Biodegradable Implantation Material: Mechanical Properties and Surface Corrosion Mechanism of Mg-1Ca-0.5Zr Alloy

metals

Article Biodegradable Implantation Material: Mechanical Properties and Surface Corrosion Mechanism of Mg-1Ca-0.5Zr Alloy

Yen-Ting Chen, Fei-Yi Hung * and Jie-Cheng Syu

Department of Materials Science and Engineering, National Cheng Kung University, Tainan 701, Taiwan * Correspondence: [email protected]; Tel.: +886-6-2757-5756-2950; Fax: +886-6234-6290

Received: 19 June 2019; Accepted: 2 August 2019; Published: 6 August 2019

Abstract: Mg alloy is suitable for biomedical implants as the mechanical properties of Mg are close to those of human bone. Ca is a major element in bone and Zr has a great grain refinement effect. Hence, we developed Mg-1Ca-0.5Zr alloy (XK105) as a biodegradable biomaterial and investigated its mechanical properties and surface corrosion mechanism. The results showed that heat treatment made the secondary phase homogeneous. Tensile tests showed that the heat treatment increased ductility, and that the tensile stress results in the extrusion direction showed better ductility than that in the transverse direction because of the fiber texture and extrusion characteristics. Electrochemistry test results showed that XK105 after heat treatment had a lower corrosion rate than that before heat treatment and that of pure Mg. XK105 after heat treatment formed a calcium phosphate layer after immersion in simulated body fluid; this layer protects Mg from corrosion. Surface roughening treatment increased corrosion because pits on the surface promoted pitting corrosion. This study developed Mg-1Ca-0.5Zr alloy as a biomedical implant material. The results can be used as a reference for the biomedical material industry.

Keywords: Mg-Ca-Zr; mechanical properties; erosion; corrosion

1. Introduction Common biomedical implant materials include metals, polymers, and ceramics. Metal-based biomedical materials have better mechanical properties than those of polymers and ceramics [1] and are thus much more widely used in practical applications. Common metal-based biomedical materials such as stainless steel [2], Ti alloys [3], and Co-based alloys [4] have elastic moduli that are very different from that of human bone. This leads to a stress shielding effect, which makes new bone tissues lack stimulation during growth and leads to insufficient recovery of bone strength [5]. Another disadvantage of common metal-based biomedical materials is their non-degradability [6]; they must be removed in a second operation after the affected area has recovered. In view of these issues, Mg alloys have been developed as next-generation biomedical materials [7]. Mg alloys have several advantages as biomaterials. Mg is an essential mineral nutrient, so it has good biosafety [8]. Mg alloys have biomechanical properties similar to those of human bone, which can avoid stress shielding effects and improve the strength of bone tissue after healing [9]. Mg is degradable in vivo, and when used for short-term implants, it does not need to be removed in a second operation [10]. The products produced by the degradation of Mg in vivo can be excreted through the kidney [11]. Mg alloy combines well with surrounding bone tissues after degradation and accelerates the formation of biological calcium phosphate [12]. Mg alloys can stimulate the differentiation and growth of osteoblasts and promote bone tissue production [13].

Metals 2019, 9, 857; doi:10.3390/met9080857 www.mdpi.com/journal/metals Metals 2019, 9, 857 2 of 15

The concentration of chloride ions in the human body is as high as 150 mmol/L, making the degradation of pure Mg very fast [14]. Such a fast degradation rate may cause slow recovery of the affected area and even lead to tissue necrosis [15]. Adding other metallic elements to form Mg alloys can solve this issue. Mg-Ca alloys have become a focus in the development of Mg-based biomedical materials [16] because Ca is not only harmless to the human body, but also an important component that makes up bone [17]. The Ca ions released during degradation help the healing of bone tissue [18]. Adding Ca to Mg alloys can enhance their mechanical properties through grain refinement and solid solution strengthening [19]. However, if the Ca content is excessive, a large amount of Mg2Ca will be precipitated, decreasing mechanical properties and corrosion resistance [20]. Mg2Ca has a lower corrosion potential than that of the Mg base, so it is preferentially corroded as a sacrificial anode and can thus increase the corrosion resistance of the Mg base [21]. However, in excess amounts, Mg2Ca forms a continuous network structure at the grain boundary, causing cracks [22]. According to previous studies, Mg alloy with 1 wt.% Ca has good mechanical properties and corrosion resistance [23]. Both Zr and Mg are hexagonal close-packed (HCP) structures and the two elements have very similar lattice constants, so adding Zr to Mg alloy can refine grains and improve mechanical properties [24,25]. However, in excess amounts, Zr aggregates into particles larger than 5 µm and reduces the corrosion resistance of Mg alloys [26]. The casting Mg-Ca-Zr alloy has better tensile strength and elongation than Mg-Ca alloy at all Ca content ratios [27]. Heat treatment is important for Mg alloys in application. The temperature of post heat treatment is generally around 380–420 ◦C for 6–12 h [28,29] to achieve the best mechanical properties. For a novel Mg alloy, the heat treatment conditions should be suggested in this range. In this study, the Mg alloy Mg-1Ca-0.5Zr (98.5 wt.% Mg, 1 wt.% Ca, and 0.5 wt.% Zr) is subjected to heat treatment to homogenize and stabilize its material properties. The heat treatment condition was 400 ◦C for 8 h at atmosphere which followed the common condition of Mg-Ca alloy [28]. A series of experiments were conducted to evaluate the feasibility of Mg-1Ca-0.5Zr alloy as a biomedical implant material. The results can be used as a reference for the biomedical material industry.

2. Material and Methods Mg-1Ca-0.5Zr was produced in a smelting furnace. A 99.9 wt.% pure Mg block was placed in a stainless-steel bowl and heated to 700 ◦C until the Mg block had completely melted into a liquid state. Then, Mg-20Ca and Mg-33Zr blocks were placed into the molten Mg. After the alloy blocks had melted completely, the molten alloy was poured into a mold and allowed to cool in air to room temperature. For all casting, argon was used as a protective gas. The casting product, Mg-1Ca-0.5Zr alloy, is denoted as XK105 following the ASTM (American Society for Testing and Materials) naming convention. For experiments, we processed XK105 into an extruded sheet, denoted as XK105-F. To enhance the mechanical properties, XK105 was subjected to heat treatment at 400 ◦C for 8 h in an air furnace. To decrease the formation of the crystalline phase, the alloy was water-quenched after heat treatment. The heat-treated XK105 is denoted XK105-H. For microstructural observation, epoxy was used to cold-mount and embed XK105-F and XK105- H. The samples were then polished using 1 and 0.3 µm aluminum oxide powder. After polishing, the samples were corroded using a solution of 4.2 g of picric acid, 10 mL of acetic acid, and 90 mL of ethanol. The microstructure was observed using light microscopy (LM) (ZEISS, Oberkochen, Germany) and scanning electron microscopy (SEM) (HITACHI, Tokyo, Japan). Energy-dispersive X-ray spectrometry (EDS) was used for elemental composition semi-quantitative analysis. This study used tensile tests to determine the tensile strength of the extruded proﬁles of XK105 and the mechanical properties after heat treatment. Three specimens of XK105-F were taken along the extrusion direction (ED) and the transverse direction (TD). After the samples were cut to an appropriate size, they were punched into a ﬁxed size using a punch press. A stretching test was conducted with an initial strain rate of 8.3 10 4 s 1 at room temperature. We used SEM and EDS to investigate the × − − stretch-breaking behavior. Metals 2019, 9, 857 3 of 15

An erosion wear test was conducted using XK105-F and XK105-H. A particle spray tester (ChunMetals 2019 Yen, 9 Testing, x FOR PEER Machines REVIEW Co., Taichung, Taiwan) was used for erosion wear and surface roughening.3 of 14 The environmental conditions in the erosion wear test were room temperature and a relative humidity ofThe 50% environmental to 70%. The conditions erosion particles in the wereerosion Al2 Owear3 (average test were diameter: room temperature 150 µm) and and SiO 2a(average relative 2 3 2 diameter:humidity 295of 50%µm), to and 70%. the The erosion erosion angle particles was 15◦ wereto 90 ◦Al. TheO particle(average rate diameter: of erosion 150 was μm) 1.2 and g/s, SiO and 2 the(average pressure diameter: was set 295 to μ 3m), kg/ cmand the(0.29 erosion MPa). angle The average was 15° jetto rate90°. The of Al particle2O3 was rate calculated of erosion as was 73 m 1.2/s, 2 2 3 andg/s, and that ofthe SiO pressure2 was 66 was m/ s.set Surface to 3 kg/cm roughening (0.29 MPa). treatment The wasaverage conducted jet rate at of erosion Al O was angles calculated of 15◦, 45 as◦, 2 and73 m/s, 90◦ .and After that surface of SiO roughening, was 66 m/s. the Surface samples roughening were observed treatm usingent was LM conducted and SEM. at erosion angles of 15°,A 45°, potentiodynamic and 90°. After polarizationsurface roughening test was, the conducted samples were using observed pure Mg, using XK105-F, LM and and SEM. XK105-H samples.A potentiodynamic In the test, a three-electrode polarization system test was was conducted used to obtain using the pure polarization Mg, XK105-F, curve. and The XK105-H reference electrodesamples. In was the a test, saturated a three-electrode calomel electrode system (SCE), was us theed saltto obtain bridge the was polarization saturated KCl curve. solution, The reference and the auxiliaryelectrode electrodewas a saturated was a platinum calomel electrode.electrode (SCE), The specimens the salt bridge were embedded was saturated in epoxy KClresin solution, and thenand carefullythe auxiliary ground. electrode The employedwas a platinum electrolyte electrode. was a The revised specimens simulated were body embedded ﬂuid (R-SBF) in epoxy solution resin [and30], whichthen carefully has ion compositionground. The andemployed concentration electrolyte similar was to a thoserevised of simulated human blood body plasma. fluid (R-SBF) The scan solution rate in the[30], polarization which has ion test composition was 1 mV/s, and from concentration2.2 to 1.2 similar V, and ato constant those of temperaturehuman blood of plasma. 37 C was The used.scan − − ◦ Therate polarizationin the polarization curve hastest awas linear 1 mV/s, Tafel from region −2.2 around to −1.2 the V, corrosion and a constant potential. temperature The corrosion of 37 current°C was canused. be The obtained polarization from the curve anodization has a linear curves Tafel using regi theon around extrapolation the corrosion method. potential. The corrosion currentThe can immersion be obtained test from wasused the anodization to evaluate thecurves corrosion using rate.the extrapolation The corrosion method. rate was calculated as: The immersion test was used to evaluate the corrosion rate. The corrosion rate was calculated as: m1 m0 Corrosion rate = − Corrosion rate = A T ∙· where m11 is the mass after immersion, m0 isis the the mass mass before before immersion, A is the total surface area of the sample,sample, andand T T is theis the immersion immersion time. time. The immersionThe immersion solution solution was R-SBF. was R-SBF. XK105-H XK105-H and surface- and roughenedsurface- roughened XK105-H XK105-H sampled weresampled immersed were immersed in R-SBF forin 1,R-SBF 3, and for 7 days1, 3, atand a constant7 days at temperature a constant oftemperature 37 ◦C. of 37 °C.

3. Results and Discussion 3.1. Metallographic Analysis 3.1. Metallographic Analysis Figure1a,b respectively show metallographic images of XK105-F and XK105-H observed using Figure 1a,b respectively show metallographic images of XK105-F and XK105-H observed using LM. LM. Crystal phases appear along ED for XK105-F. The crystal phases of XK105-H were not along ED. Crystal phases appear along ED for XK105-F. The crystal phases of XK105-H were not along ED. This This indicates that heat treatment made the crystal phases dissolve into the matrix, homogenizing indicates that heat treatment made the crystal phases dissolve into the matrix, homogenizing the material. the material. The grains were subjected to dynamic recrystallization during extrusion, and thus The grains were subjected to dynamic recrystallization during extrusion, and thus isometric crystals isometric crystals formed. The calculated average grain sizes of XK105-F and XK105-H were 14.2 and formed. The calculated average grain sizes of XK105-F and XK105-H were 14.2 and 17.1 μm, respectively. 17.1 µm, respectively.

(a) (b)

Figure 1. Light microscopy (LM) microstructural images. ( a) XK105-F and ( b) XK105-H.

SEM observation was conducted to verify the secondary phases of XK105. A Ca-rich crystal phase can be seen in Figure 2a; the phase was identified as Mg2Ca [31]. Most of the Mg2Ca phase was along ED. A Zr-rich phase can be seen in Figure 2b; the phase was identified as pure Zr particles which accumulated at grains [32]. Metals 2019, 9, 857 4 of 15

SEM observation was conducted to verify the secondary phases of XK105. A Ca-rich crystal phase can be seen in Figure2a; the phase was identiﬁed as Mg 2Ca [31]. Most of the Mg2Ca phase was along ED. A Zr-rich phase can be seen in Figure2b; the phase was identiﬁed as pure Zr particles which accumulated at grains [32]. Metals 2019, 9, x FOR PEER REVIEW 4 of 14

Figure 2.2. SecondarySecondary phase phase analysis analysis of XK105-F usin usingg scanning electron microscopy (SEM) and

energy-dispersive X-rayX-ray spectrometryspectrometry (EDS).(EDS). ((aa)) MgMg22Ca and (b)) ZrZr particles.particles.

3.2. Tensile Properties andand FractureFracture MorphologyMorphology Figure Figure3 3;and Figure Figure 44 respectively respectively show show the the tensile tensile stress-strain stress-strain curve curve and and tensile tensile stress stress of XK105. of XK105. The Thetensile tensile stress stress of XK105-F of XK105-F along along TD is TD about is about 220 MPa. 220 MPa. According According to reference to reference [33], the [33 tensile], the tensile stress stressalong alongTD of TDpure of Mg pure was Mg 173 was MPa, 173 and MPa, which and whichof Mg-0.1Ca of Mg-0.1Ca was 180 was MPa. 180 The MPa. tensile The stress tensile of stress human of humanbone is boneabout is 150 about MPa. 150 XK105 MPa. thus XK105 has thussufficien hast su tensileﬃcient strength tensile to strength apply to to stress apply load to stress in the load human in the body. human The body.yield stress The yield of XK105-F stress of along XK105-F TD was along about TD was180 aboutMPa, which 180 MPa, of pure which Mg of was pure 103 Mg MPa was and 103 MPawhich and of whichMg- 0.1Ca of Mg-0.1Ca was 92 MPa was [33]. 92 MPa The results [33]. The showed results the showed adding the Zr addingcan improve Zr can the improve tensile strength the tensile effectively. strength eFigureﬀectively. 5 shows Figure the5 elongation shows the of elongation XK105. The of elongati XK105.on The was elongation about 4–8% was and about was 4–8%caused and by wasthe addition caused byof Ca the [34]. addition Elongation of Ca [increased34]. Elongation after heat increased treatmen aftert because heat treatment of the homogenization because of the of homogenization textures and the ofsolid textures solution and of the secondary solid solution phases. of theFigure secondary 6 shows phases. the Young’s Figure mo6dulus shows of theXK105. Young’s The value modulus is about of XK105.35–55 GPa; The it value increased is about after 35–55 heat GPa;treatment it increased because after the Ca heat atoms treatment dissolved because into the matrix Ca atoms and dissolved changed intothe lattice. the matrix Although and changed the Young’s the lattice.modulus Although of XK105 the is much Young’s higher modulus than that of XK105 of human is much bone higher (3–20 GPa), than thatit has of greater human similarity bone (3–20 compared GPa), it hasto other greater medical similarity implant compared materials. to otherThe tensile medical strength implant along materials. TD is Thehigher tensile than strengththat along along ED. The TD XRD is higher pattern than (Figure that along 7) shows ED. that The the XRD mean pattern peak (Figurealong ED7) showswas (101 that1). theThis mean is due peak to the along HCP ED crystals was (slipping1011). This along is dueED, which to the HCPled to crystalsa small angle slipping between along the ED, (0001) which plane led toand a smallED. When angle the between tensile strength the (0001) transferred plane and toward ED. When ED, the the dislocatio tensile strengthns more easily transferred slipped toward compared ED, theto that dislocations toward TD. more Theref easilyore, slipped the elongation compared of toXK105 that towardalong ED TD. was Therefore, higher but the the elongation tensile stress of XK105 was alonglower EDcompared was higher to those but along the tensile TD. stress was lower compared to those along TD.

Figure 3. Tensile stress-strain curves of XK105. (H-ED: extrusion direction of XK105-H; H-TD: transverse direction of XK105-H; F-ED: extrusion of XK105-F; F-TD: transverse of XK105-F). Metals 2019, 9, x FOR PEER REVIEW 4 of 14

Figure 2. Secondary phase analysis of XK105-F using scanning electron microscopy (SEM) and

energy-dispersive X-ray spectrometry (EDS). (a) Mg2Ca and (b) Zr particles.

3.2. Tensile Properties and Fracture Morphology Figure 3; Figure 4 respectively show the tensile stress-strain curve and tensile stress of XK105. The tensile stress of XK105-F along TD is about 220 MPa. According to reference [33], the tensile stress along TD of pure Mg was 173 MPa, and which of Mg-0.1Ca was 180 MPa. The tensile stress of human bone is about 150 MPa. XK105 thus has sufficient tensile strength to apply to stress load in the human body. The yield stress of XK105-F along TD was about 180 MPa, which of pure Mg was 103 MPa and which of Mg- 0.1Ca was 92 MPa [33]. The results showed the adding Zr can improve the tensile strength effectively. Figure 5 shows the elongation of XK105. The elongation was about 4–8% and was caused by the addition of Ca [34]. Elongation increased after heat treatment because of the homogenization of textures and the solid solution of the secondary phases. Figure 6 shows the Young’s modulus of XK105. The value is about 35–55 GPa; it increased after heat treatment because the Ca atoms dissolved into the matrix and changed the lattice. Although the Young’s modulus of XK105 is much higher than that of human bone (3–20 GPa), it has greater similarity compared to other medical implant materials. The tensile strength along TD is higher than that along ED. The XRD pattern (Figure 7) shows that the mean peak along ED was (1011). This is due to the HCP crystals slipping along ED, which led to a small angle between the (0001) plane and ED. When the tensile strength transferred toward ED, the dislocations more easily slipped compared to that toward TD. Therefore, the elongation of XK105 along ED was higher but the tensile stress was lower compared to those along TD. Metals 2019, 9, 857 5 of 15

Metals 2019, 9, x FOR PEER REVIEW 5 of 14

Figure 3. Metals 2019, 9, x FORTensile PEER stress-strain REVIEW curves of XK105. (H-ED: extrusion direction of XK105-H; H-TD: transverse5 of 14 directionFigure 3. of Tensile XK105-H; stress-strain F-ED: extrusion curves of XK105-F;of XK105. F-TD: (H-ED: transverse extrusion of XK105-F).direction of XK105-H; H-TD: transverse direction of XK105-H; F-ED: extrusion of XK105-F; F-TD: transverse of XK105-F).

Figure 4. Tensile stress analysis of XK105 (YS: yield stress; UTS: ultimate tensile stress). Figure 4. Tensile stress analysis of XK105 (YS: yield stress; UTS: ultimate tensile stress). Figure 4. Tensile stress analysis of XK105 (YS: yield stress; UTS: ultimate tensile stress).

Figure 5. Elongation analysis of XK105 (UE: unif uniformorm elongation; TE: total elongation). Figure 5. Elongation analysis of XK105 (UE: uniform elongation; TE: total elongation).

Figure 6. Young’s modulus of XK105. (H-ED: extrusion direction of XK105-H; H-TD: transverse directionFigure 6. ofYoung’s XK105-H; modulus F-ED: extrusionof XK105. of (H-ED:XK105-F; extrusio F-TD: ntransverse direction of of XK105-F). XK105-H; H-TD: transverse direction of XK105-H; F-ED: extrusion of XK105-F; F-TD: transverse of XK105-F). Metals 2019, 9, x FOR PEER REVIEW 5 of 14

Figure 4. Tensile stress analysis of XK105 (YS: yield stress; UTS: ultimate tensile stress).

Metals 2019, 9, 857 6 of 15 Figure 5. Elongation analysis of XK105 (UE: uniform elongation; TE: total elongation).

Figure 6. Young’s modulus of XK105. (H-ED: extrusion direction of XK105-H; H-TD: transverse Figure 6. Young’s modulus of XK105. (H-ED: extrusion direction of XK105-H; H-TD: transverse Metalsdirection 2019, 9, xof FOR XK105-H; PEER REVIEW F-ED: extrusion of XK105-F; F-TD: transverse of XK105-F). 6 of 14 direction of XK105-H; F-ED: extrusion of XK105-F; F-TD: transverse of XK105-F).

Figure 7. Texture analysis of XK105-H. (a) XRD patterns of XK105-H on ED and TD planes and Figure 7. Texture analysis of XK105-H. (a) XRD patterns of XK105-H on ED and TD planes and (b) illustration of dislocation slip. (b) illustration of dislocation slip. Figure8a,b shows the tensile fracture morphology of XK105-H on the ED plane. Dimples can be Figure 8a,b shows the tensile fracture morphology of XK105-H on the ED plane. Dimples can be observed on the fracture surface. Mg2Ca, which did not dissolve into the matrix, appears inside the observed on the fracture surface. Mg2Ca, which did not dissolve into the matrix, appears inside the dimples. Figure8c,d shows the tensile fracture morphology of XK105-H on the TD plane. Thin wall dimples. Figure 8c,d shows the tensile fracture morphology of XK105-H on the TD plane. Thin wall structures can be observed. These structures are cleavages of brittle failure. structures can be observed. These structures are cleavages of brittle failure.

Figure 8. Tensile fracture morphology of XK105-H observed by SEM. Fracture surface on (a,b) ED plane and (c,d) TD plane.

3.3. Fracture Toughness Medical implants easily deform when subjected to a stress impact. The fracture toughness of XK105 is thus an important property. Figure 9 shows the fracture toughness of XK105. The fracture Metals 2019, 9, x FOR PEER REVIEW 6 of 14

Figure 7. Texture analysis of XK105-H. (a) XRD patterns of XK105-H on ED and TD planes and (b) illustration of dislocation slip.

Figure 8a,b shows the tensile fracture morphology of XK105-H on the ED plane. Dimples can be observed on the fracture surface. Mg2Ca, which did not dissolve into the matrix, appears inside the Metalsdimples.2019, Figure9, 857 8c,d shows the tensile fracture morphology of XK105-H on the TD plane. Thin7 ofwall 15 structures can be observed. These structures are cleavages of brittle failure.

FigureFigure 8.8. TensileTensile fracturefracture morphologymorphology ofof XK105-HXK105-H observedobserved byby SEM.SEM. FractureFracture surfacesurface onon ((aa,b,b)) EDED planeplane andand (c(c,d,d)) TD TD plane. plane. 3.3. Fracture Toughness 3.3. Fracture Toughness Medical implants easily deform when subjected to a stress impact. The fracture toughness of Medical implants easily deform when subjected to a stress impact. The fracture toughness of XK105 is thus an important property. Figure9 shows the fracture toughness of XK105. The fracture XK105 is thus an important property. Figure 9 shows the fracture toughness of XK105. The fracture toughness was lower when the stress was applied toward TD. There is a non-plane strain of XK105, and it showed that the thickness of material was not thick enough to fracture as a plane strain. LM observations indicate that the fracture cracks had more bending along ED than TD, as shown in Figure 10a,c. The cracks thus required more energy to spread, so the toughness of the sample along ED was higher than that along TD. Figure 10b,d show the fracture surface of XK105-H. A smoother fracture surface, which resulted from brittle fracture on the TD plane, can be observed. Therefore, toughness on the TD plane was lower than that on the ED plane. Figure 11 shows the secondary surface morphology and fracture mechanism. The end of the crack on the ED plane has a branch morphology; the branching of that on the TD plane is not obvious. This was caused by the brittle secondary phase easily cracking when subjected to stress; the cracks spread along TD while tensile stress was applied along ED. The density of the secondary phase was thus lower on the crack-spreading path. Therefore, the cracks on the ED plane easily branched and had much longer spreading paths than those of cracks on the TD plane. Metals 2019, 9, x FOR PEER REVIEW 7 of 14 Metals 2019, 9, x FOR PEER REVIEW 7 of 14 toughness was lower when the stress was applied toward TD. There is a non-plane strain of XK105, toughness was lower when the stress was applied toward TD. There is a non-plane strain of XK105, and it showed that the thickness of material was not thick enough to fracture as a plane strain. LM and it showed that the thickness of material was not thick enough to fracture as a plane strain. LM observations indicate that the fracture cracks had more bending along ED than TD, as shown in observations indicate that the fracture cracks had more bending along ED than TD, as shown in Figure 10a,c. The cracks thus required more energy to spread, so the toughness of the sample along Figure 10a,c. The cracks thus required more energy to spread, so the toughness of the sample along ED was higher than that along TD. Figure 10b,d show the fracture surface of XK105-H. A smoother ED was higher than that along TD. Figure 10b,d show the fracture surface of XK105-H. A smoother fracture surface, which resulted from brittle fracture on the TD plane, can be observed. Therefore, fracture surface, which resulted from brittle fracture on the TD plane, can be observed. Therefore, toughness on the TD plane was lower than that on the ED plane. Figure 11 shows the secondary toughness on the TD plane was lower than that on the ED plane. Figure 11 shows the secondary surface morphology and fracture mechanism. The end of the crack on the ED plane has a branch surface morphology and fracture mechanism. The end of the crack on the ED plane has a branch morphology; the branching of that on the TD plane is not obvious. This was caused by the brittle morphology; the branching of that on the TD plane is not obvious. This was caused by the brittle secondary phase easily cracking when subjected to stress; the cracks spread along TD while tensile secondary phase easily cracking when subjected to stress; the cracks spread along TD while tensile stress was applied along ED. The density of the secondary phase was thus lower on the stress was applied along ED. The density of the secondary phase was thus lower on the crack- spreading path. Therefore, the cracks on the ED plane easily branched and had much longer crack- spreading path. Therefore, the cracks on the ED plane easily branched and had much longer Metalsspreading2019, 9 paths, 857 than those of cracks on the TD plane. 8 of 15 spreading paths than those of cracks on the TD plane.

Figure 9.9. Fracture toughness results for XK105.XK105. (H-ED: (H-ED: extrusion extrusion direction direction of XK105-H; H-TD: Figure 9. Fracture toughness results for XK105. (H-ED: extrusion direction of XK105-H; H-TD: transversetransverse directiondirection ofof XK105-H;XK105-H; F-ED:F-ED: extrusionextrusion ofof XK105-F;XK105-F; F-TD:F-TD: transversetransverse ofof XK105-F).XK105-F). transverse direction of XK105-H; F-ED: extrusion of XK105-F; F-TD: transverse of XK105-F).

Figure 10.10. Crack morphology of XK105-H after fracture toughness test. ( aa)) LM LM image on ED plane, Figure 10. Crack morphology of XK105-H after fracture toughness test. (a) LM image on ED plane, ((bb)) SEMSEM imageimage onon EDED plane,plane, ((cc)) LMLM imageimage onon TDTD plane,plane, andand ((dd)) SEMSEM imageimage onon TDTD plane.plane. (b) SEM image on ED plane, (c) LM image on TD plane, and (d) SEM image on TD plane. MetalsMetals 20192019,, 99,, 857x FOR PEER REVIEW 98 ofof 1514 Metals 2019, 9, x FOR PEER REVIEW 8 of 14

Figure 11. Fracture mechanism of XK105-H. (a) LM secondary surface morphology on ED plane, (b) Figurefracture 11. mechanism FractureFracture mechanism illustration mechanism offor of XK105-H. ED XK105-H. plane, ( a(c) ( )aLM LM) LM secondary secondary secondary surfacesurface surface morphologymorphology morphology on TDED on plane, EDplane, plane, and(b) fracture((bd)) fracturefracture mechanism mechanismmechanism illustration illustrationillustration for ED forfor plane, EDTD plane,plane. (c) LM (c) secondary LM secondary surface surface morphology morphology on TD on plane, TD plane, and (andd) fracture (d) fracture mechanism mechanism illustration illustration for TD for plane. TD plane. 3.4. Erosion Properties and Surface Roughness 3.4. Erosion Erosion Properties Properties and Surface Roughness Figure 12 shows the relationship between erosion rate and erosion angle. For both Al2O3 or SiO2, Figure 12 shows the relationship between erosion rate and erosion angle. For both Al2O3 or SiO2, the angleFigure for 12 the shows highest the relationshiperosion rate between changed er fromosion 45° rate to and 30° erosionafter heat angle. treatment. For both This Al 2showsO3 or SiO that2, the angle for the highest erosion rate changed from 45◦ to 30◦ after heat treatment. This shows that theXK105 angle had for lower the highest hardness erosion and higher rate changed ductility from after 45° heat to treatment.30° after heat The treatment. erosion rate This obtained shows withthat XK105 had lower hardness and higher ductility after heat treatment. The erosion rate obtained with XK105Al2O3 particleshad lower was hardness higher thanand higherthat obtained ductility with after SiO heat2 due treatment. to the higher The erosionhardness rate of obtainedAl2O3. with Al2O3 particles was higher than that obtained with SiO2 due to the higher hardness of Al2O3. Al2O3 particles was higher than that obtained with SiO 2 due to the higher hardness of Al2O3. 14

14 F-SiO2 F-SiO2 12 H-SiO2 12 H-SiO2 F-Al2O3 ) 10 -4 F-Al2O3

) H-Al O 10 2 3 -4 H-Al O 8 2 3 8 6 6 4 4 Erosion Rate (g/g*10 Rate Erosion

Erosion Rate (g/g*10 Rate Erosion 2 2 0 0 10 20 30 40 50 60 70 80 90 100 10 20 30 40 50 60 70 80 90 100 Impact Angle (deg.) Impact Angle (deg.) Figure 12. Erosion rate of XK105-F and XK105-H. Figure 12. Erosion rate of XK105-F and XK105-H. Figure 12. Erosion rate of XK105-F and XK105-H. Figure 13 shows the surface roughness after the erosion test. The roughness obtained with Al2O3 and SiOFigure2 was 13 similar. shows the The surface highest roughness and lowest after roughness the erosion was test. obtained The roughness with erosion obtained angles with of 90 ◦Aland2O3 Figure 13 shows the surface roughness after the erosion test. The roughness obtained with Al2O3 15and◦, respectively.SiO2 was similar. The measurement The highest and of these lowest erosion roughness properties was obtained could be appliedwith erosion to later angles experiments of 90° and to and SiO2 was similar. The highest and lowest roughness was obtained with erosion angles of 90° and clarify15°, respectively. the relationship The measurement between anti-corrosion of these erosion ability properties and surface could roughness. be applied This to will later have experiments a beneﬁt 15°, respectively. The measurement of these erosion properties could be applied to later experiments ofto increasingclarify the the relationship bioreactivity between for medical anti-corrosion applications. ability and surface roughness. This will have a tobenefit clarify of theincreasing relationship the bioreactivity between anti-corrosion for medical applications. ability and surface roughness. This will have a benefit of increasing the bioreactivity for medical applications. Metals 2019, 9, x FOR PEER REVIEW 9 of 14

Metals 2019, 9, 857 10 of 15 Metals 2019, 9, x FOR PEER REVIEW 9 of 14

Figure 13. Surface roughness of XK105-H after erosion test.

3.5. Potentiodynamic Polarization Test The potentiodynamic polarization test was used to analyze the corrosion rate of XK105 in simulated body fluid. Surface roughening treatment can effectively increase cell adhesion [35]. Therefore, in this study,FigureFigure surface 13.13. Surfaceroughen roughness treatment ofof XK105-HXK105-Husing the afterafter particle erosionerosion erosion test.test. wear method was applied. The potentiodynamic polarization test was conducted to determine anti-corrosion ability. The3.5.3.5. Potentiodynamicelectrochemical Polarizationproperties Testof XK105 calculated from the potentiodynamic polarization test results are shown in Table 1. The results show that heat treatment improved the anti-corrosion TheThe potentiodynamicpotentiodynamic polarization test was used to analyzeanalyze the corrosioncorrosion raterate ofof XK105XK105 inin propertysimulated of bodyXK105, fluid. making Surface it better roughening than that treatmentof pure Mg. can This eff ectivelywas due increase to Mg2Ca cell being adhesion dissolved [35]. intosimulated the matrix body during fluid. heat Surface treatment, roughening and the treatm effect entof galvanic can effectively corrosion increase becoming cell inactive, adhesion which [35]. Therefore,Therefore, inin thisthis study,study, surfacesurface roughenroughen treatmenttreatment usingusing thethe particleparticle erosionerosion wearwear methodmethod waswas reducedapplied. the The corrosion potentiodynamic current and polarization potential. testThe was potentiodynamic conducted to determinepolarization anti-corrosion curve is shown ability. in Figureapplied. 14. The The potentiodynamic reduction rate of polarization XK105-F is testlowe wasr than conducted that of pure to determine Mg. XK105-H anti-corrosion had the lowestability. TheThe electrochemicalelectrochemical properties properties of XK105of XK105 calculated calculat fromed from the potentiodynamic the potentiodynamic polarization polarization test results test reduction rate. The low reduction rate represents the low formation rate of H2, which is harmless to areresults shown are inshown Table 1in. TheTable results 1. The show results that show heat treatmentthat heat treatment improved improved the anti-corrosion the anti-corrosion property theof XK105,human makingbody. The it bettercurves than of XK105-F that of pureand XK105-H Mg. This show was due passivation to Mg Ca at beinga corrosion dissolved potential into theof property of XK105, making it better than that of pure Mg. This was due2 to Mg2Ca being dissolved lessmatrix than during −1.5 V. heat Passivation treatment, was and caused the e ffbyect the of added galvanic Ca, corrosion which formed becoming a stable inactive, passivation which reducedlayer to restraininto the the matrix corrosion during compared heat treatment, to pure and Mg the[36]. effe ct of galvanic corrosion becoming inactive, which thereduced corrosion the corrosion current and current potential. and potential. The potentiodynamic The potentiodynamic polarization polarization curve is shown curve inis Figureshown 14 in. The reduction rate of XK105-F is lower than that of pure Mg. XK105-H had the lowest reduction Figure 14. The reductionTable rate 1. Electrochemicalof XK105-F is propertieslower than of XK105that of and pure pure Mg. Mg. XK105-H had the lowest rate. The low reduction rate represents the low formation rate of H , which is harmless to the human reduction rate. The low reduction rate represents the low formation2 rate of H2, which is harmless to body.the human The curvesbody. The of XK105-F curvesSample of and XK105-F XK105-HPure and Mg show XK105-H XK105-F passivation show XK105-Hpassivation at a corrosion at a potentialcorrosion ofpotential less than of 1.5 V. Passivation was causedEcorr by(V) the added−1.601 Ca, which−1.971 formed − a1.786 stable passivation layer to restrain −less than −1.5 V. Passivation was caused by the added Ca, which formed a stable passivation layer to therestrain corrosion the corrosion compared compared to pureIcorr (A) Mg to pure [36 3.69]. Mg × 10 [36].−5 1.13 × 10−4 3.85 × 10−6

Table 1. Electrochemical properties of XK105 and pure Mg.

Sample Pure Mg XK105-F XK105-H Ecorr (V) −1.601 −1.971 −1.786 Icorr (A) 3.69 × 10−5 1.13 × 10−4 3.85 × 10−6

FigureFigure 14. 14. PotentiodynamicPotentiodynamic polarization polarization curve curve of of XK105. XK105. (a (a) )Comparison Comparison of of XK105-F, XK105-F, XK105-H, XK105-H, and and

purepure Mg Mg and and ( (bb)) comparison comparison of of XK105-H XK105-H before before and and after after erosion erosion test test by by SiO SiO22 particlesparticles with with erosion erosion angleangle of of 90°. 90◦ .

Figure 14. Potentiodynamic polarization curve of XK105. (a) Comparison of XK105-F, XK105-H, and

pure Mg and (b) comparison of XK105-H before and after erosion test by SiO2 particles with erosion angle of 90°. Metals 2019, 9, 857 11 of 15

Table 1. Electrochemical properties of XK105 and pure Mg.

Sample Pure Mg XK105-F XK105-H

Ecorr (V) 1.601 1.971 1.786 − 5 − 4 − 6 Icorr (A) 3.69 10− 1.13 10− 3.85 10− Metals 2019, 9, x FOR PEER REVIEW × × × 10 of 14

For surfacesurface roughening, roughening, 90 ◦90°SiO 2SiOerosion2 erosion created created higher roughness.higher roughness. The electrochemical The electrochemical properties propertiesof surface-roughened of surface-roughened XK105 are XK105 shown are in shown Table2 .in The Table corrosion 2. The corrosion current was current higher was after higher surface after surfaceroughening roughening treatment. treatment. The surface The rougheningsurface rougheni increasedng increased the surface the area surface of the area material, of the increasing material, increasingthe corrosion the rate corrosion [37]. The rate surface-roughened [37]. The surface-roughened XK105 did not XK105 show did passivation not showaround passivation1.5 Varound in the − −potentiodynamic1.5V in the potentiodynamic polarization curvepolarization because curve the rough because surface the maderough the surface passivation made layerthe passivation easily peel layeroﬀ after easily formation. peel off after formation.

TableTable 2. ElectrochemicalElectrochemical properties properties of heat treated and surface-roughened XK105.

SampleSample XK105-H XK105-H SiO SiO2-90°2-90 ◦

EcorrEcorr(V) (V) −1.7861.786 −1.4851.485 − 6 − 4 Icorr (A) 3.85 10 −6 5.85 10−4 Icorr (A) 3.85 × 10− 5.85 × ×10 −

3.6. Immersion Immersion Test Test To understand the surface corrosion mechanism and and quantify quantify the the corrosion corrosion rate rate of of materials, materials, thisthis study study subjected subjected XK105-H XK105-H and and surface-roughened XK105 to an immersion test. The The results results of of the the immersionimmersion testtest are are shown shown in Figurein Figure 15. The15. corrosionThe corro ratesion after rate surface after rougheningsurface roughening treatment treatment increased increasedgreatly, and greatly, the corrosion and the corrosion rate of XK105-H rate of XK105- was lowerH was than lower that than of pure that Mg.of pure The Mg. corrosion The corrosion rate of rateXK105-H of XK105-H after three after days three decreased days decreased and became and became stable. Thisstable. was This due was to thedue formation to the formation of a surface of a surfacepassivation passivation layer, which layer, restrained which restrained the corrosion the ecorrosionﬀect. The phenomenoneffect. The phenomenon in which the corrosionin which ratethe corrosiondecreased rate after decreased three days after of surface-roughened three days of surface-roughened XK105 was not obvious XK105 was relatively, not obvious because relatively, it did not becausehave a passivation it did not have layer. a passivation layer.

Figure 15. Corrosion rate in immersion test (a) XK105-H and pure Mg and (b) surface- Figure 15. Corrosion rate in immersion test (a) XK105-H and pure Mg and (b) surface-roughened XK105-H. roughened XK105-H.

Figure 16 shows the secondary surface of XK105-HXK105-H and surface-roughened XK105 during the immersionimmersion test.test. XK105-HXK105-H maintained maintained the the integrity integrity of its of macroscopic its macroscopic morphology, morphology, whereas whereas the surface- the surface-roughened roughened XK105 showed XK105 pitting showed corrosion pitting corrosion after one day after of one immersion. day of immersion. The main corrosion The main mechanism corrosion mechanismof Mg alloy of was Mg pitting alloy was corrosion pitting when corrosion immersed when inimmersed R-SBF [30 in]. R-SBF The surface [30]. The roughening surface roughening treatment treatmentcreated a lotcreated of pits a lot on of the pits surface on the of surface the material. of the Thesematerial. pits These promoted pits promoted pitting corrosion, pitting corrosion, and thus andweight thus loss weight due toloss corrosion due to corrosion was observed was observed on the 7th on day the of 7th immersion. day of immersion. Metals 2019,, 9,, 857x FOR PEER REVIEW 1211 of 1514 Metals 2019, 9, x FOR PEER REVIEW 11 of 14

Figure 16. Cross-sectional observation of roughened surface during immersion test. FigureFigure 16. 16. Cross-sectionalCross-sectional observation observation of of roughe roughenedned surface surface during during immersion immersion test. test. In Figure 17a, a calcium phosphate layer can be observed on the XK105-H surface after In Figure 17a, a calcium phosphate layer can be observed on the XK105-H surface after immersion.In Figure This 17a, layer a calcium protected phosphate the material layer can against be observed corrosion. on the Calcium XK105-H phosphate surface after is immersion. the main immersion. This layer protected the material against corrosion. Calcium phosphate is the main Thiscomponent layer protected of human the bone, material so this against layer could corrosion. help he Calciumal the affected phosphate area [38]. is the In main contrast, component there was of component of human bone, so this layer could help heal the affected area [38]. In contrast, there was humana thick cover bone, layer so this on layer the surface could helpof the heal surface-ro the aﬀectedughened area XK105, [38]. In as contrast, shown in there Figure was 17b. a thick The cover a thick cover layer on the surface of the surface-roughened XK105, as shown in Figure 17b. The cover layercompound on the was surface Mg ofhydroxide the surface-roughened which does not XK105, have as a shown healing in effect. Figure The17b. formation The cover compoundof calcium compound was Mg hydroxide which does not have a healing effect. The formation of calcium wasphosphate Mg hydroxide requires which a suitable does not hy havedroxide a healing ion group eﬀect. to The react formation with [39]. of calcium The degradation phosphate requiresrate of phosphate requires a suitable hydroxide ion group to react with [39]. The degradation rate of asurface- suitable roughened hydroxide XK105 ion group was to too react fast with for [39 calcium]. The degradationphosphate rateto form; of surface- instead, roughened Mg hydroxide XK105 surface- roughened XK105 was too fast for calcium phosphate to form; instead, Mg hydroxide wasformed. too fast for calcium phosphate to form; instead, Mg hydroxide formed. formed.

Figure 17.17. SEMSEM images images and and EDS EDS analysis analysis of immersed of immersed surface. (surface.a) XK105-H (a) andXK105-H (b) surface- and Figure 17. SEM images and EDS analysis of immersed surface. (a) XK105-H and roughened(b) surface- XK105-H.roughened XK105-H. (b) surface- roughened XK105-H. Metals 2019, 9, 857 13 of 15 Metals 2019, 9, x FOR PEER REVIEW 12 of 14

To summarize the the above above results, results, Mg-1Ca-0.5Zr Mg-1Ca-0.5Zr biomedical biomedical alloy alloy has has a suitable a suitable degradation degradation rate rateand andgood good mechanical mechanical strength. strength. The The surface surface roughness roughness can can help help control control the the degradation degradation rate. Mg-1Ca-0.5ZrMg- 1Ca- 0.5Zr can can thus thus be be applied applied for for biomedical biomedical implants. implants. A A summary summary is is shown shown in in Figure Figure 18 18..

Figure 18. Summary of Mg-1Ca-0.5Zr biomedical implants.

4. Conclusions

(1)(1) TheThe brittle brittle Mg Mg2Ca2Ca phase phase of ofthe the XK105 XK105 solid solid dissolved dissolved into into the matrix the matrix after afterheat heattreatment. treatment. This madeThis made the texture the texture uniform uniform and and increased increased both both elongation elongation and and fracture fracture toughness. toughness. The Young’s modulusmodulus of of XK105 increased after heat treatment because of the solid solution of Ca atoms. (2)(2) TheThe TD TD plane of XK105 had higher strength but lower elongation than those of the ED planeplane becausebecause of of the the texture texture effects effects in in which which the the grain grain slip slip was was hard hard to start to start up on up the on ED the plane. ED plane. The fractureThe fracture toughness toughness along along ED EDwas was better better than than that that along along TD TD because because the crack-spreadingcrack-spreading directiondirection was parallel parallel to the arrangem arrangementent direction of the secondary phase. (3)(3) TheThe heat treatment of XK105 reduced galvanic corrosion due to the solid solution, and further improvedimproved anti-corrosion anti-corrosion ability. The The surface-roughened surface-roughened XK105 XK105 had a larger surface area, which enhancedenhanced the the corrosion corrosion reaction. (4)(4) AA calcium calcium phosphate layer layer formed during corrosi corrosionon on the XK105-H surface,surface, significantlysignificantly reducingreducing thethe corrosion corrosion rate. rate. The The surface-roughened surface-roughened XK105 XK105 formed formed magnesium magnesium hydroxide hydroxide instead insteadbecause because of its faster of corrosionits faster rate.corrosion The magnesium rate. The hydroxidemagnesium did hydroxide not provide did aprotective not provide effect. a (5) protectiveThis study effect. developed a biomedical implant material, Mg-1Ca-0.5Zr alloy. The results can be used (5) Thisas a referencestudy developed for the biomedicala biomedical material implant industry. material, Mg-1Ca-0.5Zr alloy. The results can be used as a reference for the biomedical material industry.

Author Contributions: Y.-T.C.: conceived and designed the analysis, collected the data, performed the analysis, wroteAuthor the Contributions: paper. F.-Y.H.: Y.-T.C.: conceived conceived and designed and design the analysis,ed the analysis, contributed collected data andthe data, analysis performed tools, performed the analysis, the analysis,wrote the revised paper. theF.-Y.H.: paper. conceived J.-C.S.: collected and designed the data, the performed analysis, contributed the analysis, data wrote and the analysis paper. tools, performed the analysis, revised the paper. J.-C.S.: collected the data, performed the analysis, wrote the paper. Funding: The authors are grateful to the Instrument Center of National Cheng Kung University and the Ministry ofFunding: Science The and authors Technology are grateful of the Republicto the Instrument of China Center (MOST of 107-2221-E-006-012-MY2)National Cheng Kung University for their and financial the Ministry support. of ConflictsScience and of Technology Interest: The of authorsthe Republic declare of China no conflicts (MOST of 107-2221-E-006-012-MY2) interest. for their financial support. Conflicts of Interest: The authors declare no conflicts of interest. References

1.ReferencesStaiger, M.P.; Pietak, A.M.; Huadmai, J.; Dias, G. Magnesium and its alloys as orthopedic biomaterials: 1. AStaiger, review. M.P.;Biomaterials Pietak, A.M.;2006, Huadmai,27, 1728–1734. J.; Dias, [CrossRef G. Ma][gnesiumPubMed and] its alloys as orthopedic biomaterials: A 2. Dikici,review. B.;Biomaterials Topuz, M. 2006 Production, 27, 1728–1734, of annealed doi:10.1016/j.biomaterials.2005.10.003. cold-Sprayed 316L stainless steel coatings for biomedical 2. applicationsDikici, B.; Topuz, and their M. in-vitroProduction corrosion of annealed response. cold-SprayedProt. Met. Phys. 316L Chem. stainless Surf. steel2018 coatings, 54, 333–339. for biomedical [CrossRef] 3. Yadroitsev,applications I.; and Karkhmalev, their in-vitro P.; Yadroitsava, corrosion I.; response. Du Plessis, Prot. A. Qualiﬁcation Met. Phys. ofChem. Ti6Al4V Surf ELI. 2018 alloy, 54, produced 333–339, by laserdoi:10.1134/S2070205118020168. powder bed fusion for biomedical applications. JOM 2018, 70, 372–377. [CrossRef] Metals 2019, 9, 857 14 of 15

4. Chen, Q.Z.; Thouas, G.A. Metallic implant biomaterials. Mater. Sci. Eng. R Rep. 2015, 87, 1–57. [CrossRef] 5. Chien, C.; Liao, T.Y.; Hong, T.F.; Kuo, T.Y.; Chang, C.H.; Yeh, M.L.; Lee, T.M. Surface microstructure and bioactivity of hydroxyapatite and fluorapatite coatings deposited on Ti-6Al-4V substrates using Nd-YAG Laser. J. Med. Biol. Eng. 2014, 34, 109–115. [CrossRef] 6. Hench, L.L.; Polak, J.M. Third-generation biomedical materials. Science 2002, 295, 1014–1017. [CrossRef] [PubMed] 7. Kojima, Y. Platform science and technology for advanced magnesium alloys. Mater. Sci. Forum 2000, 350, 3–18. [CrossRef] 8. Musso, C.G. Magnesium metabolism in health and disease. Int. Urol. Nephrol. 2009, 41, 357–362. [CrossRef] 9. Nagels, J.; Stokdijk, M.; Rozing, P.M. Stress shielding and bone resorption in shoulder. J. Shoulder Elb. Surg. 2003, 12, 35–39. [CrossRef] 10. Jurgen, V. Magnesium: Nutrition and metabolism. Mol. Asp. Med. 2003, 24, 27–37. [CrossRef] 11. Lin, D.J.; Hung, F.Y.; Jakfar, S.; Yeh, M.L. Tailored coating chemistry and interfacial properties for construction of bioactive ceramic coatings on magnesium biomaterial. Mater. Des. 2016, 89, 235–244. [CrossRef] 12. Witte, F.; Kaese, V.; Haferkamp, H.; Switzer, E.; Meyer-Lindenberg, A.; Wirth, C.J.; Windhagen, H. In vivo corrosion of four magnesium alloys and the associated bone response. Biomaterials 2005, 26, 3557–3563. [CrossRef][PubMed] 13. Pietak, A.; Mahoney, P.; Dias, G.J.; Staiger, M.P. Bone-like matrix formation on magnesium and magnesium alloys. J. Mater. Sci. Mater. Med. 2008, 19, 407–415. [CrossRef][PubMed] 14. Witte, F.; Hort, N.; Vogt, C.; Cohen, S.; Kainer, K.U.; Willumeit, R.; Feyerabend, F. Degradable biomaterials based on magnesium corrosion. Curr. Opin. Solid State Mater. Sci. 2008, 12, 63–72. [CrossRef] 15. Hong, Y.S.; Yang, K.; Zhang, G.D.; Huang, J.J.; Hao, Y.Q.; Ai, H.J. The role of Bone induction of a biodegradable magnesium alloy. Acta Metall. Sin. 2008, 44, 1035–1041. 16. Radha, R.; Sreekanth, D. Insight of magnesium alloys and composites for orthopedic implant applications—A review. J. Magnes. Alloy. 2017, 5, 286–312. [CrossRef] 17. Harandi, S.E.; Idris, M.H.; Jafari, H. Effect of forging process on microstructure, mechanical and corrosion properties of biodegradable Mg–1Ca alloy. Mater. Des. 2011, 32, 2596–2603. [CrossRef] 18. Du, H.; Wei, Z.J.; Liu, X.W.; Zhang, E.L. Effects of Zn on the microstructure, mechanical property and bio-corrosion property of Mg–3Ca alloys for biomedical application. Mater. Chem. Phys. 2011, 125, 568–575. [CrossRef] 19. Kannan, M.B.; Raman, R.K.S. In vitro degradation and mechanical integrity of calcium-containing magnesium alloys in modified-simulated body fluid. Biomaterials 2008, 29, 2306–2314. [CrossRef] 20. Lin, D.J.; Hung, F.Y.; Lui, T.S.; Yeh, M.L. Heat treatment mechanism and biodegradable characteristics of ZAX1330 Mg alloy. Mater. Sci. Eng. C 2015, 51, 300–308. [CrossRef] 21. Jeong, Y.S.; Kim, W.J. Enhancement of mechanical properties and corrosion resistance of Mg–Ca alloys through microstructural refinement by indirect extrusion. Corros. Sci. 2014, 82, 392–403. [CrossRef] 22. Kim, W.C.; Kim, J.G.; Lee, J.Y.; Seok, H.K. Influence of Ca on the corrosion properties of magnesium for biomaterials. Mater. Lett. 2008, 62, 4146–4148. [CrossRef] 23. Zheng, Y.F.; Gu, X.N.; Xi, Y.L.; Chai, D.L. In vitro degradation and cytotoxicity of Mg/Ca composites produced by powder metallurgy. Acta Biomater. 2010, 6, 1783–1791. [CrossRef][PubMed] 24. Qian, M.; StJohn, D.H.; Frost, M.T. Heterogeneous nuclei size in magnesium–zirconium alloys. Scr. Mater. 2004, 50, 1115–1119. [CrossRef] 25. Liu, G.L. The electronic structure of the microstructure of Mg-Zr alloys. Acta Phys. Sin. 2008, 57, 1043–1047. 26. Zhou, Y.L.; Li, Y.; Luo, D.M.; Wen, C.; Hodgson, P. Microstructures, mechanical properties and in vitro corrosion behaviour of biodegradable Mg–Zr–Ca alloys. J. Mater. Sci. 2012, 48, 1632–1639. [CrossRef] 27. Tamura, Y.; Sugimoto, Y.; Soda, H. Structure and mechanical properties of Mg–Ca and Mg–Ca–Zr alloys. J. Jpn. Inst. Light Met. 2013, 63, 279–285. [CrossRef] 28. Zhong, C.; Liu, F.; Wu, Y.T.; Le, J.J.; Liu, L.; He, M.F.; Zhu, J.C.; Hu, W.B. Protective diffusion coatings on magnesium alloys: A review of recent developments. J. Alloys Compd. 2012, 520, 11–21. [CrossRef] 29. Ali, Y.H.; Qiu, D.; Jiang, B.; Pan, F.S.; Zhang, M.X. Current research progress in grain refinement of cast magnesium alloys: A review article. J. Alloys Compd. 2015, 619, 639–651. [CrossRef] 30. Oyane, A.; Kim, H.M.; Furuya, T.; Kokubo, T.; Miyazaki, T.; Nakamura, T. Preparation and assessment of revised simulated body fluids. J. Biomed. Mater. Res. Part A 2003, 65, 188–195. [CrossRef] Metals 2019, 9, 857 15 of 15

31. Zeng, R.C.; Qi, W.C.; Cui, H.Z.; Zhang, F.; Li, S.Q.; Han, E.H. In vitro corrosion of as-extruded Mg–Ca alloys—The influence of Ca concentration. Corros. Sci. 2015, 96, 23–31. [CrossRef] 32. Wang, C.Q.; Sun, M.; Zheng, F.Y.; Peng, L.M.; Ding, W.J. Improvement in grain refinement efficiency of Mg–Zr master alloy for magnesium alloy by friction stir processing. J. Magnes. Alloys 2014, 2, 239–244. [CrossRef] 33. Zeng, Z.R.; Bian, M.Z.; Xu, S.W.; Davies, C.H.J.; Birbilis, N.; Nie, J.F. Effects of dilute additions of Zn and Ca on ductility of magnesium alloy sheet. Mater. Sci. Eng. A 2016, 674, 459–471. [CrossRef] 34. Gu, X.N.; Li, S.S.; Li, X.M.; Fan, Y.B. Magnesium based degradable biomaterials: A review. Front. Mater. Sci. 2014, 8, 200–218. [CrossRef] 35. Yang, G.L.; He, F.M.; Yang, X.F.; Wang, X.X.; Zhao, S.F. Bone responses to titanium implants surface-roughened by sandblasted and double etched treatments in a rabbit model. Oral Surg. Oral Med. Oral Pathol. Oral Radiol. Endodontol. 2008, 106, 516–524. [CrossRef][PubMed] 36. Hagihara, K.; Fujii, K.; Matsugaki, A.; Nakano, T. Possibility of Mg-and Ca-based intermetallic compounds as new biodegradable implant materials. Mater. Sci. Eng. C 2013, 33, 4101–4111. [CrossRef][PubMed] 37. Von Der Hoh, N.; Bormann, D.; Lucas, A.; Denkena, B.; Hackenbroich, C.; Meyer-Lindenberg, A. Influence of different surface machining treatments of magnesium-based resorbable implants on the degradation behavior in rabbits. Adv. Eng. Mater. 2009, 11, 47–54. [CrossRef] 38. Song, Y.W.; Shan, D.Y.; Han, E.H. Electrodeposition of hydroxyapatite coating on AZ91D magnesium alloy for biomaterial application. Mater. Lett. 2008, 62, 3276–3279. [CrossRef] 39. Zhang, C.Y.; Gao, J.C.; Zeng, R.C.; Liu, C.L.; Wu, X.; Wu, D. Corrosion behavior of magnesium alloy AZ31 with calcium phosphate coating in hank’s solition. J. Chin. Ceram. Soc. 2010, 38, 885–891.

© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).