosity leads to increase

242 O. Andriotis et al.: Bioceramics produced from calcium phosphate cements

© 2010 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim www.crt-journal.org

of compressive mechanical strength [20]. Bignon et al. showed that porosity decrement stops after 6 h at

1200 °C when sintering process is complete and the porosity is stabilized at a value of 30 % at about 1200-

1250 °C [12]. Our calculations showed a porosity value of 25,5% after a 12 h calcination at 1300 °C.

Fig. 3 SEM micrographs of hardened cements (a,b), and after a calcination at 600 oC (c,d) and 1300oC (e,f) respectively.

Fig. 4 Compressive strength (◊); Volume Shrinkage (○) and Estimated porosity (▲) as a function of

temperature.

It is interesting to note that in the temperature range between 37 - 800 °C the mechanical strength increases, but

not as dramatically as in the temperature range from 800 - 1300 °C. This could be explained by the

concomitant loss of most of the CDHA needle-shaped crystal structure up to 800 °C, due to conversion of

CDHA to pure β-TCP that occurs probably together with the sintering process. At 1100 °C a stabilization of

the compressive strength is observed which can be explained due to phase transformation of β-TCP to α-TCP.

After the transformation has been completed, a dramatic increase of compressive strength, because of the

sintering process, was observed.

4 Conclusion

Calcium phosphate cements consisted of low crystallinity hydroxyapatite were prepared. XRD measurements

showed that hardened cements are transformed at temperatures higher than 800 °C to β-TCP and β-TCP/α-TCP

ceramics above 1100 °C. Furthermore highly sensitive FTIR analysis revealed traces of CDHA even at 800 °C,

wile those traces are not detectable on XRD diffractograms. Heat treatment resulted in a significance increment

of the compressive strength and stiffness, as a consequence of the formation of a denser material, owing to

sintering process.

Cryst. Res. Technol. 45, No. 3 (2010) 243

www.crt-journal.org © 2010 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

Acknowledgements University of Patras research committee project KARATHEODORIS (C.165) is acknowledged for

financial support and FORTH/ICE-HT for providing SEM and XRD research facilities.

References

[1] Α. Hoffman, Adv. Drug Delivery Rev. 54, 3 (2002).

[2] S. C. Lee, C. T. Wu, S. T. Lee, and P. J. Chen. J. Clin. Neurosci. 16, 56 (2009).

[3] L. L. Hench, J. Biomed. Mater. Res. 41, 511 (1998).

[4] M. C. Chang, C. C. Ko, and W. H. Douglas. Biomaterials 24, 2853 (2003).

[5] R. W. Bucholz, Clin. Orthop. Relat R. 395, 44 (2002).

[6] K.U. Lewandrowski, S. P. Bondre, D. L. Wise, and Debra J. Trantolo, Biomed. Mater. Eng. 13, 115 (2003).

[7] R. Z. Legeros, S. Lin, R. Rohanizadeh, D. Mijares, and J. P. Legeros, J. Mater. Sci. Mater. M. 14, 201 (2003).

[8] D. D. Weiss, M. A. Sachs, C. R. Woodard, and J. Long, Term. Eff. Med. Implants 13, 41 (2003).

[9] J. P. Schmitz, J. O. Hollinger, J. Oral Maxillofac. Surg. 57, 1122 (1999).

[10] R. W. Rice, J. Mater. Sci. 28, 2187 (1993).

[11] D. M. Liu, Ceram. Int. 23, 135 (1997).

[12] A. Bignon, J. Chouteau, J. Chevalier, G. Fantozzi, J. P. Carret, P. Chavassieux, G. Boivin, M. Melin, and D.

Hartmann, J. Mater. Sci. Mater. M. 14, 1089 (2003).

[13] K. Ishikawa and K. Asaoka, J. Biomed. Mater. Res. 29, 1537 (1995).

[14] E. Dale, L. Katz, R. S. Gilmore, and R. Murty, J. Biomed. Mater. Res. 6, 221 (1972).

[15] M. P. Ginebra, F. C. M. Driessens, and J. A. Planell, Biomaterials 25, 3453 (2004).

[16] M. P. Ginebra, E. Fernandez, E. A. P. DeMaeyer, R. M. Verbeeck, M. G. Boltong, J. Ginebra, F. C. Driessens, and

J. A. Planell. J. Dent. Res. 76, 905 (1997).

[17] S. Raynaud, E. Champion, D. Bernache-Assollant, and P. Thomas, Biomaterials 23, 1065 (2002).

[18] I. R. Gibson, I. Rehman, S. M. Best, and W. Bonfield, J. Mater. Sci. Mater. M. 12, 799 (2000).

[19] E. B. Montufar, C. Gil, T. Traykova, M. P. Ginebra, and J. Planell, Key Eng. Mater. 361-363, 323 (2008).

[20] Y. Fukase, E. D. Eanes, S. Takagi, L. C. Chow, and W. E. Brown, J. Dent. Res. 69, 1852 (1990).