Mechanical properties and corrosion behavior of powder metallurgy iron-hydroxyapatite composites for biodegradable implant applications

Materials and Design 109 (2016) 556–569

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Mechanical properties and corrosion behavior of powder metallurgy iron-hydroxyapatite composites for biodegradable implant applications Mahdi Dehestani a,⁎, Erik Adolfsson b, Lia A. Stanciu a a b

Purdue University, Department of Materials Engineering, West Lafayette, IN 47907, USA Swerea IVF AB, Ceramic Materials, 431 53 Mölndal, Sweden

H I G H L I G H T S

G R A P H I C A L

A B S T R A C T

• Hydroxyapatite (HA) powders in three particle size groups (b 1 μm, 1–10 μm, 100–200 μm) were synthesized. • Nine iron-hydroxyapatite composites (HA content = 2.5, 5, 10 wt%) were fabricated via the powder metallurgy process. • Tensile strength and ductility of the composites decreased with increasing HA content and decreasing HA particle size. • In vitro corrosion rates of the composites increased with increasing HA content and decreasing HA particle size.

a r t i c l e

i n f o

Article history: Received 21 April 2016 Received in revised form 16 July 2016 Accepted 18 July 2016 Available online 20 July 2016 Keywords: Biodegradable metal Iron–hydroxyapatite composite Powder metallurgy Particle size Mechanical properties In vitro degradation

a b s t r a c t Nine Fe–HA composites were fabricated via powder metallurgy method by varying the amount (2.5, 5, 10 wt%) and particle size (b 1 μm, 1–10 μm, 100–200 μm) of hydroxyapatite (HA) as a bioactive phase in the iron (Fe) matrix. X-ray diffraction did not detect any phase changes in HA after the sintering process. Uniaxial tensile tests measured the strengths of the composites. Polarization and immersion tests estimated the corrosion rates (CR). Yield strength, tensile strength, and ductility of the composites decreased with increasing HA content and decreasing HA particle size, whereas their corrosion rates increased. The strongest composite was Fe– 2.5 wt% HA (100–200 μm) with σy = 81.7 MPa, σu = 130.1 MPa, fracture strain of 4.87%, and CR = 0.23 mmpy. The weakest composite was Fe–10 wt% HA (b1 μm) which did not exhibit plastic deformation, fractured at σu = 16.1 MPa with 0.11% strain, and showed the highest CR of 1.07 mmpy. This study demonstrates how the relative particle size between Fe and HA determines the mechanical and corrosion properties of Fe– HA composites, thereby aiding in enhancing future resorbable implant designs. The model can also be used when designing other bioactive composites (i.e. Ti–HA, Mg–HA) via powder metallurgy. © 2016 Published by Elsevier Ltd.

⁎ Corresponding author. E-mail addresses: [email protected] (M. Dehestani), [email protected] (E. Adolfsson), [email protected] (L.A. Stanciu).

http://dx.doi.org/10.1016/j.matdes.2016.07.092 0264-1275/© 2016 Published by Elsevier Ltd.

M. Dehestani et al. / Materials and Design 109 (2016) 556–569

1. Introduction The importance and usefulness of developing metallic biodegradable alloys from iron (Fe) and magnesium (Mg) for hard tissue repair and bone reconstruction in orthopedics have been highlighted extensively in the biomaterials community over the past decade [1–3]. The concept of biodegradable metals is “providing a temporary support on healing process of a diseased tissue and progressively degrade thereafter” [3]. There is no need to perform a secondary surgery for implant removal since biodegradable bone fixtures in forms of rods, plates, screws, and anchors are supposed to provide initial mechanical support and dissolve gradually into the physiological environment without inducing toxicity as the new bone tissues replace the implant during the healing process. A progressive drop in mechanical strength and integrity of the implant through degradation also facilitates a gradual load transfer from the implant to the bone, which minimizes the stress shielding effect [4]. Principal requirements in the design of biodegradable metals for orthopedics are not limited to their mechanical stability, degradation characteristics, and cytotoxicity. In order to establish bone ingrowth and promote a stable bond at the interface between bone and implant material, the implant surface needs to be bioactive to assist osteoconductivity through bone cells growth and nucleation of biological apatite [5,6]. Hydroxyapatite (HA) with stoichiometry of Ca10(PO4)6(OH)2 has been widely used in the biomaterials field as both a bioactive bone substitute and a coating material due to its excellent biocompatibility and chemical similarity to mineral phase of the human bone, which is mainly composed of calcium (Ca) and phosphorous (P) [7–12]. HA has been incorporated into titanium (Ti) matrix to develop tailored composites for orthopedic and dental applications [13–21]. HA has also been added to Fe [22] and Mg [23–28] metals to improve the bioactivity of these matrices for biodegradable implant applications. Pure Fe has been proposed for biodegradable cardiovascular stents [29] and hard tissue scaffolds [30,31]. Regardless of their slow degradation rates, iron alloys are attractive options due to their favorable mechanical properties for load-bearing applications [1–3]. Combining the inherent strength and ductility of pure iron matrix with markedly bioactive phases like HA and tricalcium phosphates (TCP) is an appealing approach for the design of composite materials for biodegradable bone fixation devices and scaffold applications. Past researchers have prepared such Fe–bioceramic composites via powder metallurgy route [22,32]. When compared with pure iron, the Fe–HA and Fe–TCP composites studied by Ulum et al. [22] exhibited increased in vitro degradation rates and a higher cellular activity was observed in their cytotoxicity test results. The implantation of these materials in sheep animal models monitored by X-ray radiography showed a consistent healing process of the tissue [22]. In another study [32], a Fe–40 vol% β-TCP composite prepared by powder injection molding showed a 28% higher corrosion rate than pure Fe after 56 days of immersion in 0.9% NaCl solution. The material experienced b1% decrease in compressive yield strength after the immersion period whereas the strength of pure iron dropped by 44%. There is currently a fast growing trend in the use of biodegradable Mg-based materials due to their excellent biocompatibility, high specific strength, safe degradation, and mechanical properties which are close to those of human natural bone [33]. Several studies have adopted the conventional powder metallurgy method to prepare Mg–HA composites [23–28]. Gu et al. [23] prepared Mg–HA composites (10, 20, 30 wt% HA) and demonstrated that yield strength, tensile strength and ductility of the composites decreased with increasing HA content whereas their corrosion rates increased. In a recent study [24], Khalajabadi et al. blended nanoparticles of HA and MgO with Mg matrix and investigated the microstructure, biocorrosion, biocompatibility and physical properties of the Mg/HA/MgO composite system when the amounts of constituents varied. Khalil et al. [27] utilized highfrequency induction heat sintering (HFIHS) and optimized the

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processing parameters to achieve high density nanostructured Mg1 wt% HA composites and investigated their microstructural features and mechanical properties. Witte et al. used AZ91D magnesium alloy as matrix and fabricated an AZ91D-20 wt% HA composite [28]. They improved the mechanical properties of the composite by including a hot extrusion step after sintering. Also, HA agglomeration during processing and its impacts on mechanical and corrosion properties of the Mg matrix were discussed. Despite recent advances in the field of biodegradable metal, no effort has been made to date to gain an understanding of the mechanical properties of Fe–bioceramic composites and the reported strength data have been limited to compressive strength measurements [22,32]. Moreover, a critical characteristic of raw powders like particle size and its impact on final sintered microstructure, strength, and degradation rates of materials have not been studied. In general, the majority of the works published on enhancing the bioactivity of metallic biomaterials by incorporation of HA particles have overlooked the particle size effect. Therefore, this paper aims to present a systematic study that investigates the tensile strength variation of powder metallurgy derived iron-hydroxyapatite (Fe–HA) composites based on their composition and HA particle size. Additionally, in vitro degradation behavior, microstructures and fractography of the test specimens are presented too. 2. Materials and methods 2.1. Powder processing The iron powder used in this study was ABC100.30 (Höganäs AB, Sweden), a water-atomized powder (0.002 wt% carbon and 0.04 wt% oxygen) with approximate particle size in the range of 30–200 μm. Three batches of hydroxyapatite powder having identical chemistry but different particle size distribution below 1 μm (b1 μm), 1–10 μm, and 100–200 μm were prepared. All different particle sizes used were prepared from the same powder of hydroxyapatite (Plasma Biotal, UK). A water-based suspension with high solid loading of the hydroxyapatite powder was prepared. The suspension was sprayed into liquid nitrogen followed by freeze drying to transform the frozen droplets to granules. From these, granules with a size of 125–250 μm were collected by sieving. When the collected granules were sintered at 1250 °C for 2 h, the sintering shrinkage reduced the granule size to around 100–200 μm. The remaining material with other sizes were also sintered at 1250 °C for 2 h and used to prepare the finer particle fractions. When various particle sizes below 30 μm are to be selected, sieving is no longer a suitable method. The finer particles were thus obtained by sedimentation in water. The sedimentation procedure was repeated several times to divide the powder into different fractions with a narrow particle size distribution. To finally remove the water from the particles in the sediment without agglomerates to be formed, the sediment was frozen and freeze dried. From these fractions, one with a particle size of around 10 μm was selected. The finest powder with a particle size below 1 μm was finally prepared by ball milling with zirconia milling media. The milled suspension was frozen in liquid nitrogen without any additives and freeze dried to avoid agglomeration of the fine powder particles. Finally, all powder fractions prepared were reheated to 900 °C for 2 h in air to ensure that the particles surfaces had the same chemical and thermal history, since the particles surfaces may vary depending on how the final surface of the particle was formed which could be as sintered, a fractured surface caused by milling or in contact with water for an extended time during the preparation process. 2.2. Sample fabrication Standard tensile test specimens for powder metallurgy materials were fabricated in this work according to the E8/E8M − 13a standard [34]. To design and make the relevant powder compaction die, ISO 2740 standard [35] was used. The iron (Fe) and hydroxyapatite (HA)

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powders were mixed and weighed to prepare mixtures of Fe–2.5 wt% HA, Fe–5 wt% HA, and Fe–10 wt% HA from three HA particles size groups (b 1 μm, 1–10 μm, 100–200 μm) so that nine composites were developed. Samples of pure iron (100 wt% Fe) were also made as control group. No lubricant or any additives were added to the powder mixes. Zinc stearate (Sigma-Aldrich) was applied by a brush to the die walls and punch surfaces to act as dry lubricant and eliminate the friction during ejection of compacts from the die. The mixed powders were cold compacted uniaxially (vertical direction) under 600 MPa pressure applied by a laboratory press (Central Hydrulics, Model 96188). The green compacts were sintered in a tube furnace under N2–5% H2 atmosphere at 1120 °C (heating rate 10°/min) for 1 h after which the samples were left in the furnace to cool down to room temperature. Similarly, some extra tensile bars of pure iron and Fe–2.5 wt% HA composites were sintered in Ar–5% H2 atmosphere in the same furnace. The small pellet specimens prepared for microscopy and corrosion experiments in this study had the same processing history as described.

corrosion rate (CR) were estimated by Tafel fit method via CorrView 3.10 software. 2.5.2. Static immersion tests Small tablet-like specimens (d = 12.7 mm, t = 2 mm) were ground, polished, cleaned (described in Section 2.3) and weighed before exposure to the corrosion media. Three samples were tested for each material. The samples were immersed in Kokubo's solution for 504 h (21 days) at T = 37 °C and pH = 7.40. The pH of the testing solution was checked every 3 days during the test. After the test, the specimens were cleaned according to the ASTM G1-03 standard [38] to remove the corrosion products. The cleaning solution was made by dissolving 3.5 g hexamethylene tetramine in 500 ml hydrochloric acid (HCl, sp gr 1.19) and adding water up to 1000 ml. The samples were cleaned ultrasonically in the solution for 5 min at room temperature, rinsed with deionized water and finally dried and stored in a desiccator for 24 h before weighing. The corrosion rate was calculated according to the following equation [39]:

2.3. Microstructure and characterization Pellet samples were ground with SiC abrasive papers up to 1200 grit, finely polished with colloidal silica, cleaned in ethanol ultrasonically and dried in air. A nital 4% solution was used as etchant. Optical microscopy on samples was performed using an optical microscope (Olympus GX51). A field emission scanning electron microscope (FE-SEM, Philips XL 40) equipped with an energy dispersive X-ray spectrometer (EDS, EDAX, USA) was used to investigate the morphologies of powder particles, elemental compositions of the samples, and fractography of the fractured tensile specimens. X-ray diffractometry (XRD, Bruker D8) using CuKα (40 kV, 40 mA) was conducted for 2θ angles between 20°–60° at scan speed of 1°/min for phase identification. The XRD was performed on the three groups of hydroxyapatite raw powders and also on three separate sintered pellets of Fe–10 wt% HA composites in which the powders were incorporated. Density of the sintered samples and the total (overall) porosity were measured according to the Archimedes' principle as described in the ASTM B962-13 standard [36]. The theoretical densities of the composites were calculated based on the rule of mixture and by taking the theoretical density of hydroxyapatite as 3.156 g·cm−3 and iron as 7.874 g·cm−3. 2.4. Mechanical testing The yield and tensile strengths of the samples were measured using a universal test machine (MTS Insight, 100 kN load cell). A uniaxial tensile test was run at a constant crosshead speed of 1 mm/min. The 0.2% strain offset method was applied to determine the yield point. Eight samples for each composition were tested. 2.5. Corrosion experiments 2.5.1. Electrochemical corrosion tests Samples for electrochemical corrosion tests (10 × 10 × 2 mm3) were embedded in mounting epoxy, ground, polished, and cleaned as described in Section 2.3. The electrochemical measurements were performed in a three electrode cell using an electrochemical analyzer (SP–150 EC-LAB, BioLogic Science Instruments, France). Platinum served as the counter electrode and Ag/AgCl (Sat. KCl) as the reference electrode. Pure Fe and Fe–HA composites served as the working electrode with exposed surface area of 1 cm2. Kokubo's solution buffered with Tris-HCl was prepared as described in reference [37] and used as the electrolyte. Each sample was exposed to 100 ml of the fresh electrolyte. The potentiodynamic polarization tests were carried out at T = 37 °C and pH = 7.40 in the potential range of − 1.50 V to 0.00 V at a scanning rate of 0.166 mV/s. The dwell time prior to the commencement of scans was 30 min. The obtained polarization curves were analyzed and the corrosion potential, corrosion current density, and

CR ¼

KW ATD

where CR is the corrosion rate in millimeters per year (mmpy), K is a constant equal to 87600, W is weight loss (g), A is area (cm2), T is time of exposure (h), D is density (g·cm−3). 3. Results and discussion 3.1. Microstructure and characterization The morphologies and qualitative particle sizes of the Fe and HA powders used in this study are shown in the SEM pictures in Fig. 1. The water-atomized Fe powder particles with approximate particle size in the range of 30–200 μm had irregular external shapes and were internally compact (no porosity). HA powders in this study were divided into three size groups: (a) ultrafine with particle size below 1 μm, (b) fine with particle size between 1 and 10 μm, and (c) coarse with particle size between 100 and 200 μm. The EDS spectra in Fig. 1 and quantitative results in Table 1, which lists the Ca/P ratio in HA powders, demonstrate similar elemental compositions. The XRD patterns of the raw hydroxyapatite powders and the sintered samples of Fe–10 wt% HA composites made from the same powders are presented in Fig. 2. Typical peaks of hydroxyapatite were detected for the powders and the bulk samples. These were assigned as hydroxyapatite on the basis of Powder Diffraction File (PDF) and related reference number. Similarity of peaks and patterns for the raw powders and the sintered compacts implies no major phase change or decomposition associated with hydroxyapatite as a result of the sintering process and indicates proper incorporation of HA into the Fe matrix. Addition of different amounts (2.5 wt%, 5 wt%, 10 wt%) of HA with different particle size (b1 μm, 1–10 μm, 100–200 μm) into the Fe matrix yield nine composites of Fe–HA. The set of SEM micrographs in Fig. 3 shows the microstructure and surface of composites after being sintered, polished, and etched in nital 4% for 30 s. The microstructures revealed different distribution patterns of HA phase between Fe particles when HA amount and size varied. As it can be seen through the micrographs in Fig. 3g–i, the coarse HA particles (100–200 μm) appeared as distinct spheres and the Fe matrix developed a more integrated and uniform structure with these particles. The fine regime HA particles on the other hand, formed a network-like structure in between Fe particles (Fig. 3a–f). This is due to the fact that, the fine HA particles are difficult to be distributed evenly in the Fe matrix by dry mixing of the powders and they tend to form agglomerates which reduces the expected difference in the distribution patterns and morphologies of b1 μm HA and 1–10 μm HA in the Fe matrix.

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Fig. 1. Morphologies of the iron (Fe) and hydroxyapatite (HA) powders and their EDS spectra, (a) pure Fe, (b) HA b 1 μm, (c) HA 1–10 μm, (d) HA 100–200 μm.

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Table 1 EDS quantitative results (point analysis) obtained from the surface of hydroxyapatite powder particles. Hydroxyapatite powder

CaK (wt%)

CaK (at%)

PK (wt%)

PK (at%)

Ca/P atomic ratio

(b1 μm) (1–10 μm) (100–200 μm)

70.29 68.95 70.64

64.64 63.18 65.03

29.71 31.05 29.36

35.36 36.82 34.97

1.82 1.71 1.85

Table 2 lists the measured density and total (overall) porosity for the sintered materials. Pure Fe and the composites in b 1 μm and 1–10 μm groups achieved approximately 95% of full density. The composites in 100–200 μm group gained about 92–95% of full density depending on their HA contents, and they had relatively higher amount of porosity (5–7.5%) compared to the other groups. This is because the large HA particles increase the interparticle spaces between Fe particles which in turn creates some micro-voids in the structure. On the other hand, the fine HA particles fit in and fill the gaps between the Fe particles more efficiently due to their substantially smaller sizes. As a general trend, the amount of porosity increased slightly as the HA content increased from 2.5 wt% to 10 wt% in each group of the composites. 3.2. Mechanical properties Standard tensile test specimens for powder metallurgy materials were fabricated in this work as presented in Fig. 4. For each composition, eight samples were tested. Engineering stress-strain curves of Fe–HA composites and pure Fe are plotted in Fig. 5. The average values of PDF 00-001-1252 Fe PDF 01-074-9771 Ca4.704(H0.492(PO4)3)(OH)0.9

yield strength (0.2% strain offset), tensile strength, and fracture strain are listed in Table 3. The strength data are plotted as bar charts in Fig. 7 and Fig. 8. As a general trend, it can be clearly seen that both yield and tensile strengths of materials dropped with different magnitudes when the amount of HA increased and the HA particle size decreased. According to the strength values, adding 2.5 wt% HA to Fe matrix dropped the yield strength of Fe (96.1 MPa) by approximately 15%, 22%, and 29% for composites in which HA with particle size of 100–200 μm, 1–10 μm, and b 1 μm were used, respectively. In the same manner, the order in yield drop was about 35%, 41%, and 47% when 5 wt% HA was added. There was a substantial decrease in yield strength when 10 wt% HA was added. At this level of HA addition, while 100–200 μm group showed a yield point of 31.8 MPa which equals to 67% decrease in yield of pure Fe, the materials made with fine HA powders did not show plastic region in their stress-strain curves and fractured in a brittle manner. The stress-strain curves for these materials are plotted in a separate graph in Fig. 6. The declining trend for tensile strength was more pronounced than yield strength. For composites with 2.5 wt% HA, tensile strength declined 40%, 45%, and 54% below the Fe tensile strength (214.6 MPa) when HA particle size changed through 100–200 μm, 1–10 μm, and b1 μm groups respectively. Similarly, for composites containing 5 wt% HA the decreasing order was as 61% b 67% b 72% and for composites with 10 wt% HA as 84% b 89% b 93% when HA particle size decreased. All in all, it can be observed that varying the amount of HA as secondary phase has a more dramatic impact on strength and ductility than varying HA particle size. An interesting observation in the yield behaviors of pure Fe and Fe– HA composites was the occurrence of an upper yield point in the stress-

Iron Hydroxyapatite

Fe-10 wt% HA (100-200 µm) Composite

Intensity, arb. units

Fe-10 wt% HA (1-10 µm) Composite

Fe-10 wt% HA (< 1 µm) Composite

HA Powder (100-200 µm)

HA Powder (1-10 µm)

HA Powder (< 1 µm) 20

25

30

35

40

45

50

55

60

2θ/degree Fig. 2. X-ray diffraction patterns of the raw hydroxyapatite (HA) powders and the sintered Fe-10 wt% HA composites.

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Fig. 3. SEM images showing the hydroxyapatite (HA) phase dispersed in the iron (Fe) matrix. The surface was polished and etched in a nital 4% solution for 30 s. Fe-HA (b1 μm) 2.5 wt% (a), 5 wt% (b), 10 wt% (c) Fe-HA (1–10 μm) 2.5 wt% (d), 5 wt% (e), 10 wt% (f) Fe-HA (100–200 μm) 2.5 wt% (g), 5 wt% (h), 10 wt% (i).

strain curves. Although the yield point phenomenon is usually considered as an exception, it can occur in variety of materials [40]. The yield point effect observed in mild steel and soft iron [41] is associated with the presence of small amounts of interstitial solute atoms (e.g. carbon and nitrogen) that segregate around stationary dislocations and create Cottrell atmospheres. Dislocations in this way are locked in position and to free them from the solute atmosphere, a large stress is required to create new mobile dislocations for the yielding to occur [42]. Cottrell demonstrated by theory and experiment that carbon concentrations as low as 0.003 wt% in the iron can lead to yield point phenomenon [41]. In another study, Wain observed yield point effect in zinc crystals

containing 0.0022 wt% nitrogen [43]. The Fe powder in our study initially contained 0.002 wt% carbon and 0.04 wt% oxygen and since the materials were exposed to nitrogen atmosphere during sintering, two additional experiments were performed to further investigate the upper yield points in the samples.

Table 2 Density and total porosity of pure Fe and Fe–HA composites. Total Relative porosity densitya (%) (%)

Material

Sintered density (g·cm−3)

Pure iron

7.49 ± 0.05 95.1 ± 0.7

4.9 ± 0.7

2.5% 7.19 ± 0.03 94.7 ± 0.3 Iron + hydroxyapatite (100–200 μm) 5% 6.88 ± 0.05 93.9 ± 0.7 10% 6.35 ± 0.03 92.7 ± 0.5

5.3 ± 0.3 6.1 ± 0.7 7.3 ± 0.5

Iron + hydroxyapatite (1–10 μm)

2.5% 7.31 ± 0.04 96.1 ± 0.8 5% 7.00 ± 0.06 95.6 ± 0.8 10% 6.51 ± 0.10 95.1 ± 1.5

3.9 ± 0.8 4.4 ± 0.8 4.9 ± 1.5

Iron + hydroxyapatite (b1 μm)

2.5% 7.27 ± 0.02 95.8 ± 0.3 5% 7.10 ± 0.03 95.5 ± 1.3 10% 6.50 ± 0.13 95.0 ± 1.9

4.2 ± 0.3 4.5 ± 1.3 5.0 ± 1.9

a

Relative density = sintered density/theoretical density.

Fig. 4. Standard tensile test specimens of pure Fe and Fe–HA composites.

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240 (a) Pure Fe a (b) Fe-2.5% HA (100-200 µm)

200

(c) Fe-2.5% HA (1-10 µm)

Stress (MPa)

160

(d) Fe-2.5% HA ( < 1 µm) b

120

(e) Fe-5% HA (100-200 µm)

c

(f)

Fe-5% HA (1-10 µm)

d e

80

(g) Fe-5% HA ( < 1 µm)

f

(h) Fe-10% HA (100-200 µm)

g

40

h i j

(i)

Fe-10% HA (1-10 µm)

(j)

Fe-10% HA ( < 1 µm)

0 0

2

4

6

8

10

12

14

Strain (%) Fig. 5. Engineering stress-strain curves of pure Fe and Fe-HA composites sintered in N2–5% H2 atmosphere.

In the first experiment, some tensile bars of Fe and Fe–2.5 wt% HA composites were sintered in Ar–5% H2 atmosphere to exclude the possible influence of nitrogen. Fig. 9 shows the stress-strain curves of these materials in which the upper yield points are still present. This result implies that there were originally solute impurities in the Fe powder at some levels that were capable of creating Cottrell atmospheres in the material. Nevertheless, the influence of nitrogen on yield point phenomenon cannot be ignored since the samples sintered in N2–5% H2 atmosphere showed more pronounced and distinct upper yield points (Fig. 5) compared to those sintered in Ar–5% H2. It has been observed that heating molybdenum, cadmium, and zinc in nitrogen atmosphere has induced yield point in these metals [40]. According to the data in Table 3, there is a significant difference between the mechanical strength of the samples sintered in Ar–5% H2 versus the ones sintered N2–5% H2 which points out the influence of absorbed nitrogen into the Fe matrix from the sintering atmosphere. In the second experiment, the Fe and Fe–2.5 wt% HA composites sintered in N2–5% H2 which had exhibited the yield point phenomena were selected for tensile testing. As illustrated in Fig. 10, the samples were first loaded up to some strains in the plastic region in order to pass the yield point completely, then unloaded and retested without appreciable delay. Upon reloading, the yield point phenomenon did not occur. This is because after the first time of loading, the sample in overstrained condition contains free

(unlocked) dislocations, thus, when reloaded it shows no yield point effect [40,41]. Similarly, Cottrell eliminated the yield point effect in Armco iron by an immediate reload [44]. Based on the results of the two experiments described here, the yield point phenomenon is believed to be due to the presence of solute impurities in the Fe powder which can create Cottrell atmospheres. To date, few research works have combined iron and bioactive ceramics to demonstrate potential properties of these materials for biodegradable bone fixtures and scaffolds. Ulum et al. produced dense powder metallurgy composites by addition of 5 wt% second phase as hydroxyapatite (HA) and tricalcium phosphate (TCP) into Fe matrix and investigated degradation rates, cytotoxicity, and compressive strength of these materials. In a decreasing order of pure-Fe N Fe– HA N Fe–TCP, they reported the mean values for yield as 354 N 325 N 312 MPa and for compressive strength as 752 N 717 N 708 MPa [22]. Reindl et al. prepared dense and porous composites of Fe–TCP with 70–30, 60–40, and 50–50 compositions via powder injection molding technique and studied the degradation rates and compressive strength of materials over a long-term in vitro immersion time [32]. Currently, there is limited literature on the tensile strength of iron-bioceramic composites among the biodegradable metals under development. Bone fracture fixation devices are desired to have appropriate mechanical strength matching that of bone [33]. When it comes

Table 3 Mechanical properties of pure Fe and Fe–HA composites. Material

Yield strength σY (MPa)

Tensile strength σU (MPa)

Fracture strain (%)

Pure iron

96.1 ± 8.8

214.9 ± 5.9

13.32 ± 0.68

Iron + hydroxyapatite (100–200 μm)

2.5% 5% 10%

81.7 ± 10.2 62.1 ± 6.7 31.8 ± 3.4

130.1 ± 10.4 82.2 ± 12.1 33.9 ± 4.3

4.87 ± 1.28 2.33 ± 0.97 0.49 ± 0.17

Iron + hydroxyapatite (1–10 μm)

2.5% 5% 10%

74.9 ± 6.7 56.3 ± 7.7 –

116.8 ± 7.4 68.9 ± 6.3 21.7 ± 9.8

4.97 ± 0.63 1.99 ± 0.37 0.12 ± 0.04

Iron + hydroxyapatite (b1 μm)

2.5% 5% 10%

68.1 ± 6.7 50.8 ± 4.3 –

97.4 ± 7.6 59.6 ± 2.7 16.1 ± 1.9

3.52 ± 0.52 1.96 ± 0.46 0.11 ± 0.03

63.8 ± 3.9

154.7 ± 4.7

13.8 ± 1.4

53.3 ± 1.4 47.1 ± 1.3 43.7 ± 1.1

94.1 ± 6.6 83.9 ± 4.6 70.5 ± 3.7

4.4 ± 0.8 3.9 ± 0.3 3.3 ± 0.1

Pure irona

Iron + 2.5% hydroxyapatitea a

Sintered in Ar-5% H2 atmosphere.

100–200 μm 1–10 μm b1 μm

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Tensile Test (Fe-10 wt% HA) (h) 100-200 µm

(i) 1-10 µm

(j) < 1 µm Fe-2.5% HA

40 Fe-5% HA

h

Fe-10% HA

Stress (MPa)

30

i 20

j Fig. 8. Variation of tensile strength (σu) in the materials.

10

0 0

0 .2

0. 4

0 .6

0. 8

Strain (%) Fig. 6. Engineering stress-strain curves of Fe–10 wt% HA composites sintered in N2–5% H2 atmosphere.

to design of materials for bone repair, it should be considered that tensile strength is as equally important as compressive strength considering the anisotropic nature of bone and the complex loading conditions to which it is subjected. For instance, during normal loading, femur bones experience bending moments which create both tensile and compressive stresses in different regions of the bone [45–48]. There is a large variation in reported values of both the tensile and compressive strength of bone in the literature, mainly because mechanical properties of bone tissue depend on several factors such as: where the bone is taken from (i.e. fibular bone, tibial bone, femoral bone, etc.), amount of mineral phase, water content, age, and gender of person [49]. These details are usually ignored in many papers when strength values of synthesized materials are compared to bone strength values. In order to establish a baseline, in this article we compared the strength data of our materials with strength of human femoral cortical bone obtained from the reference [50]. The main function of cortical bone or sometimes referred to as compact bone is structural support for body weight and movement. The femur or thigh bone as the longest and strongest bone in the skeleton receives extreme loads during daily activities [48]. Table 4 lists the average values of yield strength, tensile strengths, and fracture strains of human femur obtained from 23 year, 31 year, and 63 year old male adults [50]. Using this data to evaluate the strength of dense Fe–HA composites we processed in this study, it can be

Fe-2.5% HA Fe-5% HA

Fe-10% HA

Fig. 7. Variation of yield strength (σy) in the materials.

observed that yield strengths of all materials including pure Fe are lower than yield strength of human femur bone. When it comes to tensile strength, it is only the Fe–2.5 wt% HA group that contains composites with strength levels relatively close to that of cortical bone. In particular, Fe–2.5 wt% HA (100–200 μm) with 130.1 MPa strength stands close to 135 MPa value, the cortical bone tensile strength. Qualitatively speaking, human cortical bone exhibits linear plastic behavior in tension and after showing a distinct yield point, it fractures at somewhat small strains [51]. This fact is reflected in the small difference between values of yield and tensile strength of femur in Table 4 and the relatively small fracture strain of 3.1%. Longitudinal ultimate tensile strains of cortical bone can reach up to 5% in young adults and fall to about 1% in elderly individuals [52]. Fe–HA composites in this study showed fracture strains in a range from 0.11% for the most brittle material to 4.87% for the most ductile material as presented in Table 3. Fig. 5 finely shows how the plastic region in stress-strain curves of materials shrinks and ductility falls as the amount of HA is increased and the HA particle size is decreased. 3.3. Fractography Fractography of fractured tensile test specimens shed light on interparticle neck development during sintering process and provided visual experimental evidence on how HA particle size can influence the strength and ductility of Fe–HA composites. The strongest composite, Fe–2.5 wt% HA (100–200 μm), and the weakest one, Fe–10 wt% HA (b1 μm), were selected for fractography. The SEM micrographs Fig. 11(a–d) are related to the first sample and Fig. 11(e–h) to the latter sample. In Fig. 11(a) and Fig. 11(b), coarse HA particles with approximate size of 100 μm surrounded by Fe matrix can be viewed. Close examination of neighboring Fe matrix revealed large, deep dimples and cleavage-like characteristics on the surface of Fe particles in Fig. 11(c) and Fig. 11(d), which are signs of plastic deformation under tensile mode and indications of strong interface bonding between Fe particles developed during sintering and densification. Networks of equiaxed dimples were present all over the fracture surface and could be easily observed at low magnification. Looking at the interface decohesion between Fe matrix and the HA coarse particle in Fig. 11(b), it is speculated that the boundary regions between Fe matrix–HA particles could be origins of microcrack nucleation during tensile loading and serve as major failure sites within Fe matrix. HA particles in this manner embrittled the Fe matrix such that the fracture strain of pure Fe from 13% dropped to 4.9% in the Fe–2.5 wt% HA (100–200 μm) composite. Fractography of Fe–10 wt% HA (b1 μm) in Fig. 11(e–h) showed the surface of Fe particles almost entirely covered by fine HA particles below 1 μm in size. Owing to their high surface area, fine HA powder particles have shielded the surface of Fe particles and dramatically

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160 a

Stress (MPa)

120 (a) Pure Fe b

(b) Fe-2.5% HA (100-200 µm)

c

80

(c) Fe-2.5% HA (1-10 µm)

d

(d) Fe-2.5% HA ( < 1 µm)

40

0 0

2

4

6

8

10

12

14

Strain (%) Fig. 9. Engineering stress-strain curves of pure Fe and Fe–2.5 wt% HA composites sintered in Ar–5% H2 atmosphere.

reduced their interparticle contact areas for neck development during sintering. Isolated small-and-shallow dimples representing weak interparticle bonding were observed locally and infrequently on the fractured surface. Poor strength and low ductility of Fe–10 wt% HA (b1 μm) composite is reflected in its stress-strain curve in Fig. 6. The base strength of this material is expected to stem partly from limited sintered Fe particles attachments and partly from partial sintering of HA particles. The sintering condition used in this work (1 h, 1120 °C) is suitable for iron powder metallurgy, but not adequate for full densification of HA. Hydroxyapatite is conventionally sintered at 1250 °C for 2 h in air [53,54]. Observations in fractography of Fe–HA composites demonstrated that smaller HA particle size corresponds to higher surface area of HA phase in Fe matrix and this correlation determined the extent of which Fe powder particles could make interparticle area contacts and develop them during sintering. This factor seems to play a more important role than porosity in determining the strength since the composites in 100–200 μm group contained higher porosity but still showed higher strength levels compared to the other groups. In short, the lower the HA content and the larger the HA particle size, the stronger the composites become, and vice versa. This explanation matches the material strengths variation trend illustrated in Fig. 5. However, it should be noted that, in addition to HA amount and HA particle size, the dispersion pattern and the morphology of HA phase in the Fe matrix are some

other factors that can affect the strengths of composites. Ideally, if the fine HA particles do not form agglomerates as a consequence of dry mixing of powders, HA would form a thin continuous phase between Fe particles that will contribute to an increased brittleness and reduced strength, further beyond what was measured for the materials in this study. Also, it should be emphasized that if the particle size of the Fe is changed, then the HA particle sizes used in this study may no longer give the same results. In that case, the relative particle size between the two phases (Fe and HA) will determine how the properties of the composite will be. For instance, if the Fe and HA powder particles are selected in the same size range, this will most likely form a typical composite material in which the two phases are evenly distributed according to the volume fraction of each phase.

3.4. Corrosion behavior Electrochemical tests and static immersion tests were performed at T = 37 °C and pH = 7.4 in Kokubo's solution as electrolyte. The ionic composition of Kokubo's solution is provided in Table 5. Fig. 12 displays potentiodynamic polarization curves of Fe–HA composites with pure Fe as the control. The average values of corrosion potential, corrosion current density, and corrosion rate (CR) are summarized in Table 6. The corrosion rates of the materials obtained from polarization and

140

120

Stress (MPa)

100 (a) Fe

80

(b) Fe–2.5% HA (100-200 µm)

(a) (b) 60

(c) Fe–2.5% HA (1-10 µm)

(c)

(d) Fe–2.5% HA ( < 1 µm) Reload Curve

(d)

40

20

0 0

1

0

1 0 Strain (%)

1

0

1

Fig. 10. Yield point phenomena observed in the engineering stress-strain curves of Fe and Fe–2.5 wt% HA composites sintered in N2–5% H2 atmosphere. Upon reloading the same specimens, the yield point phenomena were not observed. Dashed line represents the stress-strain curve for the reload test.

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Fig. 11. Fractography of the surface of fractured tensile test specimens, Fe–2.5 wt% HA (100–200 μm) – (a-d), Fe–10 wt% HA (b1 μm) – (e-h), the arrows indicate HA particles and signs of plastic deformation under tensile mode (i.e. dimples) on the surface of Fe particles.

immersion tests are plotted in Fig. 13 and Fig. 14 respectively, and indicate the same trends. In general, the corrosion rates of all Fe–HA composites were higher than pure iron. Fe–2.5 wt% HA (100–200 μm) composite showed the lowest CR = 0.23 mmpy and Fe–10 wt% HA (b1 μm) had the highest CR = 1.07 mmpy among the materials. As it can be seen clearly in Fig. 12, the slopes of anodic curves became sharper and the polarization curves shifted toward higher current densities; reflecting the higher dissolution of Fe matrix as a consequence of increasing HA content and reducing HA particle size. The corrosion potentials of all composites stayed relatively close to the Fe corrosion potential, except for the Fe–

10 wt% HA (b 1 μm) and Fe–10 wt% HA (1–10 μm) materials that showed lower corrosion potentials compared to pure Fe. The increase in corrosion rate of Fe because of addition of HA observed in this study is similar to Ulum et al. results [22] where in vitro degradation tests in Kokubo's solution showed that the corrosion rates of Fe–5 wt% HA and Fe–5 wt% TCP were higher than pure Fe. Some potential reasons for the increased corrosion rates observed among the materials are provided here. When the influence of HA particle size on corrosion rate is considered, based on the corrosion rates listed in Table 6, it can be seen that the composites in b1 μm and 1–10 μm groups showed considerably higher corrosion rates than

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Fe-10% HA

Fe-5% HA

Fe-2.5% HA

Fig. 13. Estimated corrosion rates for pure Fe and Fe–HA composites obtained from potentiodynamic polarization tests.

the shrinkage of HA during sintering even though HA is not densified completely. While the shrinkage of metallic Fe matrix is negligible, HA agglomerates may shrink away from the surrounding Fe matrix and leave some interfacial defects behind. The pores and micro crevices at the Fe–HA interface can become channels for further penetration of corrosive electrolyte into the Fe matrix and increase the chance of localized and crevice corrosion besides the commonly occurring uniform corrosion. The influence of HA agglomerates on corrosion of other metallic substrates has been mentioned in the literature. In a case study in which the corrosion behavior of powder metallurgy Mg–HA–MgO composites were evaluated in simulated body fluids [24,25], it was discussed that the local agglomerates of HA and their adjacent pores can cause enhanced dissolution of Mg metal matrix. When coarse HA particles were used (100–200 μm), a more uniform distribution of HA phase could be achieved and the Fe–HA interface as shown in Fig. 3(g–i), exhibited more integrity, and lesser and smaller

Fe-10% HA

Fe-5% HA

Fe-2.5% HA

Fig. 12. Potentiodynamic polarization curves of the Fe–HA composites immersed in Kokubo's solution with pure Fe as control sample.

their counterparts in 100–200 μm group. The difference in corrosion rates can be explained according to the different dispersion patterns and morphologies of HA phase in the Fe matrix, presented in Fig. 3. As it can be seen in the SEM images in Fig. 3 (c & f), fine HA particles tend to form agglomerates in the microstructure. The HA agglomerates have voids inside them, also, some porosity and micro crevices are initiated at the agglomerate interface with the surrounding Fe matrix due to

Fig. 14. Estimated corrosion rates for pure Fe and Fe–HA composites obtained from immersion (weight loss) test.

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567

Fig. 15. Microstructure of (a) Pure Fe, (b) Fe–5 wt% HA (100–200 μm), (c) Fe–5 wt% HA (1–10 μm), (d) Fe–5 wt% HA (b1 μm), etched in nital 4% for 30 s. The micrographs (b) and (c) show the HA phase embedded well in the Fe matrix while the micrograph (d) shows HA detachment from the Fe matrix. The red arrows in (d) point the cavities created on the surface after HA detachment. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

defects. Generally, it can be inferred that, when HA content is increased and HA particle size is reduced (more HA surface area), the Fe–HA interface regions as well as the size and number of associated defects are increased too, which results in higher corrosion rates. In addition to the HA distribution pattern, another factor that can affect the corrosion rates of Fe–HA composites is the interfacial bond strength between Fe matrix and HA phase. The set of optical micrographs in Fig. 15 shows the surface of pure Fe and different Fe–HA composites etched in nital 4% for 30 s. Micrograph (d) is related to the microstructure of Fe– 5 wt% HA (b1 μm) in which some HA phase has fallen out of the Fe matrix, thus, more Fe substrate is exposed to the corrosion media. HA particle debonding from Fe matrix was occasionally observed for the composites of (b1 μm) group through surface polishing step and also during electrochemical corrosion tests in which minor amount of HA particles were found at the bottom of electrolyte container at the end of the experiments. It seemed that the ultrafine HA particles (b 1 μm) did not develop adequate interfacial bond and integrity with the Fe matrix compared to the other two HA powders in which good integrity and no loose particles were observed as shown in Fig. 15(b & c). It is highly possible that more Fe surface exposure and surface porosity generated due to the HA debonding at some local spots on the surface were also contributing factors in the high corrosion rates observed for the composites of (b 1 μm) group. Similarly, increased corrosion rate due to HA particle detachment from matrix was observed and reported by Anawati et al. for Ti–HA composites during specimen preparation and electrochemical corrosion tests [55]. When the influence of overall porosity of the sintered sample on the corrosion rate is considered, Table 4 Mechanical properties of human femur bone [50]. Type of loading

Loading direction

Yield strength σY (MPa)

Ultimate strength σU (MPa)

Fracture strain (%)

Tension

Longitudinal Transverse

114 –

135 53

3.1 0.7

although the composites in 100–200 μm group showed approximately 3% higher porosity, they still indicated lower corrosion rates due to the integrity, less defects, and no HA detachment at Fe–HA interface in their microstructure. In general, the slight increase in porosity of the composites as their HA contents increase corresponds to the increased corrosion rates observed. It must be stated that the in vitro degradation rates reported in this study can be incomparably higher than the actual in vivo rates for iron implants. The in vivo experiments for iron may show too slow [56] or no significant degradation [57] over an appreciable implantation time. Much of the inconsistencies and controversies related to the comparability of in vitro and in vivo results stem from the fact that it is not easy to simulate the complex physiological environment in vitro [58]. Some of the few but important factors that stand for lower in vivo corrosion rates are formation of biological phosphate layers [57,59] and protein adsorption [60] that inhibits the corrosion of implants.

4. Conclusion Considering the development of iron-based biodegradable materials for bone repair and hard tissue reconstruction, iron-hydroxyapatite (Fe–HA) composites are potentially good candidates due to their cytocompatibility, bioactivity, and the favorable impact of HA on bone healing and formation when these composites are implanted [22]. In this context, this work combined pure Fe powder and HA powders with three different particle size groups (b 1 μm, 1–10 μm, 100–200 μm) to produce powder metallurgy Fe–HA composites (HA Table 5 Ion concentrations in Kokubo's SBF and human blood plasma [37]. mmol·l−1

Na+

K+

Human blood plasma Kokubo's solution

142.0 142.0

5.0 1.5 5.0 1.5

Mg2+ Ca2+ Cl¯ 2.5 2.5

103.0 103.0

HCO3¯ HPO2− SO2− 4 4 27 4.2

1.0 1.0

0.5 0.5

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Table 6 Electrochemical data and estimated corrosion rates of pure Fe and Fe–HA composites. The rates were obtained from potentiodynamic polarization tests and immersion (weight loss) tests in Kokubo's solution. Material

ECorr (V)

iCorr (μA.cm−2)

CR – polarization (mmpy)

CR – immersion (mmpy)

Pure iron

−0.63 ± 0.03

17.31 ± 4.50

0.20 ± 0.05

0.11 ± 0.05

Iron + hydroxyapatite (100–200 μm)

2.5% 5% 10%

−0.55 ± 0.06 −0.68 ± 0.03 −0.61 ± 0.02

19.87 ± 3.28 25.90 ± 4.93 38.41 ± 4.43

0.23 ± 0.04 0.30 ± 0.06 0.46 ± 0.03

0.24 ± 0.01 0.36 ± 0.08 0.46 ± 0.07

Iron + hydroxyapatite (1–10 μm)

2.5% 5% 10%

−0.60 ± 0.04 −0.59 ± 0.03 −0.79 ± 0.06

23.26 ± 4.58 48.99 ± 7.99 70.63 ± 6.18

0.27 ± 0.05 0.57 ± 0.09 0.82 ± 0.07

0.25 ± 0.03 0.54 ± 0.01 0.88 ± 0.08

Iron + hydroxyapatite (b1 μm)

2.5% 5% 10%

−0.67 ± 0.05 −0.69 ± 0.06 −0.79 ± 0.05

43.84 ± 5.95 66.45 ± 6.34 92.38 ± 9.17

0.51 ± 0.07 0.77 ± 0.07 1.07 ± 0.11

0.32 ± 0.07 0.67 ± 0.06 1.03 ± 0.09

content = 2.5, 5, 10 wt%) and evaluated their mechanical properties and in vitro corrosion behaviors. The Fe–HA composites showed lower values of yield strength, tensile strengths, and ductility compared to those of pure Fe. It was demonstrated that the magnitude of drop in mechanical properties clearly depends not only on HA content, but also on HA particle size and HA dispersion pattern in the Fe matrix. While using the Fe powder mentioned in this study, the addition of lower amount of HA and larger-size HA particles, resulted in stronger and more ductile composites. Among the materials, only the Fe–2.5 wt% HA (100–200 μm) composite exhibited mechanical properties close to that of human femoral cortical bone, however, this does not imply that the rest of compositions are not suitable for bone repair and scaffold applications. Generally speaking, in selection of materials for such applications, some factors such as target area of implantation, mechanical design aspects (i.e. size of bone defect, geometry of implant, stress analysis), type of bone (cortical or trabecular) must be taken into account and there may be applications where stress levels and tolerance limits are not required to be as high as the human femoral cortical bone which was used as the baseline for comparison in this study. The in vitro corrosion rates of all the Fe–HA composites fabricated in this study were higher than pure Fe. As a general trend, increasing the HA content and reducing the HA particle size, created composites with higher corrosion rates which is beneficial for biodegradable implant applications. However, it should be noted that, since the in vitro corrosion rates are usually rough estimates of the in vivo rates, the realistic degradation behavior of materials should certainly be verified by further in vivo experiments. This study presents a model which demonstrate how the relative particle size between the two phases (Fe and HA) can determine the mechanical and corrosion properties of Fe–HA composites. The model can give insight into design strategies for processing of other bioactive composites (i.e. Ti–HA, Mg–HA) if powder metallurgy is selected as fabrication route and adjustable material properties are desired. The HA powder particles used in this study had perfectly or nearly spherical geometries. Further research can be done to investigate the influence of other HA morphologies such as rod-like, plate-like, needle-like, clusters, etc. Acknowledgements The authors would like to thank Prof. David Johnson at the Department of Materials Engineering at Purdue University who provided insight and expertise in the area of mechanical properties. We thank Dr. Ola Bergman, material development manager at Höganäs in Sweden for kindly providing the iron powder. Also, special thanks to Michael J. Kelley, graduate teaching assistant at Artisan and Fabrication Lab (AFL) at Purdue University for his constructive collaboration on machining operations involved in this work.

References [1] H. Li, Y. Zheng, L. Qin, Progress of biodegradable metals, Prog. Nat. Sci. Mater. Int. 24 (2014) 414–422. [2] Y.F. Zheng, X.N. Gu, F. Witte, Biodegradable metals, Mater. Sci. Eng. R 77 (2014) 1–34. [3] H. Hermawan, Biodegradable Metals: From Concept to Applications, Springer, Heidelberg, Germany, 2012. [4] L. Tan, X. Yu, P. Wan, K. Yang, Biodegradable materials for bone repairs: a review, J. Mater. Sci. Technol. 29 (2013) 503–513. [5] A.J. Salinas, M. Vallet-Regi, Bioactive ceramics: from bone grafts to tissue engineering, RSC Adv. 3 (2013) 11116–11131. [6] R.Z. LeGeros, Calcium phosphate-based osteoinductive materials, Chem. Rev. 108 (2008) 4742–4753. [7] P. Habibovic, M.C. Kruyt, M.V. Juhl, S. Clyens, R. Martinetti, L. Dolcini, N. Theilgaard, C.A. van Blitterswijk, Comparative in vivo study of six hydroxyapatite-based bone graft substitutes, J. Orthop. Res. 26 (2008) 1363–1370. [8] L. Sun, C.C. Berndt, C.P. Grey, Phase, structural and microstructural investigations of plasma sprayed hydroxyapatite coatings, Mater. Sci. Eng. A 360 (2003) 70–84. [9] H. Aoki, Medical Applications of Hydroxyapatite: Bone Mineral, Drug Delivery System, Cancer & HIV, IVH & CAPD, Dental Implant. Ishiyaku EuroAmerica, 1994. [10] L.L. Hench, Bioceramics: from concept to clinic, J. Am. Ceram. Soc. 74 (1994) 1487–1510. [11] Z. Jianguo, Z. Xingdong, C. Muller-Mai, U. Gross, The early host and material response of hydroxyapatite/β-tricalcium phosphate porous ceramics after implantation into the femur of rats, J. Mater. Sci. Mater. Med. 5 (1994) 243–251. [12] H. Yamasaki, Osteogenic response to porous hydroxyapatite ceramic under the skin of dogs, Biomaterials 13 (1992) 308–314. [13] A. Arifin, A.B. Sulong, N. Muhamad, J. Syarif, M.I. Ramli, Powder injection molding of HA/Ti6Al4V composite using palm stearin as based binder for implant material, Mater. Des. 65 (2015) 1028–1034. [14] D. Bovand, M. Yousefpour, S. Rasouli, S. Bagherifard, N. Bovand, A. Tamayol, Characterization of Ti-HA composite fabricated by mechanical alloying, Mater. Des. 65 (2015) 447–453. [15] M.R. Raza, A.B. Sulong, N. Muhamad, M.N. Akhtar, J. Rajabi, Effects of binder system and processing parameters on formability of porous Ti/HA composite through powder injection molding, Mater. Des. 87 (2015) 386–392. [16] A. Arifin, A.B. Sulong, N. Muhamad, J. Syarif, M.I. Ramli, Material processing of hydroxyapatite and titanium alloy (HA/Ti) composite as implant materials using powder metallurgy: a review, Mater. Des. (2014) 165–175. [17] C.Q. Ning, Y. Zhou, Correlations between the in vitro and in vivo bioactivity of the Ti/ HA composites fabricated by a powder metallurgy method, Acta Biomater. 4 (2008) 1944–1952. [18] C.Q. Ning, Y. Zhou, On the microstructure of biocomposites sintered from Ti, HA and bioactive glass, Biomaterials 25 (2004) 3379–3387. [19] E.S. Thian, N.H. Loh, K.A. Khor, S.B. Tor, Microstructures and mechanical properties of powder injection molded Ti-6Al-4V/HA powder, Biomaterials 23 (2002) 2927–2938. [20] C. Chu, X. Xue, J. Zhu, Z. Yin, Fabrication and characterization of titanium-matrix composite with 20 vol% hydroxyapatite for use as heavy load-bearing hard tissue replacement, J. Mater. Sci. Mater. Med. 17 (2006) 245–251. [21] M. Eriksson, M. Andersson, E. Adolfsson, E. Carlström, Titanium–hydroxyapatite composite biomaterial for dental implants, Powder Metall. 49 (2006) 70–77. [22] M.F. Ulum, A. Arafat, D. Noviana, A.H. Yusop, A.K. Nasution, M.R. Abdul Kadir, H. Hermawan, In vitro and in vivo degradation evaluation of novel iron-bioceramic composites for bone implant applications, Mater. Sci. Eng. C 36 (2014) 336–344. [23] X. Gu, W. Zhou, Y. Zheng, L. Dong, Y. Xi, D. Chai, Microstructure, mechanical property, bio-corrosion and cytotoxicity evaluations of Mg/HA composites, Mater. Sci. Eng. C 30 (2010) 827–832. [24] S. Zamani Khalajabadi, M.R. Abdul Kadir, S. Izman, M. Marvibaigi, The effect of MgO on the biodegradation, physical properties and biocompatibility of a Mg/HA/MgO nanocomposite manufactured by powder metallurgy method, J. Alloys Compd. 655 (2016) 266–280.

M. Dehestani et al. / Materials and Design 109 (2016) 556–569 [25] S. Zamani Khalajabadi, M.R. Abdul Kadir, S. Izman, M. Marvibaigi, Fabrication, biocorrosion behavior and mechanical properties of a Mg/HA/MgO nanocomposite for biomedical applications, Mater. Des. 88 (2015) 1223–1233. [26] S. Zamani Khalajabadi, M.R. Abdul Kadir, S. Izman, M.Z. Mohd Yusop, Facile fabrication of hydrophobic surfaces on mechanically alloyed-Mg/HA/TiO2/MgO bionanocomposites, Appl. Surf. Sci. 324 (2015) 380–392. [27] K.A. Khalil, A.A. Almajid, Effect of high-frequency induction heat sintering conditions on the microstructure and mechanical properties of nanostructured magnesium/hydroxyapatite nanocomposites, Mater. Des. 36 (2012) 58–68. [28] F. Witte, F. Feyerabend, P. Maier, J. Fischer, M. Stormer, C. Blawert, W. Dietzel, N. Hort, Biodegradable magnesium–hydroxyapatite metal matrix composites, Biomaterials 28 (2007) 2163–2174. [29] M. Peuster, P. Wohlsein, M. Brügmann, M. Ehlerding, K. Seidler, C. Fink, H. Brauer, A. Fischer, G. Hausdorf, A novel approach to temporary stenting: degradable cardiovascular stents produced from corrodible metal–results 6–18 months after implantation into New Zealand white rabbits, Heart 86 (2001) 563–569. [30] N. Mohd Daud, N. Boon Sing, A.H. Yusop, F.A. Abdul Majid, H. Hermawan, Degradation and in vitro cell–material interaction studies on hydroxyapatite-coated biodegradable porous iron for hard tissue scaffolds, J. Orthop. Trans. 2 (2014) 177–184. [31] A.H. Yusop, A.A. Bakir, N.A. Shaharom, M.R. Abdul Kadir, H. Hermawan, Porous biodegradable metals for hard tissue scaffolds: a review, Int. J. Biomater. 2012 (2012). [32] A. Reindl, R. Borowsky, S.B. Hein, J. Geis-Gerstorfer, P. Imgrund, F. Petzoldt, Degradation behavior of novel Fe/β-TCP composites produced by powder injection molding for cortical bone replacement, J. Mater. Sci. 49 (2014) 8234–8243. [33] M.P. Staiger, A.M. Pietak, Jerawala Huadmai, George Dias, magnesium and its alloys as orthopedic biomaterials: a review, Biomaterials 27 (2006) 1728–1734. [34] Standard Test Methods for Tension Testing of Metallic Materials, E8/E8M−13a. [35] Sintered Metal Materials, Excluding Hard Metals – Tensile Test Pieces, International Standard ISO 2740. [36] Standard Test Methods for Density of Compacted or Sintered Powder Metallurgy (PM) Products Using Archimedes' Principle, B962–13. [37] T. Kokubo, H. Takadama, How useful is SBF in predicting in vivo bone bioactivity? Biomaterials 27 (2006) 2907–2915. [38] Standard Practice for Preparing, Cleaning, and Evaluating Corrosion Test Specimens, G1–03. [39] Standard Practice for Laboratory Immersion Corrosion Testing of Metals, G31–72. [40] E.O. Hall, Yield Point Phenomena in Metals and Alloys, Macmillan, London, 1970. [41] A.H. Cottrell, B.A. Bilby, Dislocation theory of yielding and strain ageing of iron, Proc. Phys. Soc. Sect. A 62 (1949) 49–62. [42] G.T. Van Rooyen, The stress and strain distribution in a propagating Liiders front accompanying the yield-point phenomenon in iron, Mater. Sci. Eng. 3 (1968) 105–117. [43] H.L. Wain, Some observations on the yield point in zinc, Proc. Phys. Soc. Sect. B 65 (1952) 886–896.

569

[44] A.T. Churchman, A.H. Cottrell, Yield phenomenon and twinning in α-iron, Nature 167 (1951) 943. [45] J.D. Currey, The structure and mechanics of bone, J. Mater. Sci. 47 (2012) 41–54. [46] E. Verhulp, B. van Rietbergen, R. Huiskes, Load distribution in the healthy and osteoporotic human proximal femur during a fall to the side, Bone 42 (2008) 30–35. [47] J.H. Cole, M.C.H. van der Meulen, Whole bone mechanics and bone quality, Clin. Orthop. Relat. Res. 469 (2011) 2139–2149. [48] E.J. Cheal, M. Spector, W.C. Hayes, Role of loads and prosthesis material properties on the mechanics of the proximal femur after total hip arthroplasty, J. Orthop. Res. 10 (1992) 405–422. [49] H. Fonseca, D. Moreira-Goncalves, H.J. Coriolano, J.A. Duarte, Bone quality: the determinants of bone strength and fragility, Sports Med. 44 (2014) 37–53. [50] D.T. Reilly, A.H. Burstein, The elastic and ultimate properties of compact bone tissue, J. Biomech. 8 (1975) 393–405. [51] M. Kutz, Standard Handbook of Biomedical Engineering & Design, McGraw-Hill, 2003 (chpt. 8: Bone Mechanics p. 8.6). [52] R.W. McCalden, J.A. McGeough, M.B. Barker, C.M. Court-Brown, Age-related changes in the tensile properties of cortical bone: the relative importance of changes in porosity, mineralization, and microstructure, J. Bone Joint Surg. 75A (1993) 1193–1205. [53] S. Ramesh, C.Y. Tan, R. Tolouei, M. Amiriyan, J. Purbolaksono, I. Sopyan, W.D. Teng, Sintering behavior of hydroxyapatite prepared from different routes, Mater. Des. 34 (2012) 148–154. [54] G. Muralithran, S. Ramesh, The effects of sintering temperature on the properties of hydroxyapatite, Ceram. Int. 26 (2000) 221–230. [55] H. Anawati, H. Tanigawa, H. Asoh, T. Ohno, M. Kubota, S. Ono, Electrochemical corrosion and bioactivity of titanium–hydroxyapatite composites prepared by spark plasma sintering, Corros. Sci. 70 (2013) 212–220. [56] M. Peuster, C. Hesse, T. Schloo, C. Fink, P. Beerbaum, C. von Schnakenburg, Longterm biocompatibility of a corrodible peripheral iron stent in the porcine descending aorta, Biomaterials 27 (2006) 4955–4962. [57] A. Drynda, T. Hassel, F.W. Bach, M. Peuster, In vitro and in vivo corrosion properties of new iron–manganese alloys designed for cardiovascular applications, J. Biomed. Mater. Res. Part B 103 (2015) 649–660. [58] A.H.M. Sanchez, B.J.C. Luthringer, F. Feyerabend, R. Willumeit, Mg and Mg alloys: how comparable are in vitro and in vivo corrosion rates? A review, Acta Biomater. 13 (2015) 16–31. [59] F. Witte, V. Kaese, H. Haferkamp, E. Switzer, A. Meyer-Lindenber, C.J. Wirth, H. Windhagen, In vivo corrosion of four magnesium alloys and the associated bone response, Biomaterials 26 (2005) 3557–3563. [60] M.A. Khan, R.L. Williams, D.F. Williams, The corrosion behavior of Ti–6Al–4V, Ti– 6Al–7Nb and Ti–13Nb–13Zr in protein solutions, Biomaterials 20 (1999) 631–637.