Materials Science and Engineering C 33 (2013) 618–625
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Materials Science and Engineering C
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Microstructure, mechanical property and corrosion behaviors of interpenetratingC/Mg-Zn-Mn composite fabricated by suction casting
X. Wang a,⁎, L.H. Dong a, X.L. Ma b, Y.F. Zheng a,c
a Center for Biomedical Materials and Engineering, Harbin Engineering University, Harbin 150001, Chinab School of Materials Science and Engineering, Harbin University of Science and Technology, Harbin 150080, Chinac State Key Laboratory for Turbulence and Complex System and Department of Advanced Materials and Nanotechnology, College of Engineering, Peking University, Beijing 100871, China
A novel interpenetrating C/Mg-Zn-Mn composite was fabricated by infiltrating Mg-Zn-Mn alloy into porouscarbon using suction casting technique. The microstructure, mechanical properties and corrosion behaviorsof the composite have been evaluated by means of SEM, XRD, mechanical testing and immersion test. Itwas shown that the composite had a compact structure and the interfacial bonding between Mg-Zn-Mnalloy and carbon scaffold was very well. The composite had an ultimate compressive strength of (195±15)MPa, which is near with the natural bone (2–180 MPa) and about 150-fold higher than that of the originalporous carbon scaffold, and it still retained half of the strength of the bulk Mg-Zn-Mn alloy. The corrosion testindicated that the mass loss percentage of the composite was 52.9% after 30 days′ immersion in simulatedbody fluid (SBF) at 37±0.5 °C, and the corrosion rates were 0.043 mg/cm2h and 0.028 mg/cm2h after 3 and7 days′ immersion, respectively. The corrosion products on the composite surface were mainly Mg(OH)2 andhydroxyapatite (HA).
Carbon, in all its forms, such as carbon nanotubes, carbon fabric,carbon-carbon composites, glassy carbon, pyrocarbons and diamond-like layers, is considered as a promising material for biomedical applica-tion thanks to its interesting mechanical, physical, chemical propertiesand to its biocompatibility with living bone and other tissue [1]. In addi-tion to the advantages of carbon materials, porous carbon fabricated byusing polyurethane foam template has many special features, such aslight weight (0.2-0.8 g/cm3), large external surface area with open cellstructure and adjustable thermal and electrical conductivity [2]. Mean-while, the porous carbon is useful as a template for new tissue growthand can provide temporary structural support as a delivery vehicle forcells and/or bioactive molecules [3, 4]. Lewandowska-Szumiel et al. [5]found that porous carbon fiber reinforced carbon (CFRC) compositewas partially degraded when it was in contact with bone and appearedsubstantially tissue compatible. Turgut et al. [6] also prepared porouscarbon with controllable pore size and distribution with high ratio ofopen porosity for the purpose of using it as a bone implant material.Histological examination of the engineered constructs revealed that thetissue adaptation and bone compatibility of the carbon foam werefound to be satisfactory. Progression of connective tissue formation intothe carbon implant was observed without any sign of cytotoxicity andincompatibility during the postoperative follow-ups. However, the low
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strength and brittleness of porous carbon limits its clinical applicationin repairing bone defect. Metals have played an important role as ortho-pedic biomaterials owing to their strength and resilience [7], so thebiocomposite materials can be designed by combining with alloys toimprove the mechanical properties of porous carbon.
Mg alloys show potential application as degradable hard tissueimplant materials due to their good biocompatibility and mechanicalproperties, in comparison with biodegradable polymeric implantmate-rials currently in use [8–10]. Meanwhile, Mg alloys have the advantagethat they do not show the undesirable effect of stress shielding, whichoccurs normally in conventional metallic implants made of stainlesssteel or titanium alloy which are widely used for bone implants[11–13]. Yin [14] reported that the tensile strength, the elongationand the yield strength of the as-cast Mg-Zn-Mn alloys increased withthe increasing of Zn content, and the highest mechanical property wasobtained at 3.0 wt.% Zn. When Zn content of the as-cast Mg-Zn-Mn al-loys increased to 3.0 wt.%, the corrosion resistance property apparentlydecreased. Xu et al. [15] have also fabricated Mg-Mn-Zn alloys withgood mechanical and corrosion resistance properties. Meanwhile, invitro degradation, hemolysis andMC3T3-E1 cell adhesion of biodegrad-ableMg-6wt.%Zn alloywere researched. The alloy zinc elevated the cor-rosion potential and charge transfer resistance of Mg, and the in vitrodegradation rate ofMg-6wt.%Zn alloywas lower than that of high purityMg in SBF. Hemolysis and adhesion of cells display good biocompatibil-ity ofMg-Zn alloy [12]. However, the rapid degradation rate ofMg alloysin human bio-environment limits their clinical applications [16, 17].Therefore, it is important to improve the corrosion resistance of Mg
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alloys in order to enlarge their applicationfield. Recently,manymethods,such as element alloying [14, 15, 18], surface modification [19, 20] anddesign of composite materials [21–24], have been used to improve thecorrosion resistance of Mg alloys. Although the element alloying canincrease the bio-corrosion resistance of Mg alloys to some extent, itcannot significantly improve the bone response to Mg alloys, especiallythe bone response at the early stage [15]. Some coating can providethe Mg alloy with a significantly better surface cytocompatibility andimproved osteoconductivity and osteogenesis in the early postoperationperiod, but combination of coating and matrix alloy is poor [19, 20].Biocomposite materials can be designed to obtain a wide range of me-chanical and biological properties [25]. Gu et al. [21] prepared Mg/HAmetal matrix composite by powder metallurgy method and found thecorrosion rate of Mg/HA composites increased with the increment ofHA content and Mg/10HA extract showed no toxicity to L-929 cells.Magnesium fluorapatite nanocomposite (AZ91-20FA) was also madevia the blending-pressing-sintering method. The addition of FA nanoparticles to magnesium alloy can reduce the corrosion rate and acceler-ate the formation of bone-like apatite layer and in turn provide betterprotection for matrix alloy [26]. Among the techniques currently avail-able for the processing of metal matrix composites, the properties ofthe composite fabricated by liquid metal infiltration technique (a tech-nique of infiltration of a molten metal into a porous ceramic preform)can be adjusted by free selection of different porous scaffolds as well asthe molten metals [27, 28]. In our previous study, the corrosion rate ofthe MgCa-HA/TCP composite fabricated by liquid metal infiltration wasslower than that of the bulk MgCa alloy alone. And the diluted extractswith 50 and 10% concentrations of the resulting composite exhibitedgrade I cytotoxicity to L-929 andMG63 cells [29]. However, few attemptshave been made so far to combine porous carbon with Mg alloys toprepare a new biomedical material.
This paper aims to combine the bothmerits by preparingMg-Zn-Mnalloys and porous carbon composite in order tomake up for each other'sdeficiencies of properties. Firstly, porous carbon is fabricated by carbon-ize polyurethane foam modified by thermoset phenolic resin athigh temperature, and then biphase continuation interpenetratingcomposite are fabricated by infiltrating the molten Mg-Zn-Mn alloyinto the porous scaffold using self-made suction casting equipment.The microstructure, mechanical properties, in vitro corrosion behaviorof C/ Mg-Zn-Mn composite were characterized to evaluate the feasibil-ity of the resulting composite as potential biomaterials.
2. Materials and methods
2.1. Materials and preparation
The interconnected porous carbon was prepared by carbonizingphenolic resin coated commercial thermoplastic polyurethane sponges(PU, 40 ppi). PU sponges were dipped into the tetrachloroethylenesolution for about 30 minutes and then washed with distilled waterbefore drying at 40 °C for at least 24 h. The phenolic resin was used asmodifier for PU sponges. The cleaned PU sponges were dipped intothe phenolic resin slurry and gently squeezed to allow the slurry toadhere to the sponges before drying at 150 °C for at least 3 h. Finally,the phenolic resin slurry-coated PU sponges were carbonized at900 °C for 2 hours in a vacuum heating furnace and the final porouscarbon were removed from the furnace after it had cooled down toambient temperature. The heating rate of vacuum heating furnacewas of 5 °C/min applied up to 900 °C.
The Mg-3 wt.%Zn-1 wt.%Mn (expressed as Mg-Zn-Mn alloy in thefollowing) were melted by using commercial pure Mg (≥99.99%),pure Zn (≥99.99%), and Mg-10Mn master alloy in a high-puritygraphite crucible by resistance furnace under a mixed gas atmosphereof SF6 and CO2. The C/Mg-Zn-Mn composite were fabricated byinfiltrating the molten Mg-Zn-Mn alloy into the porous scaffoldusing self-made vacuum suction casting equipment, and the diagram
of experimental equipment was shown in Fig. 1. During the fabrica-tion process, the porous carbon was preheated at 150 °C andpositioned in the vacuum suction casting equipment, which waspre-pumped to −0.07~−0.05 MPa. The matrix Mg-Zn-Mn alloywas melted in the furnace under a mixed gas atmosphere of SF6 andCO2 and the temperature was kept at 700~720 °C. Then the infiltra-tion of the molten Mg-Zn-Mn alloy was driven by the vacuum andheld for 2 min while the melt was solidified. Finally, the C/Mg-Zn-Mncomposite was removed from the suction casting equipment after ithad cooled down to ambient temperature.
For the compression and immersion tests, the rectangularcomposite samples (with the size of 5 mm×5 mm×8 mm) wereground and mechanically polished up to 2000 grit, then ultrasonicallycleaned in acetone, absolute ethanol and distilled water and thendried in air.
2.2. Microstructure characterization
Scanning electron microscopy (SEM) equipped with energy dis-perse spectrometer (EDS) was employed for the microstructure iden-tification of the sintered porous carbon and changes on the surface ofthe C/Mg-Zn-Mn composite samples before and after immersion. Andtheir chemical compositions were also analyzed by EDS. The phases ofsintered porous carbon, as-cast composite samples and their corro-sion products were identified by the X-ray diffractometer (XRD)using the Cu Kα radiation at the step size of 0.02° with a scanningspeed of 2° min-1.
2.3. Porosity measurement
The porosity of porous carbon scaffold was measured usingArchimedes Principle. At first, each scaffold was submerged indeionized water, degassed under vacuum, and suspended from ananalytical balance to obtain scaffold wet weight (mwet). Then thescaffolds were dried in the oven with air blower at room tempera-ture overnight. The dimensions of each specimen were further mea-sured using a Vernier callipers to produce a total volume (V) of thescaffold, and the porous scaffold dry weight (mdry) was also deter-mined with an analytical balance. All weights were in grams. Thus,the porosity (θ) was calculated using the formula [30]:
θ ¼ 100 1−mdry−mwet
ρ⋅V
� �ð1Þ
in which ρ is the density of deionized water( ~1 g/cm3). Five parallelspecimens were taken for every scaffold and the mean value of theporosity was achieved.
2.4. Compression testing
Uniaxial compression testing was conducted on an Instron 3365testing machine at a constant nominal strain rate of 7×10-4 s-1 at roomtemperature. The test samples with the size of 3 mm×3 mm×6 mmwere prepared according to the ASTM-E9-09 [31]. Five identical sampleswere used for the compressive tests.
2.5. Electrochemical test
A conventional three electrode cell was used: the composite sam-ple (exposed area of 1 cm2) was the working electrode, an Ag/AgClelectrode was used as reference and a platinum foil was used ascounter electrode. The electrochemical tests were carried out at37±0.5 °C in SBF solution using an electrochemical workstation(SI1287). SBF solution was prepared by dissolving the reagentsNaCl, NaHCO3, KCl, K2HPO4 · 3H2O, MgCl2 · 6H2O, CaCl2 · 2H2O andNa2SO4 into distilled water. The solution was buffered to pH 7.4 at
Fig. 1. Diagram of experimental equipment of vacuum suction casting.
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37±0.5 °C with HCl and Tris (hydroxy-methyl) aminomethane. Con-centration (in mmol/L) of various ions in the SBF was: 142 Na+, 5 K+,1.0 Mg2+, 2.5 Ca2+, 109Cl-, 27 HCO3
-, 1.0 HPO42-, 1.0 SO4
2-. The polari-zation curve was measured at a scanning rate of 1 mV s-1 in thepotentiodynamic polarization tests. The volume of the SBF used inthe electrochemical test was approximately 150 mL.
2.6. Immersion test
The immersion test was carried out in SBF solution at 37±0.5 °Caccording to ASTM-G31-72 [32]. The test samples with the size of5 mm×5 mm×8 mmwere prepared. After 3, 7, 15 and 30 days’ immer-sion, the samples were removed out of the solution, gently rinsed withdistilled water, and dried in air. Microstructure and composition charac-terizations of the corroded composite sample surface were examined bySEM, EDS and XRD. It should be noted that the immersion solution of15 days and 30 days were refreshed every other 2 days, in order tomaintain the same composition and pH value as the human bodyfluid. Moreover, the corrosion mass loss was calculated from the massloss of the specimen weighed by an analysis balance with the sensitivityof ±0.01 mg. The mass loss percentage was calculated according tothe formula: mass loss percentage=(mbefore immersion-mafter immersion)/mbefore immersion, where m means mass. The pH meter (E-900, luosu,Shanghai) was used to record the change of pH value of the SBF solutionduring the soaking experiment. An average of three measurements wastaken for each group.
The corrosion rate of the composite was calculated according tothe Ref. [33] using the formula: CR=m/St, where CR is the corrosionrate (in mg/cm2h), m is the mass loss (in mg), S is the exposedsurface area of the specimen (in cm2), t is exposed time (in h).
2.7. Statistical analysis
Statistical analysis was conducted to evaluate the difference incorrosion rate by the analysis of variance (ANOVA). The statisticalsignificance was defined as 0.05.
3. Results and discussion
3.1. Microstructures of porous carbon and composite
The microstructure of carbonized carbon scaffold, with the totalporosity 92%, was shown in Fig. 2(a). It revealed that the pore ofporous carbon prepared by PU sponge template was open with sizeabout 460–750 μm and the cell geometry was remarkably uniform.
It was evident that there were no micro and mesopore structure ex-cept hollow structure formed due to the decomposition of PU spongeduring the pyrolysis process (inset in Fig. 2(a)) found in the surface ofporous carbon. In addition, no cracks appeared in the junction area ofthe porous cell. Thus the resulting intercontinuous porous carbonwould have better mechanical strength. The complete infiltration ofthe molten Mg-Zn-Mn alloy into the carbon scaffold was successfullyachieved and the typical SEM image of the resulting composite wasshown in Fig. 2(b), the brighter phase is Mg-Zn-Mn and the darkeris carbon. Compared with the microstructure of porous carbon, itcan be found that the composite had a closed-cell and compact struc-ture as well as the interfacial bonding between Mg-Zn-Mn alloy andcarbon scaffold was very well without discernible debonding ormicro-crack. This compact structure also demonstrated the efficientinfiltration of the Mg-Zn-Mn alloy throughout the porous carbon. Inaddition, it can be seen that it was also a typical casting microstruc-ture containing small grains, and the grain boundary could be seenin the alloy. The second phase was mainly distributed in the grainboundary and the interdendrite of matrix alloy. Zhang et al. [11]demonstrated that this second phase was Mg7Zn3. When Chang etal. [34] studied the microstructures of Al(Mg)/Al2O3 interpenetratingcomposites produced by a pressureless infiltration technique, theyalso found that grain boundaries could be seen in the Al(Mg) alloy,and suggested that there was an upper size limit for the alloy grainwhen in particularly large ceramic foam cells.
Fig. 3 showed the XRD analysis results of the porous carbon,Mg-Zn-Mn alloy and C/Mg-Zn-Mn composite. The XRD patternsexhibited that the porous material fabricated in this present study wasvitreous carbon (as shown in Fig. 3(a)). Fig. 3(b) showed that the pre-dominant phases of the Mg-Zn-Mn alloy were α-Mg and the secondMg7Zn3 phase. According to the binary Mg-Zn diagram [35], it can beobserved that up to 2wt.%Zn completely dissolve in α-Mg matrix.Thus, with the increasing of Zn content, the second Mg7Zn3 phase wasfound in Mg-3Zn-1Mn alloy [11]. Meanwhile, in addition to the phasesofα-Mg andMg7Zn3, no other phase appeared in C/Mg-Zn-Mn compos-ite (as shown in Fig. 3(c)), and also no carbon peaks appeared, this isbecause the amount of carbon is too little to be detected.
3.2. Mechanical property of composite
Fig. 4 showed the representative compression stress strain curvesof the C/Mg-Zn-Mn composite. The average compressive strength andthe elongation for the resulting composite were (195±15) MPa and(20±1)%, which is near with the natural bone (2–180 MPa) andabout 150-fold higher than that of the original porous carbon scaffold
Fig. 2. SEM images of (a) the carbonized porous carbon and (b) C/Mg-Zn-Mn composite.
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(around 1.45 MPa) [inset in Fig. 4], but retained half of the strength ofthe bulk Mg-Zn-Mn alloy (around 370 MPa). That is, the C/Mg-Zn-Mncomposite shows a great improvement in the strength and theelongation in comparison with the original porous carbon scaffold.MgCa-HA/TCP composite were investigated in our previous study,that the crushing of ceramic foams is based on bending of the struts[29]. The interpenetration of the Mg-Zn-Mn alloy stabilizes theceramic struts and partially prevents strut bending, resulting in amajor improvement of the mechanical properties of the ceramicscaffold. Many factors influence the compressive strength of thecomposite, such as compactness, the interfacial bonding of matrixalloy and inforcement, volume fraction and distribution uniformityof inforcement. Among these, the compactness has the biggest impacton the compressive property of the fabricated composite. For theresulting C/Mg-Zn-Mn composite, the interfacial bonding is well (asshown in Fig. 2(b)), but the hollow structure of carbon scaffold[inset in Fig. 2(a)] reduced the composite's compactness to someextent, coupled with the brittle carbon being intercontinuous, thusleads to the reduced mechanical property compared to the matrixMg-Zn-Mn alloy. Zeschky et al. [36] thought when load transferredfrom the infiltrated alloy to ceramic skeleton, crack would occur be-tween the infiltrated metal and the ceramic skeleton during the com-pression process, so AZ91 infiltrated Si-O-C ceramic foams presentedreduced mechanical property compared to the AZ91 bulk alloy.For this kind of metal-ceramic interpenetrating composite, anotherstudy showed crack propagated preferentially through the brittle
Fig. 3. XRD patterns of the (a) porous carbon, (b) Mg-Zn-Mn alloy and (c) C/Mg-Zn-Mncomposite.
ceramic phase, when it encountered the ductile metal, the metaldeformed and the crack was bridged by the metal [34].
Table 1 summarizes the ultimate compressive strength of differentmaterials. It should be noted that the reported alloy including theMg-Zn and Mg-Ca and the Mg-3Zn-1Mn alloy as control in thisstudy demonstrated greatly higher ultimate compressive strengththan that of the natural bone. However, the values of the ultimatecompressive strength of some other porous materials such as porousMg, porous polymers and so on were too lower to be suitable fornatural bone.
3.3. In vitro immersion tests
3.3.1. Electrochemical measurementThe representative potentiodynamic polarization curve of
C/Mg-Zn-Mn composite in SBF at 37±0.5 °C was shown inFig. 5, compared with Mg-Zn-Mn alloy. The experimental compositeclearly indicated less negative corrosion potential value (−1.50 V)than the bulk Mg-Zn-Mn matrix alloy (−1.69 V), and there was nosignificant difference in the corrosion current density of the experimen-tal composite (14.47 μA cm-2) and the bulk Mg-Zn-Mn matrix alloy(13.88 μA cm-2). Also, there existed a longer passivation stage in thepolarization curve of composite than that of matrix alloy sample, indi-cating the better corrosion resistance of the C/Mg-Zn-Mn composite.Razavi et al. [26] also found that there existed a long passivation stagein the polarization curve of AZ91-20FA composite sample while there
Fig. 4. Compressive behaviors of three replicate C/Mg-Zn-Mn composite samples, withMg-Zn-Mn matrix alloy and carbon scaffold as controls.
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was no passivation stage in the curve of AZ91 sample, and thought thatthe addition of FA nano particles drops the cathodic current density andslows the rate of hydrogen evolution.
3.3.2. Surface morphology and phase composition of compositeFig. 6 illustrated the morphologies of the C/Mg-Zn-Mn composite
after 3, 7, 15 and 30 days′ immersion in SBF solution at 37±0.5 °C, re-spectively. The surface of the composite was rugged as well as coveredwith the similar spherical protective corrosion products after 3 and7 days′ immersion, and the scaffold structure was visible obviouslyafter 7 days′ immersion, as shown in Fig. 6(a, b). After 15 days ofimmersion, a compact corrosion product layer was clearly observedon the surface of the C/Mg-Zn-Mn composite [Fig. 6(c)]. However,many cracks on the surface of the infiltrated Mg-Zn-Mn alloy could beseen [Fig. 6(c, d)]. Some deep corrosion pits were also observed on thesurface of the infiltratedMg-Zn-Mn alloy [Fig. 6(d)], revealing the disso-lution of the Mg-Zn-Mn alloy. The different results were caused by thefilm change, in the SBF solution, the Mg(OH)2 is connected with someH2O molecule to form the hydrate form of Mg(OH)2·nH2O. When thesamples are dried in the air, the film shrinks due to dehydration. Then,lots of cracks are formed on the film surface [24].
The EDS results of the C/Mg-Zn-Mn composite surface after 15 days′immersion in SBF solution was shown in Fig. 7. It indicated that thecorrosion product were mainly composed of O, P, C, Mg, Ca and asmall amount of Na elements. The atomic ratio of Ca to P (Ca/P) wasabout 1.84, which is close to 1.67 (the Ca/P atomic ratio of HA) andsuggested the deposition of a compound containing Ca and P elements.TheCa/P atomic ratio of the calciumphosphate crystals depended on thematerial composition and the time of in vitro degradation and were inthe range of 1.52-2.0, while the Ca/P ratio in native cancellous bone inthe ilium was 1.83 and the Ca/P ratio in the regenerated cancellous
Fig. 5. Potentiodynamic polarization curves of the C/Mg-Zn-Mn composite in SBF, withthe Mg-Zn-Mn alloy as a control.
bone at 4 months was 1.76 [41]. The presence of the Ca-P coating wasconsistentwith the in vivo studies onMg-Zn-Mn alloy [42] and indicatedthat the composite had a good osteoconductivity and osteoinductivity[43]. Fig. 8 showed the XRD patterns of the composite samples afterimmersing in SBF solution for different immersion periods. Results re-vealed that the corrosion products were mainly magnesium hydroxide[Mg(OH)2] and hydroxyapatite (HA). The in vitro testing solution wasSBF, containing inorganic ions such as Cl-, H2PO4
- and Ca2+. Thus, thecorrosion products were mineral-like and the corrosion layer includedHA, Mg(OH)2 and other amorphous magnesium-substituted calciumphosphates [40].
The greatest distinction of the morphologies between the compos-ite and the Mg-Zn-Mn alloy after different immersion times might beascribed to the existence of carbon scaffold in the Mg alloy. Crevicecorrosion was found along the interface between Mg alloy and carbonscaffold. It can be seen from the Fig. 2(a) that the width of the liga-ments connecting the porous cell was about 90 μm. It can be inferredthat the width of crevice in the composite, as shown in Fig. 2(b),might also be around 90 μm. However, it is the most sensitive areaoccurring the crevice corrosion when the width of the crevice is be-tween 25–100 μm [44]. In addition, the existence of the Cl- ions intothe SBF solution accelerated this reaction. The infiltrated Mg alloyinside and outside the crevice consisted of the corrosion cell. TheMg alloy inside the crevice was regarded as micro anode sites andthat outside the crevice as macro cathode sites, giving rise to thecrevice corrosion [40]. The crannies were often covered with thecorrosion products in the crevice corrosion process so that the carbonscaffold could not be observed in the SEM images (Fig. 6(c, d)).
For the infiltrated Mg-Zn-Mn alloy, the second phase Mg7Zn3 hadhigher corrosion potentials than α-Mg matrix, therefore, Mg matrixcould be regarded as the micro anode sites and the second phaseMg7Zn3 as themicro cathode sites, resulting in microgalvanic corrosion[11]. On the other hand, it is well known that Mg is a relatively reactivealloy, when they are immersed in the SBF solution, Mg in aqueoussolution dissolves according to the following equations [40]:
Thus, during the immersion period, Mg(OH)2 was deposited on thesurface of the C/Mg-Zn-Mn composite. Nevertheless, the existence ofthe chloride ions (Cl-) would react with Mg(OH)2 and formed a moreresoluble MgCl2 so that the protective film was destroyed and it wasobserved that there were a number of small corrosion pits, as shownin Fig. 6(b). Witte [45] also reported that occasionally localised corro-sion of AZ91D-HA composite was interpreted as a break down andreformation of the protective surface layer. The breakdown of theprotective layer was responsible for the observed localised corrosionand occurred randomly during the whole test period.
3.3.3. Corrosion rate of compositeFig. 9 showed the variation of the mass loss percentage and the
corrosion rate of the immersed composite samples in SBF as a functionof the immersion time. It could be seen that the mass loss of thecomposite was stable in the first 150 h and then rapidly increased.The mass loss percentage reached about 52.9% during the overallimmersion time. It should be noted that after 30 day′ immersion, tothe naked eye, the large part of the composite specimens had beenbroken into tiny particles. From Fig. 9(b), the corrosion rate was0.043 mg/cm2h and 0.028 mg/cm2h after 3 and 7 days′ immersion,
Fig. 6. The morphologies of C/Mg-Zn-Mn composite after (a) 3, (b) 7, (c) 15 and (d) 30 days′ immersion in SBF solution.
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respectively. It indicated that the corrosion rate of the compositedecreased with the increase of the immersion time during the initialmeasuring period, but it increasedwith further increasing of immersiontime, and then reached a steady state after 20 days. Razavi et al. [24, 26]found that the corrosion rate of AZ91-20FA composite decreased withthe increase in the immersion time, and corrosion rate after 5 and
Fig. 7. EDS results of C/Mg-Zn-Mn composite after 15 days′ immersion in SBFsolution.
72 h was 0.8 and 0.08 mg/cm2h, respectively. This retardation of thecorrosion of the AZ91-FA nanocomposites can be explained by the accu-mulation of the corrosion products (magnesiumhydroxide, aswell as cal-cium magnesium phosphate and calcium carbonate). They formed aprotective layer on the surface of the AZ91-FA nanocomposites, retardingthe corrosion process. Zhang et al. [12] also reported that degradation rate
Fig. 8. XRD patterns of C/Mg-Zn-Mn composite after 3, 7, 15 and 30 days′ immersion in SBFsolution.
Fig. 9. (a) Mass loss percentage and (b) Corrosion rate of C/Mg-Zn-Mn composite versus immersion time.
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of Mg-Zn alloy after 30 days immersion in SBF was lower than that of3 days. The reason is that the Mg(OH)2 passive layer formed on the sam-ples surfaces and prohibited the degradation. When the immersion timeof the samples is less than 7 days, Mg(OH)2 passive layer formed on thesurface of C/Mg-Zn-Mn composite retards degradation, but with theincreasing of immersion time, the existence of the chloride ions (Cl-) inthe SBF solution would transform Mg(OH)2 into soluble MgCl2 [40], sothat the protective film was destroyed, and the corrosion rate wasaccelerated to some extent. Meanwhile, the exposed interface betweenporous carbon skeleton and Mg-Zn-Mn alloy (as seen in Fig. 6(b)) raisedexposed surface area of the samples, to some extent, this also acceleratedthe corrosion rate. As timewent on, the rate of the formation ofMg(OH)2is equal to that of the destruction of MgCl2, so the corrosion rate wassteady.
3.3.4. pH value of immersion solutionFig. 10 showed the pH values of the SBF solution incubating the
C/Mg-Zn-Mn composite at different immersion times. It revealedthat when the immersion time was not longer than 50 hours, thepH value increased quickly, i.e. from original 7.4 jumped rapidly to10.8 at 50 hours, and then the pH value stabilized at around 10.75and increased no more. In order to maintain the same compositionand pH value as the human body fluid, the immersion SBF solution
Fig. 10. pH values of the SBF solution incubating C/Mg-Zn-Mn composite at differentimmersion time.
for 15 days and 30 days were refreshed every other 2 days, so thepH value of SBF was not tested during the immersion interval.But from the Fig. 10, the pH value remained mostly a constantafter 50 hours, and the similar results were obtained in the otherimmersion tests [12, 29, 40].
The reason why the pH value of SBF solution incubating the com-posite raised rapidly in the early immersion period is that the signif-icant cathodic reaction occurred near the surface (Eq. (3)). When thereactions among all the ions get equilibrium, the pH values of the SBFsolution kept at a relative stable level. According to the Eq. (4), theMg(OH)2 film was expected to form in the following reaction. Theinsoluble protective corrosion layer (i.e. Mg(OH)2 and calcium phos-phate) formed during the stage of immersion in SBF can retard thedegradation. Thus in the early immersion period, the mass changeof the composite was slower than the next immersion. In addition,during the corrosion process, the HA would be formed (Fig. 9). Onone hand, the dissolved Mg(OH)2 film could provide favorable sitesfor HA [46]. On the other hand, the increasing pH value promotedthe HA nucleation by increasing the supersaturation of the SBF solu-tion with respect to HA. Consequently, a number of HA nuclei wereformed on the surface film and then HA would grow spontaneouslyby consuming the calcium and phosphate ions from the surroundingfluid [47].
Compared with the Mg-Zn-Mn alloy and porous carbon and otherbiomedical Mg composite [6, 11, 21, 29], the mechanical property andcorrosion behaviors of the C/Mg-Zn-Mn composite can satisfy therequirement of implant materials, so this composite is a very promisingcandidate for bone substitute.
4. Conclusion
The interpenetrating C/Mg-Zn-Mn composite for biodegradablebone implant material was fabricated by the suction casting method.Compared to the single porous carbon scaffold, the C/Mg-Zn-Mncomposite exhibited higher mechanical property (increased about150-fold). The mass loss percentage of the resulted composite was52.9% after 30 days′ immersion in the SBF solution at 37±0.5 °C.The corrosion rates were 0.043 mg/cm2h and 0.028 mg/cm2h after3 and 7 days′ immersion, respectively. The immersion test also indi-cated that carbon scaffold was left behind due to the much slowerdegradation rate of the ceramic scaffold. The corrosion products ofthe composite surface were mainly Mg(OH)2 and HA. Good proper-ties of the C/Mg-Zn-Mn composite indicated the possibility for newbiomaterials.
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Acknowledgement
The authors would like to thank the National Natural ScienceFoundation of China(51101039) and the Natural Science Foundationof Heilongjiang Province(E201004)
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