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Materials Science and Engineering C 51 (2015) 294–299
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Materials Science and Engineering C
j ourna l homepage: www.e lsev ie r .com/ locate /msec
Structure–property relationships of iron–hydroxyapatite ceramic matrixnanocomposite fabricated using mechanosynthesis method
Jamillah Amer Nordin a, Djoko Hadi Prajitno b, Syafiqah Saidin a, Hadi Nur c,d,⁎, Hendra Hermawan e,⁎a Faculty of Biosciences and Medical Engineering, Universiti Teknologi Malaysia, Johor Bahru 81310, Malaysiab Nuclear Technology Center for Materials and Radiometry, National Nuclear Energy, Bandung 40132, Indonesiac Centre for Sustainable Nanomaterials, Ibnu Sina Institute for Scientific and Industrial Research, Universiti Teknologi Malaysia, Johor Bahru 81310, Malaysiad Department of Physics, Institut Sains dan Teknologi Nasional, Jl. Moh. Kahfi II, Jagakarsa, Jakarta Selatan 12640, Indonesiae Department of Mining, Metallurgical and Materials Engineering & CHU de Québec Research Center, Laval University, Québec City G1V 0A6, Canada
⁎ Corresponding authors.E-mail addresses: [email protected] (H. Nur), hen
(H. Hermawan).URL: http://www.hadinur.com (H. Nur).
http://dx.doi.org/10.1016/j.msec.2015.03.0190928-4931/© 2015 Elsevier B.V. All rights reserved.
a b s t r a c t
a r t i c l e i n f oArticle history:Received 4 September 2014Received in revised form 1 February 2015Accepted 15 March 2015Available online 17 March 2015
Keywords:Bovine hydroxyapatiteIronMillingNanocompositeHardness
Hydroxyapatite (HAp) is an attractive bioceramics due to its similar composition to bonemineral and its ability topromote bone–implant interaction. However, its low strength has limited its application as load bearing im-plants. This paper presented a work focusing on the improvement of HAp mechanical property by synthesizingiron (Fe)-reinforced bovine HAp nanocomposite powders via mechanosynthesis method. The synthesis processwas performed using high energy milling at varied milling time (3, 6, 9, and 12 h). The samples were character-ized by X-ray diffraction (XRD), Fourier transform infrared (FT-IR), and scanning electron microscopy (SEM). Itsmechanical properties were investigated bymicro-Vicker's hardness and compression tests. Results showed thatmilling timedirectly influenced the characteristics of the nanocomposite powders. Amorphous BHApwas formedafter 9 and 12 h milling in the presence of HPO4
2− ions. Continuous milling has improved the crystallinity of Fewithout changing theHAp lattice structure. The nanocomposite powders were found in spherical shape, agglom-erated and dense after longer milling time. The hardness and Young's modulus of the nanocomposites were alsoincreased at 69% and 66%, respectively, as themilling timewas prolonged from 3 to 12 h. Therefore, the improve-ment of the mechanical properties of nanocomposite was attributed to high Fe crystallinity and homogenous,dense structure produced by mechanosynthesis
© 2015 Elsevier B.V. All rights reserved.
1. Introduction
Over the past decades, hydroxyapatite (HAp) has been used inmed-ical implant applicationsmainly as bone substitute and coating onmetalor polymer implants [1,2]. Hydroxyapatite is mainly composed of calci-um (Ca) and phosphorus (P)which can be found in human bone [3,4]. Itis a bioactive material that promotes osseointegration, thus making itfavourable for use on bone–implant interface [5,6]. Ions released duringthe dissolutionofHAp stimulate the process of bone growth and remod-elling [7]. This dissolution is more pronounced with amorphous HApcompared to crystalline HAp as reported by Xue et al. [7]. However,the brittle and low strength properties of HAp have been identified asthe root cause for bone substitute and coating failures [8–10]. Attemptsto overcome this limitation have been done recently, mostly by devel-oping HAp composites. The HAp brittleness is ameliorated by incorpo-rating titanium (Ti), mangan (Mn), and zirconium oxide (ZrO2) into
HAp, where this metallic reinforcement has improved HAp densifica-tion and crystallinity [9,11,12].
Iron (Fe) is another potential element to enhance the strength ofHAp. It has been used as a framework structure for HAp bone fillingdevice [13], and for improvingHAp radiopacity in drug delivery applica-tions [10,14]. Similar to Ti and Mn, Fe is added in HAp matrix as rein-forcement for high compression resistance to inhibit failure over arange of deformation [13]. Previously, Fe–HAp composite was synthe-sized via wet-chemical method [11]. By wet-chemical method, Ajesshet al. [14] found that X-ray diffraction spectra of sintered iron oxide–HAp indicated the presence of intermediate phase which was responsi-ble for the improvement of mechanical and biological properties of theHAp. However, the wet chemical method requires costly chemical re-agents and involves long chemical routes [11]. Therefore, another feasi-ble method to synthesis Fe–HAp composite by offers simpler and moredirect process is required.
Mechanosynthesis is a direct mechanical route method [9], which isbased on mechanical activation effect resulting from collided balls dur-ingmilling step [15]. Awork onCaP–Ti composite reported by Silva et al.[16] showed that increased in milling time has resulted in smaller grainsize composite without reaching its melting point. This method also
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produced nano-size particles which is beneficial in providing large sur-face area for bone substitution and integration [17]. To this end,mechanosynthesis was chosen over wet-chemical method as a top-down synthesis from bulk material to smaller pieces using mechanicalform of energy [17].
In this work, we aimed to synthesize 30wt.% Fe–HAp ceramic nano-composite using mechanosynthesis method at different milling times.Reinforcement of less than 15% Fe in hard matrix will not improve themechanical property of HAp, while the reinforcement of more than50% Fe will change the properties of HAp to metal behaviour [18].Chang et al. [19] proved that 33 wt.% Fe was the optimum proportionfor HAp reinforcement [19]. Therefore, 30 wt.% Fe was chosen and rein-forced into HAp in order to improve its mechanical toughness whilemaintaining the HAp ceramic properties. The synthesized Fe–HApnanocomposites were then subjected to several characterization analy-ses andmechanical tests to investigate the effects of milling time on thechemical functionality, crystallite size, lattice strain, morphology, andhardness.
2. Materials and methods
2.1. Sample preparation
Hydroxyapatite was prepared from bovine bone based on the workby Barakat et al. [6]. Firstly, portion of femur bones were taken fromslaughtered cows, cut into 5 cm length, and cleaned with distilledwater and acetone. The small bone pieces were boiled overnight anddried at 160 °C for 48 h to remove any fatty substances and impurities.Thermal decomposition was then performed on the dried bonesthrough heating at 350 °C for 3 h to ensure that all fatty substanceswere removed. The bones were continuously heated at 750 °C for an-other 4 h to fully transform the bone into HAp. Finally, the chalkywhite HApwas crushed into powdered formbyusingmortar and pestle.By this stage, the powdered formof HAp is known as bovine hydroxyap-atite (BHAp).
The 30 wt.% Fe–BHAp nanocomposite was prepared by adding 1.5 gFe powders (particle size 450 μm, Goodfellow, UK) to 3.5 g BHAp. The
Fig. 1. Schematic view of me
powder mixtures were ball milled by using high energy milling (SPEXSamplePrep 8000M Mixer/Mill, USA) at 3, 6, 9, and 12 h (Fig. 1). Themilling was operated with Teflon vial and zirconia balls. The ball ratioand rotational speed were set at 1:10 and 1200 rpm, respectively. Theball milled nanocomposite powders were then collected and stored ina desiccator for characterization and testing. The density of sampleswas measured by Archimedes method.
2.2. Characterization analyses
Chemical composition and crystalline phases of the nanocompositepowders were identified by an X-ray diffractometer (XRD, X'pert Dif-fractometer, Philips, USA) using CuKα radiation at λ of 0.15406 nm.Data were collected in 2θ range (20°–80°) at 0.05°/min. The diffractionpeak analyses were conducted by using XPowder software (Ver.2004.04.41 PRO, USA). Crystallite size and lattice strain were calculatedbased on Scherer's equations [15].
D ¼ KλFWHMhklð Þ � cosθ
ð1Þ
E2 ¼ FWHMð Þ24 tanθ2� � ð2Þ
where D is the crystallite size; K is the broadening constant (0.9); λ isthewavelength of X-ray beam; FWHM is the full width at halfmaximumfor the diffraction peak under consideration (rad); hkl is the Miller'splane family where (002) and (110) were applied for HAp and Fe, re-spectively; θ is the diffraction angle (°); and E is the lattice strain.
The presence of functional groups in the nanocomposite powderswas investigated by Fourier transform infrared spectroscopy (FTIR,Prestige 21 Shimadzu, Japan). The FTIR spectra were obtained usingKBr disks at 1:100 “sample-to-KBr” ratio. The transmittance spectrosco-py range was set at 4000–400 cm−1 by 16 scans.
Morphology and particle size were then visualized by a field emis-sion scanning electron microscope (FESEM, Zeiss Supra 35 VP,
chanosynthesis method.
Fig. 3. FTIR spectra of milled nanocomposite powders at 3, 6, 9, and 12 h.
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Germany) at 25 k× magnification. The nanocomposite powders werepreviously sputter-coated with ultrathin gold film for 10 min to reduceelectron discharge during the analyses.
2.3. Mechanical testing
The nanocomposite powders were compacted into pellets (∅ =8 mm, thickness = 5 mm) at 129 kg/cm2. The pellets were sintered at900 °C for 2 h under argon atmosphere. The pellet surfaces were flat-tened with abrasive SiC paper grit #1000. Next, hardness measurementwas performed on the flattened pellets by using a micro-Vicker's hard-ness tester (Matsuzawa DVK II, Japan). A continuous load of 1.9 N wasapplied to each pellet for 15 s according to ASTM E384 standard [18].Five indentations were performed on each set of pellets for significancepurposes. The sintered pellets were also subjected to axial compressiontests by using a universal testingmachine (Instron 8874, USA) to furthervalidate themechanical properties data. A crushing load of 2.5 kNwith acrosshead speed of 0.02 mm/s was given to a set of pellets to measurethe Young's modulus value.
3. Results
3.1. XRD analysis
Fig. 2 shows XRD spectra ofmilled nanocomposite powders at 3, 6, 9,and 12 h. The BHAp and Fe peaks were determined based on JCPDSdatasheet: HAp: JCPDS#09-0432 and Fe: JCPDS#6-0696. The XRD spec-tra showed that highly crystalline BHAp powders were obtained after3 h milling. Several new peaks of BHAp were detected at 2θ = 28.05°,48.11°, 50.47°, 51.29°, 52.06°, and 53.19° after 6 h milling. These peaksremained at approximately similar intensities after 9 h milling. Howev-er, continuousmilling that reaches up to 12 h resulted in the disappear-ance of some BHAp peaks at 2θ = 28.05° and 48.11°.
The intensity of Fe peak at 2θ=44.62° showed an increasewhen themilling processes were performed from 3 to 12 h, indicating that the Fepowders becamemore crystallized after being subjected to longer mill-ing time. Furthermore, new Fe peak was observed at 2θ=65.02° whenthe nanocomposites were milled from 6 to 12 h. There were no impuri-ties in the nanocomposites for all milling times as confirmed by the XRDspectra.
3.2. FTIR analysis
Fig. 3 shows FTIR spectra of milled nanocomposite powders at 3,6, 9, and 12 h. Broad peaks between 3800 and 2800 cm−1 were no-ticed on all FTIR spectra which indicated that the vibration of
Fig. 2. XRD patterns of milled nanocomposite powders at 3, 6, 9, and 12 h.
adsorbed water (OH−) in the apatite lattice accords with HAp chemicalformula (Ca10(PO4)6(OH)2). The OH− peaks broadened as the millingtime increased due to the transformation of BHAp from crystalline toamorphous phase.
There are four types of vibration modes representing the phosphate(PO4
3−) functional group, ν1, ν2, ν3, and ν4 of HAp [20]. Those vibrationmodes were found on the FTIR spectra, thus confirming that the BHAppowders were composed of phosphate group. The major peak of phos-phate group (ν3 vibration mode) was between 1047 and 1091 cm−1.The peak at 964 cm−1 represented the ν1 symmetric P–O stretching vi-brationmode. The peaks at 569, 601, and 630 cm−1 corresponded to theν4 vibration while the peak at 473 cm−1 corresponded to weak phos-phate band (ν2 vibration mode) [20,21]. Interestingly, a new peak ofHPO4
2− at 1215 cm−1 was observed on the FTIR spectrum at 9 hmillingdue to the hydration of phosphate ions. This peak becamemore intenseat 12 h milling.
The carbonate (CO32−) functional group was indicated by two types
of vibration modes, ν2, and ν3. The ν2 stretching mode was found at873 cm−1 whereas the ν3 stretching mode was found at 1414 and1460 cm−1, similar to the results by Rehman and Bonfield [20]. The in-tensity of ν3 bandwas found to have increased at 6 hmilling, although itwas further reduced at 9 and 12 hmilling. These results were caused bythe formation of crystal apatite at 6 h [20] and the formation of HPO4
2−
at 9 and 12 h milling.
3.3. SEM analysis
Fig. 4 shows morphology images of milled nanocomposite powdersat 3, 6, 9, and 12 h. The particles were transformed from flakes to almostspherical shapes after being milled for a longer period of time. The par-ticle size also decreased from approximately 240 nm to 97 nm, whichcauses the formation of smaller pores between the nanoparticles. Thenanoparticles were also found to be agglomerated after longer millingprocess due to ball collision [15].
3.4. Mechanical testing
Hardness and Young's modulus value is defined as the ability of thematerial to resist plastic deformation, which implies on the materialstrength and material properties such as yield strength [22]. Fig. 5shows the micro-hardness measurement on the sintered nanocompos-ite pellets where the hardness of thematerial showed an increase as themilling time was increased. The 12 hmilling resulted in nanocompositehardness of 210 Hv which was 38% higher than the nanocompositehardness at 3 h milling (124 Hv). Furthermore, results from the
Fig. 4. SEM images of milled nanocomposite powders at (a) 3, (b) 6, (c) 9, and (d) 12 h.
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compression test demonstrated that the Young's modulus of 30 wt.%Fe–BHAp nanocomposite milled from 3 h (28 GPa) to 12 h (46 GPa)was also increased.
4. Discussion
4.1. Crystallite size and lattice strain
This study confirmed that the mechanical activation at various timeshas produced different crystallographic properties of the nanocompositeswithout being subjected to any melting point temperature. The changes
Fig. 5. Micro-hardness of sintered nanocomposite pellets produced at various millingtimes.
in crystallographic properties were caused by collision between ball–ball and ball–vial. Mechanical impact from the ball onto the nanocompos-ite powders led to the deformation of particle structure and refinement ofparticle size. These refined structures were influenced by crystallite sizeand lattice strain [15]. The crystallite sizewas indicatedby thepeakprofileof Full Width Half Maximum (FWHM) where changes in FWHM mayhave caused micro-strain in the lattice.
The 9 hmilling and 12 hmilling have produced low crystallite peaksand decreased the crystallite size (Fig. 2). Previous worked by Nasiri-Tabrizi et al. [23] showed similar findings where the HAp crystallinityhad increased at 40 h milling and further reduced at 80 h milling. How-ever, their experiments were performed at 650 rpm, which requiredlonger time for crystallinity transformation. Therefore, in our study,the duration of milling ranged from 3 to 12 h due to high energymillingat 1200 rpm.
Fig. 6 shows magnified XRD patterns of BHAp (hkl = 002) and Fe(hkl = 110) peaks. The diffraction of BHAp was highly increased afterthe powders were milled at 6 h compared to 3 h. However, it graduallydecreased at 9 and 12 h (Fig. 6a). The BHAp peaks also became broaderfrom 3 to 12 h, which corresponded to the breadth of FWHM. These re-sults demonstrated the reduction in BHAp crystallite size and crystallin-ity, suggesting the transformation into amorphous phase after longermilling time. The observed amorphous phasemay contribute to the for-mation of nanosize particles [17].
Alternatively, the XRD peaks at 6 to 12 h showed gradual movementtoward lower angles. This phenomenamight be attributed to Fe dopingin the crystal lattice of BHAp. High diffraction of Fe peaks were noticedafter longer milling time, presenting the doping theory. It was reportedby Anee et al. [23] that the presence of doping Fe has caused evidentshifting in the peaks. Furthermore, Wang et al. [18] explained that theshifting was due to smaller radius of Fe(III) than Ca(II) which increasedthe ability of Fe to substitute within HAp lattice. In addition, in this
Fig. 6.Magnified XRD patterns of (a) BHAp (hkl = 002) and (b) Fe (hkl = 110) peaks.
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study, the density of nanocompositeswas increased from 1.40±0.02 to1.54±0.01 g/mm3 formilling at 3 to 12 h, suggestingmore Fe doping atlonger milling time.
Crystallite size and lattice strain were then calculated from the XRDpatterns byusing Scherrer formula as shown in Fig. 7. The crystallite sizeof BHAp (Fig. 7a) decreased from 43.5 nm to 23 nm after continuousmilling from 3 to 12 h, respectively. Meanwhile, the lattice strain ofBHAp (Fig. 7b) increased from 0.15% to 0.34% after the samemilling du-ration. Changes in those parameters have modified the crystallinity de-gree [24], which lead to the transformation of BHAp powders fromcrystalline to amorphous. It is reported that amorphous powder hashigh degradation properties which will improve the bioactivity of thematerial, thus increasing the integration ability with surrounding bone.
Analysis on the crystallite size of Fe showed that the diameter wasincreased from 29.7 nm at 3 hmilling to 34.4 nm at 9 hmilling, then de-creased to 30.4 nm at 12 h milling. In contrast, the lattice strain for Fedecreased from 0.29% at 3 h milling to 0.21% at 9 h milling, then in-creased to 0.27% at 12 h milling. These results were identical to thebroaden XRD peak profile (hkl = 110) which has been discussed previ-ously. Changes in the crystallite size and lattice strain have increased thelattice defect due to the transformation of kinetic energy to thermal en-ergy during the mechanosynthesis process [17]. The lattice defect in-cluding stacking faults, dislocation, and internal stresses [17] wassubsequently caused by micro-strain.
4.2. Chemical functionality of the nanocomposites
The apatite functional groupswere observed in the FTIR spectra at allmilling times (Fig. 3). Eventually, the bands at 1215 and 1226 cm−1
Fig. 7. (a) Crystallite size measurement and (b) lattice strai
during 9 and 12 h milling respectively, were attributed to theautoprotolysis of phosphate ions (PO4
3−). These ions reactedwith atmo-spheric water (H2O) to gain hydrogen ions (H+) ions by producing hy-droxyl ions (OH−) [25] as shown in Eq. (3) below:
PO3−4 þH2O→OH− þHPO2−
4 : ð3Þ
This reaction explained the presence of PO43− and OH− spectra that
indicate the formation of amorphous phase. The amorphization by con-tinuous milling was further confirmed with the appearance of HPO4
2−
ions at 9 and 12 h milling.
4.3. Mechanical properties
The depth indention of Vicker's microhardness and Young's modu-lus data described themechanical properties of the Fe–BHAp nanocom-posites. The hardness of 30 wt.% Fe–BHAp showed an improvementafter the nanocomposites were subjected to longer milling time. LowYoung's modulus and large standard deviation of micro-hardness dataat 3 and 6 h milling (Fig. 5) were attributed to uneven shape (flakeand spherical) of the nanocomposite powders. After longer millingtime (9 and 12 h), the powders were transformed into more homoge-nous spherical shape (Fig. 4). This homogenous spherical shape causedthe powders to be highly dense, compact, and structured, resulting inhigh Young's modulus and consistent Vicker's hardness measurementwith low standard deviation (Fig. 5). Therefore, increased in millingtime has refined the microstructure and formed evenly reinforcementdistribution due to the mechanism of plastic deformation, fracture,and coldwelding [27]. Sintering temperature did not affect the variation
n of milled nanocomposite powders at 3, 6, 9, and 12 h.
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in hardness improvement as constant and similar temperature wasadapted for all samples.
5. Conclusion
Mechanosynthesis method is a feasible and effective process to pro-duce naturally-derived HAp composites with improved mechanicalproperties. The 30 wt.% Fe–BHAp nanocomposite powders were suc-cessfully synthesized via this method at different milling times (3, 6, 9,and 12 h). The crystallite size of Fe and BHApwas reduced after 9 hmill-ing, resulting into amorphous BHAp and crystalline Fe powders. Thehigh crystallinity of Fe powders at longer milling time has enhancedthe hardness and high Young's modulus of the BHAp nanocomposite.
Acknowledgements
The authors would like to acknowledge ZAMALAHUTM scholarship,Tier-1 grant provided by the Malaysian Ministry of Education throughUniversiti Teknologi Malaysia and fonds de demarrage from UniversitéLaval and CHU de Québec Research Center.
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