M A T E R I A L S C H A R A C T E R I Z A T I O N 6 1 ( 2 0 1 0 ) 1 3 6 3 – 1 3 7 0

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Fabrication and characterization of magnesium–fluorapatitenanocomposite for biomedical applications

M. Razavi⁎, M.H. Fathi, M. MeratianBiomaterials Group, Department of Materials Engineering, Isfahan University of Technology, Isfahan, 84156-83111, Iran

A R T I C L E D A T A

⁎ Corresponding author. Tel.: +98 913 110 893E-mail addresses: m.razavi@ma.iut.ac.ir,

1044-5803/$ – see front matter © 2010 Elsevidoi:10.1016/j.matchar.2010.09.008

A B S T R A C T

Article history:Received 17 April 2010Received in revised form7August 2010Accepted 28 September 2010

Recent studies indicate that there is a high demand for magnesium alloys with adjustablecorrosion rates, suitable mechanical properties, and the ability for precipitation of a bone-like apatite layer on the surface of magnesium alloys in the body. An approach to thischallenge might be the application of metal matrix composites based onmagnesium alloys.The aim of this work was to fabricate and characterize a nanocomposite made of AZ91magnesium alloy as the matrix and fluorapatite nano particles as reinforcement. Amagnesium–fluorapatite nanocomposite was made via a blending–pressing–sinteringmethod. Mechanical, metallurgical and in vitro corrosion measurements were performedfor characterization of both the initial materials and the composite structure. The resultsshowed that the addition of fluorapatite nano particle reinforcements to magnesium alloyscan improve the mechanical properties, reduce the corrosion rate, and accelerate theformation of an apatite layer on the surface, which provides improved protection for theAZ91 matrix. It is suggested that the formation of an apatite layer on the surface ofmagnesium alloys can contribute to the improved osteoconductivity of magnesium alloysfor biomedical applications.

© 2010 Elsevier Inc. All rights reserved.

Keywords:X-ray diffractionScanning electron microscopyMagnesium alloyFluorapatite nano particlesNanocomposite material

1. Introduction

Biodegradable metals have been successfully used in biomate-rial science [1,2]. The purposes of biodegradable implants aresupport tissue regeneration and healing by degrading thematerials and concurrently replacing the implants with con-trolled corrosion rate through the surrounding tissues [3–6].

Magnesium has been recently recognized as a biomaterialfor bone substitutes due to its excellent properties such asgood biocompatibility, biodegradability, high strength com-pared to polymers, high ductility compared to bioceramics [4]and relatively low elastic modulus [7–9]. The closer elasticmodulus of magnesium alloys to natural bone compared toother metallic materials could significantly reduce the“stress shielding” existing in the metallic bone implants [5].While fast corrosion kinetics is generally beneficial inbiodegradable implants, the problem for magnesium alloys

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is that their corrosion rate is too high [10], thereby losingtheir mechanical integrity before the tissues have sufficienttime to heal [11], and inhibiting apatite nucleation on thesurface [12]. Therefore, the main research activities arefocused on how to increase the strength of magnesiumalloy and how to protect magnesium from fast corrosion [12].Recent studies indicate that there is a high demand formagnesium alloys with controlled corrosion rate, suitablemechanical properties and the ability for precipitation ofbone-like apatite layer on its surface in the body [7,9]. Anapproach tomeet such requirements could be the applicationof the metal matrix composite (MMC) based on magnesiumalloys [13,14]. The advantage of using nanocomposites asbiomaterials is the adjustable mechanical properties (elasticmodulus and compressive strength), as well as the adjustablecorrosion properties obtained by choosing appropriate com-posite materials [15].

(M. Razavi).

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Calcium is known to reduce the susceptibility of magne-sium to corrosion, even if a negligible amount of it is added tomagnesium [16,17]. As a natural bone composition, hydroxy-apatite (HA: Ca10(PO4)6(OH)2) is known to possess a lowsolubility in body environment [13]. In vitro results haveshown that fluorapatite (FA: Ca10(PO4)6 F2) nano particles couldprovide lower dissolution, and better cell adherence than HA,and significantly improve phosphates activity that lead toimproved osteoconductivity [18,19]. Also, FA nano particlescould provide sufficient low levels of fluoride for improvingbone formation [20,21].

The aim of the present research was fabrication andcharacterization of magnesium–fluorapatite (AZ91-20FA)nanocomposite material for biomedical applications.

2. Materials and Methods

2.1. Materials Production

The magnesium–fluorapatite nanocomposite material wasproduced by mixing AZ91 magnesium alloy powder and 20 %wt of FA nano particles as reinforcement. In order to produceAZ91 magnesium alloy powder, the AZ91 (ASTM B93) billetwith the chemical composition mentioned in Table 1 waschipped via machining process. The chips were crushed topowder by mechanical ball milling under argon atmosphere.The ball milling parameters were selected as follows: ball/powder ratio: 15:1, rotational speed: 250 rpm, and time: 12 h.The FA nano particles were produced by mixing calciumhydroxide (Ca(OH)2), phosphorous pentoxide (P2O5) and calci-um fluoride (CaF2) by planetary high energy ball milling usingzirconia vial and zirconia balls with ball/powder ratio of 35:1,rotational speed of 300 rpm, and the time of 6 h [18]. Thecomposite fabrication was performed through a blending–pressing–sintering method. AZ91 magnesium alloy powderwith 20 %wt of FA nano particles was mixed by ball milling(ball/powder ratio: 5:1, rotational speed: 150 rpm, and time:15 min.). Uniaxial pressing was applied at 880 MPa pressureand the pressed materials were sintered at 400 °C for 1.5 hunder argon atmosphere. The obtained samples were ma-chined in order to prepare disks, 12 mm in diameter and 5 mmin height.

2.2. Mechanical and Metallurgical Testing

Vickers hardness (HV0.1) was measured at five replicates, witha BUHLER microhardness tester (load 500 g, load time 30 s).

A compressive test was performed according to ASTM E9,for AZ91 and AZ91-20FA samples. The compressive yieldstrength at 0.2% offset was determined from the stress–straindiagram.

Table 1 – Chemical composition of initial AZ91magnesium alloy.

Element Al Zn Mn Fe Cu Mg

(Wt %) 8.63 0.59 0.17 <0.05 <0.05 Balance

The bulk density relative to theoretical density (relativedensity) measurements was performed in accordance withArchimedes’ law on the three randomly selected polishedsamples in ethanol.

Phase structure analysis was carried out by X-ray diffrac-tometer (XRD, Philips Xpert) using Ni filtered Cu kα (λ cu kα=0.154186 nm, radiation at 40 kV) over the 2 θ range of 10–90°. Theobtainedexperimental patternswere compared to the standardscompiled by Joint Committee on Diffraction Pattern andStandards (JCDPS). The FA crystallite size was estimated bybroadening XRD peaks using Williamson–Hall formula [22]:

B cos θ = 0:89λ=D + 2ε sin θ ð1Þ

whereλ is thewavelength of the X-ray (nm), B is the fullwidth ofdiffraction peak under consideration (rad) in the middle of itsheight, considered after computer fitting of the X-ray data usingGaussian line shape, and θ is Bragg's angle (º). Thus,whenB cos θwas plotted against sin θ, a straight line was obtained with theslope of 2ε and the intercept as (0.89 λ/D). Thus, the crystallitesize, d (nm), and the lattice strain, ε could be calculated. Thecrystallite size determination of each sample was repeated twotimes for two groups of peaks; one group was (002), (211) and(300), and another was (222), (004) and (213) miller's planesfamily; their average being reported as crystallite size, whiletypical standard deviation was 2 nm.

Scanning electron microscopy (SEM, Philips XL 30: Eindho-ven, The Netherlands) which has been equipped with EDSanalyzer utilizing ZAF corrections and elemental mapping,and transmission electron microscopy (TEM, Philips CM 200FEG: Eindhoven, The Netherlands) operating at acceleratingvoltage of 200 kV, were utilized for microstructure study. Inorder to prepare sample, the cross sections were wet groundwith water down to a grinding size of 2400, polished withalumina solution, and then etched in sequence with HFetchant (10 ml HF (48%), 90 ml water) for 2 s and a solutionwith a composition of 0.6 g picric acid, 10 ml ethanol (96%) and90 ml water for 15 s [23].

2.3. In Vitro Corrosion Measurements

2.3.1. Specimen Preparation for Corrosion TestsThe disk samples 12mm in diameter and 5 mm in height weremachined from the produced rods, and the two flat surfaceswere wet ground with 1200 grid SiC paper and cleanedultrasonically in ethanol for 15min, prior to the corrosion tests.

2.3.2. Electrochemical TestThe electrochemical corrosion test was carried out in a Ringersolution at 37 °C. The experimental set consisted of a conven-tional three-electrode cell containing the working electrode, asaturated calomel electrode (SCE) and a platinum mesh as thecounter. The test was carried out using an Ametek potentiostat(model PARSTAT 2273), and the potentiodynamic polarizationwas run at a scan rate of 0.5 mV/s. The corrosion current wasestimated by the linear fit and Tafel extrapolation method.

2.3.3. Immersion TestImmersion testing of the obtained samples was performed byimmersing them in simulated body fluid (SBF) as this

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procedure has been widely used to prove the similaritybetween in vitro and in vivo behavior of certain compositions.The standard SBF solution was prepared according to Koku-bo's protocol [24]. The inorganic ion concentrations instandard SBF were nearly equal to those in human bloodplasma as shown in Table 2.

Each prepared sample was placed in a sterilized bottlecontaining 20 ml SBF solution. The immersion test wasperformed in a temperature controlled shaking water bathfor different time periods (5, 24, 48, and 72 h) under thephysiological condition of pH 7.41 at 36.5±0.5 °C. Upon thecompletion of the pre-selected immersion times, the sampleswere filtered, rinsed with distilled water and dried in air. Thefunctional groups of precipitated layer on the surface ofsamples in SBF solution were analyzed by Fourier transforminfrared spectroscopy (FTIR, JASCO-680 PLUS) in a mid-IRspectrum range of 400–4000 cm−1. The samples were weighedbefore and after the immersion test by a balance having0.001 g of accuracy to calculate the weight gain (weight afterimmersion − weight before immersion). The immersedsamples were cleaned using 180 g/l chromic acid to removethe surface corrosion product and then rinsed with ethanol,dried in air, and finally weighed to calculate the weight loss(weight before immersion − weight after clean). The corrosionrate was calculated using Eq. (2) [25].

CR = W =At ð2Þwhere CR is the corrosion rate in mg/cm2/h, W is the weightloss in mg, A is the original surface area exposed to thecorrosive media in cm2, and t is the exposure time in hours.

3. Results and Discussion

3.1. Mechanical and Metallurgical Testing

Measurements ofmicrohardness showed that FAnanoparticlesinfluenced the overall hardness of the composite material,compared to standard magnesium alloys in general and couldbe attributed primarily to: (a) the presence of relatively harderceramic particles in thematrix [26], and (b) ahigher constraint tothe localized matrix deformation during indentation due totheir presence [27]. The Vickers hardness formagnesiumalloys,which is reported to have HV0.1 values up to 80 [8], wasmeasured in this study for the AZ91 magnesium alloy and

Table 2 – Nominal ion concentrations of SBF incomparison with those in human blood plasma.

Ion Ion concentrations (Mm)

Blood plasma SBF

Na+ 142.0 142.0K+ 5.0 5.0Mg2+ 1.5 1.5Ca2+ 2.5 2.5CI− 103.0 147.8HCO3− 27.0 4.2HPO4

2− 1.0 1.0SO4

2− 0.5 0.5pH 7.2-7.4 7.40

showedanaverageHV0.1 valueof 80. The sizeanddistributionofFA nano particles can change the overall hardness of thecomposite material compared to standard magnesium alloys[13]. The HV0.1 of the AZ91-20FA is 95±2, depending on thecontribution of FA nano particles to nanocomposite material.Measurements of micro and nano-hardness (Vickers Hardness)have been shown in previous studies to have a strong positiverelationship with yield stress and ultimate stress [28,29]. Theincreased mechanical characteristics are suitable for a bioma-terial for large degradable implants in load bearing applications[13].

Fig. 1 shows the compressive stress–strain curves for AZ91andAZ91-20FAsamples. The compressiveyield strengthat 0.2%offset was determined to be 88±5 MPa for AZ91 and 107±4 MPafor AZ91-20FA. The AZ91-20FA sample seemed to be in higheragreement than AZ91 with the characteristics of the naturalbone, having compressive yield strength of 130–180 MPa [8].Since mechanical stability in body environment and controlledcorrosion rate are two main goals of large degradable implantsin load bearing applications, it is desirable to have a biomaterialwith compressive yield strength similar to natural bone [13]. So,this biomaterial is more suitable for large degradable implantsin load bearing applications compared to a biomaterial withlower compressive yield strength than natural bone.

The relative density of samples based on Archimedes lawwas measured to be 92.7±0.5% for AZ91 and 95.3±0.4% forAZ91-20FA. Thus, an increase in bulk density was observed byadding the FA nano particles as reinforcement. This was dueto filling of porosities in intersections of grains of matrix viareinforcements [13]. It is worth mentioning that the presenceof porosities in the samples with no significant influence onmechanical properties and corrosion rates can help tissuehealing and formation of the bone tissue in the body [11].

Fig. 2 shows X-ray diffraction pattern of the initial FA nanoparticles. The X-ray diffraction pattern showed the peaksexpected for FA. The strong peaks of FA could be seen in therange of 2θ=25–55°. Williamson–Hall method used for calcu-lation gave 28±2 nm for “d” parameter, which confirms thatFA is a nano scale material.

Fig. 3 shows X-ray diffraction patterns of FA, AZ91 andAZ91-20FA samples. The X-ray diffraction pattern of AZ91-20FA exhibits dominant α-Mg peaks. There are also peaks of β-Mg17Al12 phase originating from the AZ91 matrix, and thedefinite peaks of FA.Weak indications of the presence of other

Fig. 1 – Compressive stress-strain curves for AZ91 and AZ91-20FA samples.

Fig. 2 – X-ray diffraction pattern of initial FA nano particles.

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phases have also been found, which corresponds to thepositions expected for calcium hydrogen phosphates orcalcium phosphates.

Fig. 4 shows (a), (b) SEM photomicrographs of AZ91-20FAsample in twomagnifications and (c) TEM photomicrograph ofAZ91-20FA sample. From SEM photomicrograph, it reveals arelatively uniform distribution of the FA nano particles in theAZ91 matrix although some FA agglomerates can be seen inthe matrix. From TEM photomicrograph, it reveals the size ofthe FA nano particles as reinforcement in the AZ91 matrix.

Fig. 5 shows EDS analysis of (a) Mg matrix and (b) FAreinforcements. Since the detection zone of EDS beam is biggerthan the average size of FA, the EDS peaks for FA particles(Fig. 5b) inevitably include compositional information of Mgmatrix near particles, too.

Fig. 6 shows the results of elementalmapping of AZ91-20FAsample by EDS. The map shows homogenous distribution ofelements in the structure. It can be seen that the magnesiumhas almost covered the entire surface of the sample. Otherelements such as calcium and phosphorous, which corre-spond to the dispersed FA in the structure, are partiallyoverlapped. The images of elemental mapping consist of colordots and the dots density depends on the concentration andatomic number of the elements [30]. Since the concentrationof P ion is lower than Ca ion in FA and the P ion is lighter thanCa ion, the dots density of P ion is lower than Ca ion andpartially overlapped. Also, the results of elemental mappingdata can confirm this assumption that there is an increase in

Fig. 4 – (a, b) SEM photomicrographs of AZ91-20FA sample intwomagnifications and (c) TEMphotomicrographofAZ91-20FAsample.

Fig. 3 – X-ray diffraction patterns of FA, AZ91 and AZ91-20FAsamples.

Fig. 6 – Elemental mapping of AZ91-20FA sample by EDS.

Fig. 5 – EDSanalysis of (a)Mgmatrixand (b) FA reinforcements.

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relative density due to filling of porosities at grain boundariesand intersections of grains of the matrix via reinforcements.

3.2. In Vitro Corrosion Measurements

3.2.1. Electrochemical TestFig. 7 illustrates the electrochemical polarization curves of theAZ91 and AZ91-20FA samples in Ringer solution at 37 °C.Table 3 summarizes the electrochemical data obtained fromFig. 7. As shown, a significant change of the corrosion current(Icorr) and corrosion potential (Ecorr) could be observed byadding FA nano particles as reinforcement. The corrosionresistance of AZ91-20FA is higher than AZ91. This observationis confirmed by the values of their corrosion currents andpotentials. The main difference between AZ91 and AZ91-20FAis the presence of the FA nano particles — the phase withhigher corrosion resistance [19] — at the AZ91 magnesiumalloy matrix grain boundaries.

3.2.2. Immersion TestFig. 8 shows SEM photomicrographs of (a) AZ91 and (b) AZ91-20FA samples after 72 h immersion time in the SBF solution. Itcan be seen from Fig. 8a that corrosion has occurred especiallyin grain boundaries of AZ91 sample without formation ofwhite particles on the corroded surfaces and grain boundaries.Fig. 8b shows the formation of white particles of cauliflowershape on the entire surface of AZ91-20FA sample. Thus, theaddition of FA nano particles to AZ91 magnesium alloys asreinforcement can accelerate the formation of white particleson the surface of composite samples.

Fig. 9 shows EDS analysis of the white particles on thesurface of the AZ91-20FA sample after 72 h immersion time inthe SBF solution (Fig. 8b). The results indicated that theprecipitated layer was mainly composed of O, P, Mg, and Ca.

A FTIR spectrum of precipitated layer on the surface ofAZ91-20FA sample after 72 h immersion (Fig. 8b) is shown in

Fig. 7 – Electrochemical polarization curves of the AZ91 andAZ91-20FA samples in Ringer solution at 37 °C.

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Fig. 10. According to Fig. 10, the layer precipitated on thesurface of AZ91-20FA samples after 72 h immersion in SBFsolution contained some CO3

2− groups in PO43− sites. This kind

Fig. 8 – SEM photomicrographs of (a) AZ91 and (b) AZ91-20FAsamples after 72 h immersion in SBF solution.

of phosphate is more similar to biological apatite and could bemore suitable for bone replacement materials [19]. Also, thereis two OH− groups in the FTIR spectrum that confirms thepresence of Mg(OH)2 layer on the surface.

Magnesium is a very active metallic element. Corrosionreactions can happen in the SBF solution at pH 7.4. When theimmersion time is less than 1 h, the electrochemical corrosionmechanism of magnesium in the neutral corrosive mediumsis [31]:

Mg + 2H2O→Mg OHð Þ2 + H2

The metal Mg is transformed into an Mg(OH)2 film. In theSBF solution, the Mg(OH)2 is connected with some H2Omolecule to form the hydrate form of Mg(OH)2.nH2O. Whenthe samples are dried in the air, the film shrinks due todehydration. Then, lots of cracks are formed on the filmsurface. Thus, the surface layer with many cracks (Fig. 8a) ismainly composed of Mg(OH)2.

The SEMphotomicrographs ofAZ91 andAZ91-20FA samples(Fig. 8) clearly confirm that during the immersion, a reactionbetweenmagnesiumalloy and solutionhappens on the surface.As a result, the samples surface is covered by a reaction layer fortheAZ91 andAZ91-20FA samples. The results of EDS (Fig. 9) andFTIR spectrum (Fig. 10) illustrate that bone-like apatite hasprecipitated on the surface of AZ91-20FA samples in the SBFsolution. In vivo studies have already shown that a calciumphosphate reaction layer formsaround themagnesium implant[9]. According to Fig. 8 higher amount of apatite formationaround the AZ91-20FA sample may be due to several reasons.First, thismay be attributed to a reduced rate of corrosion as thecomposite materials decreased the amount of direct contactwith SBF solution, because the reinforcement can provideconditions for precipitation of a passive layer on the surface[13]. In addition, large amounts of magnesium ion releasedduring corrosionof theAZ91 samplepossibly inactivatedapatiteformation [12], resulting in less apatite formation around theAZ91 sample if compared to the AZ91-20FA samples. Thus, thehigh levels of apatite formation on the AZ91-20FA sample maybe explained by the release of low levels of magnesium ion andoperating of FA nano particles as nucleation sites for apatiteformationonthesurfaceofAZ91-20FA,whichhasbeenreportedto enhanceosteoblastic activity and thus generate a stimulatoryeffect on the growth of new bone tissues [5].

Fig. 11 shows (a) weight gain and (b) weight loss ofimmersed AZ91 and AZ91-20FA samples in the SBF solutionas a function of the immersion time. It can be seen from

Fig. 9 – EDS analysis of white particles on the surface of AZ91-20FA sample after 72 h immersion time in SBF solution.

Fig. 10 – FTIR spectrum of precipitated layer on the surface ofAZ91-20FA sample after 72 h immersion in SBF.

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Fig. 10a that about 0.01 and 0.04 g weight gain is observed after5 h immersion for AZ91 and AZ91-20FA samples, respectively,and the weight gain increases with the increase of theimmersion time. After 72 h immersion, about 0.08 and 0.85 gweight is gained for AZ91 and AZ91-20FA samples, respec-tively. Weight gain measurement and the SEM surfacemicrostructure (Fig. 8.) clearly confirm that a reaction betweenthe magnesium alloy and solution happens on the surfaceduring immersion test. As a result, the surface is covered by areactive layer or product. The weight gain of AZ91-20FA ishigher than AZ91. More weight gain of AZ91-20FA can be due

Fig. 11 – (a)Weight gain and (b) weight loss of immersedAZ91and AZ91-20FA samples in SBF solution as a function of theimmersion time.

to more precipitation of apatite on the surface of AZ91-20FA(Fig. 8b), compared to AZ91 (Fig. 8a). After the removal of thecorrosion products with chromic acid, the weight loss(Fig. 11b) indicated an increase with the increase of theimmersion time for all samples. About 0.03 and 0.02 g weightloss is observed after 5 h immersion for AZ91 and AZ91-20FAsamples, respectively, and the weight loss increased with theincrease of the immersion time. After 72 h immersion, about0.055 and 0.036 g weight is lost for AZ91 and AZ91-20FAsamples, respectively. The localized corrosion attack isresponsible for the major weight loss in the immersion test[17]. Theweight loss of AZ91-20FA is less than AZ91, due to theformation of apatite layer on the surface that would operate asa passive layer and reduce the corrosion of the AZ91 matrix[12].

Fig. 12 shows corrosion rate of the immersed AZ91 andAZ91-20FA samples in the SBF solution as a function of theimmersion time. The conversion of the weight loss intocorrosion rates revealed that in AZ91 and AZ91-20FA samples,the corrosion rate decreases with the increase of theimmersion time. The corrosion rates for the AZ91 sampleswere approximately 1.3 mg/cm2/h and 0.25 mg/cm2/h after 5and 72 h of immersion times, respectively. However, for theAZ91-20FA samples, the corrosion rate after 5 and 72 himmersion times was 0.8 and 0.08 mg/cm2/h, respectively.The corrosion rates of all samples decreases with the increaseof the immersion time due to the formation of Mg(OH)2passive layer on the samples surface [31]. In this study, thecorrosion behavior of AZ91-20FA samples was stronglydetermined by the formation of apatite passive layer on thesamples surface, as shown for MMCs in different corrosiveenvironments [13]. While corrosion occurred in SBF solution,an apatite layer was found on the AZ91-20FA samplesconsisting of phosphate and carbonate group (Fig. 10). Theseprecipitations seemed to reduce the local corrosion of theAZ91-20FA samples. In contrast, the uncovered areas of AZ91(the area without any apatite precipitated layer) exhibitedhigher corrosion rates than AZ91-20FA (Fig. 12).

Further studies need to focus on combinations of differentmagnesium alloys as the MMC matrix and on the selection ofbiocompatible nano particles as reinforcement.

Fig. 12 –Variation of the corrosion rate of immersedAZ91 andAZ91-20FA samples in SBF solution as a function of theimmersion time.

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4. Conclusion

A magnesium–fluorapatite nanocomposite material wasmade via a blending–pressing–sintering method. The corro-sion behavior of magnesium alloy in the SBF solution can bedescribed by the dissolving of magnesium through reactionsbetween magnesium and solution, and also by the precipitat-ing of apatite on the magnesium surface. The addition of FAnano particle reinforcements to magnesium alloys canimprove the mechanical properties, reduce corrosion rate,and accelerate the formation of an apatite layer on the surface,which provides improved protection for the AZ91 matrix. It issuggested that the formation of an apatite layer on the surfaceof magnesium alloys can contribute to the improved osteo-conductivity of magnesium alloys for biomedical applications.

Acknowledgement

The authors would like to extend their gratitude for thesupporting provided by Isfahan University of Technology,Iran.

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