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Corrosion fatigue of a magnesium alloy in modified simulated body fluid
Engineering Fracture Mechanics xxx (2014) xxx–xxx
Contents lists available at ScienceDirect
Engineering Fracture Mechanics journal homepage: www.elsevier.com/locate/engfracmech
Corrosion fatigue of a magnesium alloy in modified simulated body fluid Sajjad Jafari a, R.K. Singh Raman a,b,⇑, Chris H.J. Davies a a b
Department of Mechanical & Aerospace Engineering, Monash University (Melbourne), VIC 3800, Australia Department of Chemical Engineering, Monash University (Melbourne), VIC 3800, Australia
a r t i c l e
i n f o
Article history: Available online xxxx Keywords: Implant materials Magnesium alloys Environmental cracking High cycle fatigue Hydrogen embrittlement
a b s t r a c t For magnesium (Mg) alloys to be used as temporary biodegradable implants it is essential to establish their resistance to body fluid-assisted cracking. In this paper the fatigue behaviour of a common magnesium alloy, AZ91D, is studied in air and in modified simulated body fluid (m-SBF), and the effect of different electrochemical conditions on corrosion fatigue life is investigated. The alloy was found to be susceptible to corrosion fatigue. Results suggest inclusions and corrosion pits to be the crack initiation sites, and hydrogen embrittlement to play a dominant role in cracking of AZ91D Mg alloy in m-SBF. Ó 2014 Elsevier Ltd. All rights reserved.
1. Introduction Metallic biomaterials such as titanium alloys, stainless steels and cobalt-chromium alloys are widely used implant materials due to their corrosion resistance and strength [1,2]. However, when these materials are used as temporary implants, their retention in the body becomes unnecessary after they have fulfilled their function, and a removal surgery is required which increases the health care cost as well as inconvenience to the patient [3]. Moreover, the mechanical properties of these alloys are considerably different from those of human bone which results in the problem of ‘stress shielding’ and consequent reduction in bone density [4]. Magnesium (Mg) alloys are suitable as potential temporary biomedical implants because they are biodegradable and can completely dissolve in the body [5–8], which eliminates the need for secondary surgery to remove an implant. Magnesium is also biocompatible, is essential to human metabolism, and in addition any excess Mg is harmlessly excreted [4,9]. Furthermore magnesium has mechanical properties much closer to bone and this mitigates stress shielding [8]. These properties make magnesium and its alloys suitable as temporary orthopaedic implants (e.g. bone plates and screws) and cardiovascular implants (e.g. stents). In spite of these advantages Mg alloys have found very little actual use in implants, primarily because they tend to corrode too quickly in chloride solutions including physiological environments and therefore lose their mechanical integrity before accomplishing their purpose [10–13]. Orthopaedic and cardiovascular implants generally experience cyclic loading [14,15], which along with the corrosive physiological environment can cause corrosion assisted cracking. Depending on the nature of loading, corrosion assisted cracking includes stress corrosion cracking (tensile loading) and corrosion fatigue (cyclic loading). Stress corrosion cracking (SCC) of Mg alloys has been widely investigated including in physiological environments [16–20] but corrosion fatigue (CF)
⇑ Corresponding author at: Department of Mechanical & Aerospace Engineering, Bldg 31, Monash University – Clayton Campus (Melbourne), VIC 3800, Australia. Tel.: +61 399053671. E-mail address: [email protected] (R.K. Singh Raman). http://dx.doi.org/10.1016/j.engfracmech.2014.07.007 0013-7944/Ó 2014 Elsevier Ltd. All rights reserved.
Please cite this article in press as: Jafari S et al. Corrosion fatigue of a magnesium alloy in modified simulated body fluid. Engng Fract Mech (2014), http://dx.doi.org/10.1016/j.engfracmech.2014.07.007
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Nomenclature Symbols
rLH rCE
fatigue strength in air fatigue strength in corrosive environment
Abbreviations Al aluminium BEI back scattered electron imaging CF corrosion fatigue Cl chloride EDX energy-dispersive X-ray spectroscopy HEPES 2-(4-(2-hydroxyethyl)-1-piperazinyl) ethanesulfonic acid ICP-AES inductively coupled plasma atomic emission spectroscopy Mg magnesium m-SBF modified simulated body fluid NDE negative difference effect OCP open circuit potential SCC stress corrosion cracking SEM scanning electron microscopy
of Mg alloys in human body fluid has received little attention even though CF fractures cause catastrophic failures of biomedical implants [14,21–24]. There are two main mechanisms for SCC of Mg alloys: hydrogen induced cracking (hydrogen embrittlement), and dissolution assisted cracking [25,26]. But there is limited mechanistic understanding of the role of electrochemical polarization on corrosion fatigue cracking of Mg alloys. Aluminium (Al) containing Mg alloys are widely used due to the beneficial roles of Al in corrosion resistance and mechanical properties [27,28]. The AZ91D alloy exhibits good biocompatibility and causes no harm to the surrounding tissues [6,29,30]; in vivo and in vitro studies on this alloy have shown no adverse toxicity due to Al. However, there are also reports to suggest that Al causes Alzheimer’s disease, muscle fibre damage, as well as Al3+ ions combining with inorganic phosphate, causing phosphate deficiencies in the body [27,31–33]. Consequently, it is unlikely that AZ91D can be used for implants. However, AZ91D is the most investigated alloy for corrosion assisted cracking in chloride solutions, and the present study on CF of this alloy in the physiological environment provides an improved mechanistic understanding, and baseline data for magnesium alloys that can be actually used as biodegradable implants. 2. Experimental procedure 2.1. Materials Magnesium alloy AZ91D alloy was received in sand-cast form, the chemical composition of which was analysed by inductively coupled plasma atomic emission spectroscopy (Table 1). The tensile properties of the AZ91D are listed in Table 2. 2.2. Fatigue and corrosion fatigue tests Fatigue samples with 6 mm gauge diameter and 15 mm gauge length (ASTM E466 [34]) were tested (Fig. 1a). The specimen gauge was abraded in the loading direction with 1200 and 2500 grit emery papers, and then polished with 1 lm diamond paste followed by cleaning with ethanol and de-ionised water. In order to obtain an appropriate alignment of the specimen and also save the amount of material at two shoulders, a sample holder (made of high strength steel) with collet chuck was used for gripping (Fig. 1b). A corrosion chamber made of acrylic was attached to the sample (Fig. 2). Any leakage of solution from the chamber was prevented by using two O-rings fitted to the shoulders of the fatigue specimens. Modified simulated body fluid (m-SBF) used for corrosion fatigue experiments has a chloride ion concentration close to that of blood plasma and was buffered with 2-(4-(2-hydroxyethyl)-1-piperazinyl) ethanesulfonic acid (HEPES) at a pH of 7.4 (Table 3) [35].
Table 1 Chemical composition of AZ91D magnesium alloy. Element
Mg
AL
Zn
Mn
Cu
Fe
Ni
Si
Be
Wt.%
Bal
8.89
0.78
0.20
0.002
0.002
<0.001
<0.01
<0.001
Please cite this article in press as: Jafari S et al. Corrosion fatigue of a magnesium alloy in modified simulated body fluid. Engng Fract Mech (2014), http://dx.doi.org/10.1016/j.engfracmech.2014.07.007
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S. Jafari et al. / Engineering Fracture Mechanics xxx (2014) xxx–xxx Table 2 Mechanical properties of AZ91D magnesium alloy.
AZ91D
Yield strength (MPa)
Tensile strength (MPa)
Elongation (%)
105 ± 10
180 ± 12
±2
Fig. 1. (a) Fatigue specimen (dimensions are in mm) and (b) sample holder, collet chuck and nut.
Fig. 2. Schematic of the corrosion fatigue rig experimental setup.
Table 3 Amount of reagents for preparing 1000 ml of the m-SBF solution. Reagents
Amount
NaCl NaHCO3 Na2CO3 KCl K2HPO43H2O MgCl26H2O 0.2 mol l 1 NaOH HEPES CaCl2 Na2SO4 1 mol l 1 NaOH
5.403 g 0.504 g 0.426 g 0.225 g 0.23 g 0.311 g 100 ml 17.892 g 0.293 g 0.072 g 15 ml
Please cite this article in press as: Jafari S et al. Corrosion fatigue of a magnesium alloy in modified simulated body fluid. Engng Fract Mech (2014), http://dx.doi.org/10.1016/j.engfracmech.2014.07.007
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Axial fatigue tests were performed using a servo-hydraulic Instron testing machine with sinusoidal wave form and a stress ratio of 1 (fully reversed) with a frequency of 5 Hz. The tests continued either until complete failure of the samples or were stopped if the specimen did not fail in 107 cycles (run-out). The temperature of m-SBF was maintained at 37 °C using a thermostat. To simulate the in vitro flow of body fluid, a submersible pump was used for continuous circulation of the fluid through corrosion chamber. The solution flow rate was adjusted to 120 ml/min. In order to understand the mechanism of cracking, a few sets of experiments were performed under different imposed electrochemical conditions, controlled using a potentiostat in conjunction with a saturated calomel electrode as reference electrode and a platinum mesh counter electrode. 2.3. Fractography The fracture surfaces were examined using scanning electron microscopy (SEM) equipped with energy-dispersive X-ray spectroscopy (EDX) and back scattered electron imaging (BEI). In order to remove corrosion products off the fracture surface, the samples tested in m-SBF were cleaned using 20% wt. CrO3 and 10% wt. AgNO3 solution before fractography. 3. Results 3.1. Microstructure of as-cast AZ91D The microstructure of the as-received AZ91D alloy is mainly composed of a-Mg phase and discontinuous secondary Mg17Al12 phase (Fig. 3). Castings of Mg alloy generally incorporate inclusions and these were found to be MgO. The expected AlMn phase and Mg17Al12 are also observed in the microstructure of the alloy. MgO inclusion can act as nucleation sites for the AlMn phase during solidification of the AZ91D [36]. 3.2. Fatigue and corrosion fatigue testing The as-cast AZ91D magnesium alloy has a fatigue limit of 57 MPa in air, and a fatigue strength of 17 MPa at 5 105 cycles in m-SBF (Fig. 4). Fractographs of a sample tested in air (Fig. 5) show the three stages of fatigue failures; i.e., crack initiation, crack propagation and final overload failure. The fatigue crack nucleated from the MgO inclusions (arrows in Fig. 5a and c). Close to the initiation site the fracture surface is relatively flat whereas as stress intensity increases away from this site a serrated and faceted path is observed. Cleavage striations are seen on the surface of individual facets at high magnification (Fig. 5b), indicative of the progress of the crack front. The mechanical overload failure (Fig. 5d) was characterized by quasi-cleavage fracture. Samples tested in m-SBF (Fig. 6) show corrosion pit formation in the specimen gauge length and subsequent fatigue crack that developed from the pit (Fig. 6b). Ductile striations are formed close to the initiation site (Fig. 6c) while a serrated path is observed farther from initiation site (Fig. 6d). The stepped interlocking features are probably brittle type failure involving hydrogen [37]. Mechanical overload of this sample shows a quasi-cleavage failure (Fig. 6e). 3.3. Corrosion fatigue testing under different electrochemical conditions The potentiodynamic polarization curve for AZ91D tested in m-SBF (Fig. 7) identifies the cathodic and anodic regimes for fatigue testing under continuous cathodic and anodic charging conditions. Our experiments were performed at 1.85 and 1.45 V vs. SCE. When the applied stress levels were greater than the fatigue limit in air (57 MPa), the corrosion fatigue (CF) lives were similar for the specimens tested under cathodic charging condition and those tested at open circuit potential (OCP) in m-SBF (Fig. 4). However, there was a remarkable difference in CF lives under the two conditions when tests were conducted at the stress levels lower than the fatigue limit in the air (Fig. 4). In this regime, at constant stress level the specimens tested under cathodic charging (i.e., when anodic dissolution is minimal) showed considerably longer fatigue lives. On the other hand, anodic charging somewhat shortened the fatigue lives of the samples in m-SBF at all applied stresses. The samples that failed under different electrochemical conditions at 67 MPa showed more localized corrosion/pitting (producing a greater susceptibility to cracking) at the anodic potential (Fig. 8a) compared to the cathodic potential at which fewer pits formed (Fig. 8b). We see transgranular cracking for the samples tested under cathodic charging condition (Fig. 8d and e), which is attributed to the predominance of mechanism involving hydrogen diffusing to the high stress intensity areas ahead of crack tip causing localized embrittlement [17,37]. 4. Discussion 4.1. Fatigue behaviour in air and m-SBF The AZ91D showed a fatigue limit in air (Fig. 4), which occurs when defects initiate cracks by enhancing stress intensity [38,39] and our observations indicate that it is the MgO that is likely to be responsible for this (Fig. 5a). This is supported by others [40]. Apart from triggering crack initiation, casting defects and inclusions can also accelerate crack propagation. Please cite this article in press as: Jafari S et al. Corrosion fatigue of a magnesium alloy in modified simulated body fluid. Engng Fract Mech (2014), http://dx.doi.org/10.1016/j.engfracmech.2014.07.007
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Fig. 3. Microstructural analysis of the tested AZ91D: (a) backscatter electron micrograph, and EDX analysis of: (b) Al–Mn intermetallic, (c) Mg17Al12 intermetallic and (d) MgO inclusion.
The effect of a corrosive environment on fatigue strength can be described by rLHrLHrCE where rLH is fatigue strength in air, and rCE fatigue strength in the corrosive environment [41]. On this measure the fatigue strength in air reduced by 70% when tests were carried out in m-SBF, indicating a high sensitivity of the fatigue strength of AZ91D Mg alloy to the m-SBF environment. It is well-established that magnesium alloys are susceptible to pitting in chloride (Cl ) containing solutions [41] including in simulated body fluid [12]. The drastic reduction in fatigue strength under m-SBF environment is mainly attributed to corrosion pit formation [42]. However, our work differs from others [42] in that corrosion pits appear to be crack initiation sites for all the samples tested in m-SBF, even for those tested at stress amplitudes greater than the fatigue limit in air. 4.2. Effect of polarization on corrosion fatigue cracking Under cathodic charging conditions when the stress levels were lower than the fatigue limit in air, the fatigue life increases considerably, whereas fatigue life was similar to those at OCP when stress levels were higher than the fatigue limit in air. This behaviour is attributed to the difference in predominance of pitting in the two stress regimes. First, magnesium alloys suffer pitting even under cathodic charging condition though with a considerably less intensity (Fig. 8d). At high stress Please cite this article in press as: Jafari S et al. Corrosion fatigue of a magnesium alloy in modified simulated body fluid. Engng Fract Mech (2014), http://dx.doi.org/10.1016/j.engfracmech.2014.07.007
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Fig. 4. S–N curves for AZ91D magnesium alloy under different conditions.
Fig. 5. Micrographs of the sample that failed at 95,970 cycles at 71 MPa in air: (a) crack initiation from MgO inclusions, (b) crack propagation farther from initiation site, (c) EDX analysis of inclusion at the initiation site and (d) mechanical overload failure region (arrows show the inclusions at the bottom of voids).
amplitudes, the required size of a pit to develop into a fatigue crack will be small, and hence the pits developed under the cathodic charging condition is sufficient to facilitate the fatigue cracking. But, such pit depths are possibly not sufficient for causing fatigue cracking at low stress amplitudes. Please cite this article in press as: Jafari S et al. Corrosion fatigue of a magnesium alloy in modified simulated body fluid. Engng Fract Mech (2014), http://dx.doi.org/10.1016/j.engfracmech.2014.07.007
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Fig. 6. Fracture features of the sample failed at 80,512 cycles at 58 MPa in m-SBF: (a) overview of the fracture surface and indication of initiation sites (dark arrows), (b) crack initiation from corrosion pit in the specimen gauge, (c) striation formation close to the initiation site, (d) facet features farther from initiation site and (e) mechanical overload failure.
On the other hand, the anodic charging somewhat shortened the fatigue life over the entire stress amplitude range. Generally, an anodic increase of the applied potential causes an increase in the anodic dissolution rate, with concurrent suppression of cathodic activity such as hydrogen evolution. Interestingly, in the case of magnesium alloys, as result of the negative difference effect (NDE) [43], hydrogen evolution occurs even at anodic over-potential. Such hydrogen evolution and the concurrent pitting explain the greater tendency for corrosion fatigue cracking under anodic charging condition.
4.3. General considerations Orthopaedic implants need to maintain their mechanical integrity (20–25 MPa) for about 5 105 cycles [14]. However, we have shown that as-cast AZ91D Mg alloy possesses a fatigue strength of 17 MPa in m-SBF, which is considerably lower than required. It is probable that the casting defects reduced the fatigue and corrosion fatigue resistance of the AZ91D alloy, indicating the need for a defect- and inclusion-free alloy. The decrease in the corrosion fatigue strength of as-cast AZ91D magnesium alloy at 5 105 cycles was about 70% of that of similar sample tested in air, which supports the need to assess corrosion fatigue behaviour of candidate wrought-Mg implant alloys in physiological environments, before putting implants of such alloys into service.
Please cite this article in press as: Jafari S et al. Corrosion fatigue of a magnesium alloy in modified simulated body fluid. Engng Fract Mech (2014), http://dx.doi.org/10.1016/j.engfracmech.2014.07.007
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Fig. 7. Potentiodynamic polarization plot for AZ91D tested in m-SBF.
Fig. 8. Fractographs of corrosion fatigue samples failed at 67 MPa under: (a) anodic charging, (b) cathodic charging condition, and crack initiation under: (c) anodic charging, (d) cathodic charging, and (e) evidence of transgranular cracking for cathodically charged specimen.
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5. Conclusion The main conclusions of this study are: 1) Fatigue cracking of AZ91D Mg alloy in m-SBF is a mechano-chemical process which is profoundly influenced by the electrochemical condition of the corrosive environment. At stresses higher than the fatigue limit in air, the loss of fatigue resistance in m-SBF is attributed to the mechanical breakdown of the protective film. In the low stress regime, the electrochemical condition (pitting) has an opportunity to play a dominant role, thereby considerably lowering the fatigue resistance. 2) Pits are initiation sites for corrosion fatigue cracking and their nucleation and growth is influenced by electrochemical conditions and applied stress levels. Our results suggest that enhancing the pitting resistance will improve the corrosion fatigue behaviour of Mg alloys in m-SBF. 3) Hydrogen embrittlement was the main crack propagation mechanism for AZ91D magnesium alloy in m-SBF. Fractography revealed transgranular cracking as the predominant mode of corrosion fatigue of AZ91D in m-SBF.
Acknowledgements Author Sajjad Jafari would like to acknowledge Monash University Institute of Graduate Research (MIGR) for providing research scholarships. He thanks Dr. Lokesh Choudhary for providing training in electrochemical experimental techniques. He also acknowledges support from the Department of Mechanical & Aerospace Engineering, Monash University, Australia and the facility and services of Monash Centre for Electron Microscopy (MCEM). References [1] Okazaki Y, Gotoh E. Metal release from stainless steel, Co–Cr–Mo–Ni–Fe and Ni–Ti alloys in vascular implants. Corros Sci 2008;50:3429–38. [2] Niinomi M, Nakai M, Hieda J. Development of new metallic alloys for biomedical applications. Acta Biomater 2012;8:3888–903. [3] Gunde P, Hänzi AC, Sologubenko AS, Uggowitzer PJ. High-strength magnesium alloys for degradable implant applications. Mater Sci Engng, A 2011;528:1047–54. [4] Staiger MP, Pietak AM, Huadmai J, Dias G. Magnesium and its alloys as orthopedic biomaterials: a review. 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Contents lists available at ScienceDirect
Engineering Fracture Mechanics journal homepage: www.elsevier.com/locate/engfracmech
Corrosion fatigue of a magnesium alloy in modified simulated body fluid Sajjad Jafari a, R.K. Singh Raman a,b,⇑, Chris H.J. Davies a a b
Department of Mechanical & Aerospace Engineering, Monash University (Melbourne), VIC 3800, Australia Department of Chemical Engineering, Monash University (Melbourne), VIC 3800, Australia
a r t i c l e
i n f o
Article history: Available online xxxx Keywords: Implant materials Magnesium alloys Environmental cracking High cycle fatigue Hydrogen embrittlement
a b s t r a c t For magnesium (Mg) alloys to be used as temporary biodegradable implants it is essential to establish their resistance to body fluid-assisted cracking. In this paper the fatigue behaviour of a common magnesium alloy, AZ91D, is studied in air and in modified simulated body fluid (m-SBF), and the effect of different electrochemical conditions on corrosion fatigue life is investigated. The alloy was found to be susceptible to corrosion fatigue. Results suggest inclusions and corrosion pits to be the crack initiation sites, and hydrogen embrittlement to play a dominant role in cracking of AZ91D Mg alloy in m-SBF. Ó 2014 Elsevier Ltd. All rights reserved.
1. Introduction Metallic biomaterials such as titanium alloys, stainless steels and cobalt-chromium alloys are widely used implant materials due to their corrosion resistance and strength [1,2]. However, when these materials are used as temporary implants, their retention in the body becomes unnecessary after they have fulfilled their function, and a removal surgery is required which increases the health care cost as well as inconvenience to the patient [3]. Moreover, the mechanical properties of these alloys are considerably different from those of human bone which results in the problem of ‘stress shielding’ and consequent reduction in bone density [4]. Magnesium (Mg) alloys are suitable as potential temporary biomedical implants because they are biodegradable and can completely dissolve in the body [5–8], which eliminates the need for secondary surgery to remove an implant. Magnesium is also biocompatible, is essential to human metabolism, and in addition any excess Mg is harmlessly excreted [4,9]. Furthermore magnesium has mechanical properties much closer to bone and this mitigates stress shielding [8]. These properties make magnesium and its alloys suitable as temporary orthopaedic implants (e.g. bone plates and screws) and cardiovascular implants (e.g. stents). In spite of these advantages Mg alloys have found very little actual use in implants, primarily because they tend to corrode too quickly in chloride solutions including physiological environments and therefore lose their mechanical integrity before accomplishing their purpose [10–13]. Orthopaedic and cardiovascular implants generally experience cyclic loading [14,15], which along with the corrosive physiological environment can cause corrosion assisted cracking. Depending on the nature of loading, corrosion assisted cracking includes stress corrosion cracking (tensile loading) and corrosion fatigue (cyclic loading). Stress corrosion cracking (SCC) of Mg alloys has been widely investigated including in physiological environments [16–20] but corrosion fatigue (CF)
⇑ Corresponding author at: Department of Mechanical & Aerospace Engineering, Bldg 31, Monash University – Clayton Campus (Melbourne), VIC 3800, Australia. Tel.: +61 399053671. E-mail address: [email protected] (R.K. Singh Raman). http://dx.doi.org/10.1016/j.engfracmech.2014.07.007 0013-7944/Ó 2014 Elsevier Ltd. All rights reserved.
Please cite this article in press as: Jafari S et al. Corrosion fatigue of a magnesium alloy in modified simulated body fluid. Engng Fract Mech (2014), http://dx.doi.org/10.1016/j.engfracmech.2014.07.007
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Nomenclature Symbols
rLH rCE
fatigue strength in air fatigue strength in corrosive environment
Abbreviations Al aluminium BEI back scattered electron imaging CF corrosion fatigue Cl chloride EDX energy-dispersive X-ray spectroscopy HEPES 2-(4-(2-hydroxyethyl)-1-piperazinyl) ethanesulfonic acid ICP-AES inductively coupled plasma atomic emission spectroscopy Mg magnesium m-SBF modified simulated body fluid NDE negative difference effect OCP open circuit potential SCC stress corrosion cracking SEM scanning electron microscopy
of Mg alloys in human body fluid has received little attention even though CF fractures cause catastrophic failures of biomedical implants [14,21–24]. There are two main mechanisms for SCC of Mg alloys: hydrogen induced cracking (hydrogen embrittlement), and dissolution assisted cracking [25,26]. But there is limited mechanistic understanding of the role of electrochemical polarization on corrosion fatigue cracking of Mg alloys. Aluminium (Al) containing Mg alloys are widely used due to the beneficial roles of Al in corrosion resistance and mechanical properties [27,28]. The AZ91D alloy exhibits good biocompatibility and causes no harm to the surrounding tissues [6,29,30]; in vivo and in vitro studies on this alloy have shown no adverse toxicity due to Al. However, there are also reports to suggest that Al causes Alzheimer’s disease, muscle fibre damage, as well as Al3+ ions combining with inorganic phosphate, causing phosphate deficiencies in the body [27,31–33]. Consequently, it is unlikely that AZ91D can be used for implants. However, AZ91D is the most investigated alloy for corrosion assisted cracking in chloride solutions, and the present study on CF of this alloy in the physiological environment provides an improved mechanistic understanding, and baseline data for magnesium alloys that can be actually used as biodegradable implants. 2. Experimental procedure 2.1. Materials Magnesium alloy AZ91D alloy was received in sand-cast form, the chemical composition of which was analysed by inductively coupled plasma atomic emission spectroscopy (Table 1). The tensile properties of the AZ91D are listed in Table 2. 2.2. Fatigue and corrosion fatigue tests Fatigue samples with 6 mm gauge diameter and 15 mm gauge length (ASTM E466 [34]) were tested (Fig. 1a). The specimen gauge was abraded in the loading direction with 1200 and 2500 grit emery papers, and then polished with 1 lm diamond paste followed by cleaning with ethanol and de-ionised water. In order to obtain an appropriate alignment of the specimen and also save the amount of material at two shoulders, a sample holder (made of high strength steel) with collet chuck was used for gripping (Fig. 1b). A corrosion chamber made of acrylic was attached to the sample (Fig. 2). Any leakage of solution from the chamber was prevented by using two O-rings fitted to the shoulders of the fatigue specimens. Modified simulated body fluid (m-SBF) used for corrosion fatigue experiments has a chloride ion concentration close to that of blood plasma and was buffered with 2-(4-(2-hydroxyethyl)-1-piperazinyl) ethanesulfonic acid (HEPES) at a pH of 7.4 (Table 3) [35].
Table 1 Chemical composition of AZ91D magnesium alloy. Element
Mg
AL
Zn
Mn
Cu
Fe
Ni
Si
Be
Wt.%
Bal
8.89
0.78
0.20
0.002
0.002
<0.001
<0.01
<0.001
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S. Jafari et al. / Engineering Fracture Mechanics xxx (2014) xxx–xxx Table 2 Mechanical properties of AZ91D magnesium alloy.
AZ91D
Yield strength (MPa)
Tensile strength (MPa)
Elongation (%)
105 ± 10
180 ± 12
±2
Fig. 1. (a) Fatigue specimen (dimensions are in mm) and (b) sample holder, collet chuck and nut.
Fig. 2. Schematic of the corrosion fatigue rig experimental setup.
Table 3 Amount of reagents for preparing 1000 ml of the m-SBF solution. Reagents
Amount
NaCl NaHCO3 Na2CO3 KCl K2HPO43H2O MgCl26H2O 0.2 mol l 1 NaOH HEPES CaCl2 Na2SO4 1 mol l 1 NaOH
5.403 g 0.504 g 0.426 g 0.225 g 0.23 g 0.311 g 100 ml 17.892 g 0.293 g 0.072 g 15 ml
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Axial fatigue tests were performed using a servo-hydraulic Instron testing machine with sinusoidal wave form and a stress ratio of 1 (fully reversed) with a frequency of 5 Hz. The tests continued either until complete failure of the samples or were stopped if the specimen did not fail in 107 cycles (run-out). The temperature of m-SBF was maintained at 37 °C using a thermostat. To simulate the in vitro flow of body fluid, a submersible pump was used for continuous circulation of the fluid through corrosion chamber. The solution flow rate was adjusted to 120 ml/min. In order to understand the mechanism of cracking, a few sets of experiments were performed under different imposed electrochemical conditions, controlled using a potentiostat in conjunction with a saturated calomel electrode as reference electrode and a platinum mesh counter electrode. 2.3. Fractography The fracture surfaces were examined using scanning electron microscopy (SEM) equipped with energy-dispersive X-ray spectroscopy (EDX) and back scattered electron imaging (BEI). In order to remove corrosion products off the fracture surface, the samples tested in m-SBF were cleaned using 20% wt. CrO3 and 10% wt. AgNO3 solution before fractography. 3. Results 3.1. Microstructure of as-cast AZ91D The microstructure of the as-received AZ91D alloy is mainly composed of a-Mg phase and discontinuous secondary Mg17Al12 phase (Fig. 3). Castings of Mg alloy generally incorporate inclusions and these were found to be MgO. The expected AlMn phase and Mg17Al12 are also observed in the microstructure of the alloy. MgO inclusion can act as nucleation sites for the AlMn phase during solidification of the AZ91D [36]. 3.2. Fatigue and corrosion fatigue testing The as-cast AZ91D magnesium alloy has a fatigue limit of 57 MPa in air, and a fatigue strength of 17 MPa at 5 105 cycles in m-SBF (Fig. 4). Fractographs of a sample tested in air (Fig. 5) show the three stages of fatigue failures; i.e., crack initiation, crack propagation and final overload failure. The fatigue crack nucleated from the MgO inclusions (arrows in Fig. 5a and c). Close to the initiation site the fracture surface is relatively flat whereas as stress intensity increases away from this site a serrated and faceted path is observed. Cleavage striations are seen on the surface of individual facets at high magnification (Fig. 5b), indicative of the progress of the crack front. The mechanical overload failure (Fig. 5d) was characterized by quasi-cleavage fracture. Samples tested in m-SBF (Fig. 6) show corrosion pit formation in the specimen gauge length and subsequent fatigue crack that developed from the pit (Fig. 6b). Ductile striations are formed close to the initiation site (Fig. 6c) while a serrated path is observed farther from initiation site (Fig. 6d). The stepped interlocking features are probably brittle type failure involving hydrogen [37]. Mechanical overload of this sample shows a quasi-cleavage failure (Fig. 6e). 3.3. Corrosion fatigue testing under different electrochemical conditions The potentiodynamic polarization curve for AZ91D tested in m-SBF (Fig. 7) identifies the cathodic and anodic regimes for fatigue testing under continuous cathodic and anodic charging conditions. Our experiments were performed at 1.85 and 1.45 V vs. SCE. When the applied stress levels were greater than the fatigue limit in air (57 MPa), the corrosion fatigue (CF) lives were similar for the specimens tested under cathodic charging condition and those tested at open circuit potential (OCP) in m-SBF (Fig. 4). However, there was a remarkable difference in CF lives under the two conditions when tests were conducted at the stress levels lower than the fatigue limit in the air (Fig. 4). In this regime, at constant stress level the specimens tested under cathodic charging (i.e., when anodic dissolution is minimal) showed considerably longer fatigue lives. On the other hand, anodic charging somewhat shortened the fatigue lives of the samples in m-SBF at all applied stresses. The samples that failed under different electrochemical conditions at 67 MPa showed more localized corrosion/pitting (producing a greater susceptibility to cracking) at the anodic potential (Fig. 8a) compared to the cathodic potential at which fewer pits formed (Fig. 8b). We see transgranular cracking for the samples tested under cathodic charging condition (Fig. 8d and e), which is attributed to the predominance of mechanism involving hydrogen diffusing to the high stress intensity areas ahead of crack tip causing localized embrittlement [17,37]. 4. Discussion 4.1. Fatigue behaviour in air and m-SBF The AZ91D showed a fatigue limit in air (Fig. 4), which occurs when defects initiate cracks by enhancing stress intensity [38,39] and our observations indicate that it is the MgO that is likely to be responsible for this (Fig. 5a). This is supported by others [40]. Apart from triggering crack initiation, casting defects and inclusions can also accelerate crack propagation. Please cite this article in press as: Jafari S et al. Corrosion fatigue of a magnesium alloy in modified simulated body fluid. Engng Fract Mech (2014), http://dx.doi.org/10.1016/j.engfracmech.2014.07.007
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Fig. 3. Microstructural analysis of the tested AZ91D: (a) backscatter electron micrograph, and EDX analysis of: (b) Al–Mn intermetallic, (c) Mg17Al12 intermetallic and (d) MgO inclusion.
The effect of a corrosive environment on fatigue strength can be described by rLHrLHrCE where rLH is fatigue strength in air, and rCE fatigue strength in the corrosive environment [41]. On this measure the fatigue strength in air reduced by 70% when tests were carried out in m-SBF, indicating a high sensitivity of the fatigue strength of AZ91D Mg alloy to the m-SBF environment. It is well-established that magnesium alloys are susceptible to pitting in chloride (Cl ) containing solutions [41] including in simulated body fluid [12]. The drastic reduction in fatigue strength under m-SBF environment is mainly attributed to corrosion pit formation [42]. However, our work differs from others [42] in that corrosion pits appear to be crack initiation sites for all the samples tested in m-SBF, even for those tested at stress amplitudes greater than the fatigue limit in air. 4.2. Effect of polarization on corrosion fatigue cracking Under cathodic charging conditions when the stress levels were lower than the fatigue limit in air, the fatigue life increases considerably, whereas fatigue life was similar to those at OCP when stress levels were higher than the fatigue limit in air. This behaviour is attributed to the difference in predominance of pitting in the two stress regimes. First, magnesium alloys suffer pitting even under cathodic charging condition though with a considerably less intensity (Fig. 8d). At high stress Please cite this article in press as: Jafari S et al. Corrosion fatigue of a magnesium alloy in modified simulated body fluid. Engng Fract Mech (2014), http://dx.doi.org/10.1016/j.engfracmech.2014.07.007
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Fig. 4. S–N curves for AZ91D magnesium alloy under different conditions.
Fig. 5. Micrographs of the sample that failed at 95,970 cycles at 71 MPa in air: (a) crack initiation from MgO inclusions, (b) crack propagation farther from initiation site, (c) EDX analysis of inclusion at the initiation site and (d) mechanical overload failure region (arrows show the inclusions at the bottom of voids).
amplitudes, the required size of a pit to develop into a fatigue crack will be small, and hence the pits developed under the cathodic charging condition is sufficient to facilitate the fatigue cracking. But, such pit depths are possibly not sufficient for causing fatigue cracking at low stress amplitudes. Please cite this article in press as: Jafari S et al. Corrosion fatigue of a magnesium alloy in modified simulated body fluid. Engng Fract Mech (2014), http://dx.doi.org/10.1016/j.engfracmech.2014.07.007
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Fig. 6. Fracture features of the sample failed at 80,512 cycles at 58 MPa in m-SBF: (a) overview of the fracture surface and indication of initiation sites (dark arrows), (b) crack initiation from corrosion pit in the specimen gauge, (c) striation formation close to the initiation site, (d) facet features farther from initiation site and (e) mechanical overload failure.
On the other hand, the anodic charging somewhat shortened the fatigue life over the entire stress amplitude range. Generally, an anodic increase of the applied potential causes an increase in the anodic dissolution rate, with concurrent suppression of cathodic activity such as hydrogen evolution. Interestingly, in the case of magnesium alloys, as result of the negative difference effect (NDE) [43], hydrogen evolution occurs even at anodic over-potential. Such hydrogen evolution and the concurrent pitting explain the greater tendency for corrosion fatigue cracking under anodic charging condition.
4.3. General considerations Orthopaedic implants need to maintain their mechanical integrity (20–25 MPa) for about 5 105 cycles [14]. However, we have shown that as-cast AZ91D Mg alloy possesses a fatigue strength of 17 MPa in m-SBF, which is considerably lower than required. It is probable that the casting defects reduced the fatigue and corrosion fatigue resistance of the AZ91D alloy, indicating the need for a defect- and inclusion-free alloy. The decrease in the corrosion fatigue strength of as-cast AZ91D magnesium alloy at 5 105 cycles was about 70% of that of similar sample tested in air, which supports the need to assess corrosion fatigue behaviour of candidate wrought-Mg implant alloys in physiological environments, before putting implants of such alloys into service.
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Fig. 7. Potentiodynamic polarization plot for AZ91D tested in m-SBF.
Fig. 8. Fractographs of corrosion fatigue samples failed at 67 MPa under: (a) anodic charging, (b) cathodic charging condition, and crack initiation under: (c) anodic charging, (d) cathodic charging, and (e) evidence of transgranular cracking for cathodically charged specimen.
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5. Conclusion The main conclusions of this study are: 1) Fatigue cracking of AZ91D Mg alloy in m-SBF is a mechano-chemical process which is profoundly influenced by the electrochemical condition of the corrosive environment. At stresses higher than the fatigue limit in air, the loss of fatigue resistance in m-SBF is attributed to the mechanical breakdown of the protective film. In the low stress regime, the electrochemical condition (pitting) has an opportunity to play a dominant role, thereby considerably lowering the fatigue resistance. 2) Pits are initiation sites for corrosion fatigue cracking and their nucleation and growth is influenced by electrochemical conditions and applied stress levels. Our results suggest that enhancing the pitting resistance will improve the corrosion fatigue behaviour of Mg alloys in m-SBF. 3) Hydrogen embrittlement was the main crack propagation mechanism for AZ91D magnesium alloy in m-SBF. Fractography revealed transgranular cracking as the predominant mode of corrosion fatigue of AZ91D in m-SBF.
Acknowledgements Author Sajjad Jafari would like to acknowledge Monash University Institute of Graduate Research (MIGR) for providing research scholarships. He thanks Dr. Lokesh Choudhary for providing training in electrochemical experimental techniques. He also acknowledges support from the Department of Mechanical & Aerospace Engineering, Monash University, Australia and the facility and services of Monash Centre for Electron Microscopy (MCEM). References [1] Okazaki Y, Gotoh E. Metal release from stainless steel, Co–Cr–Mo–Ni–Fe and Ni–Ti alloys in vascular implants. Corros Sci 2008;50:3429–38. [2] Niinomi M, Nakai M, Hieda J. Development of new metallic alloys for biomedical applications. Acta Biomater 2012;8:3888–903. [3] Gunde P, Hänzi AC, Sologubenko AS, Uggowitzer PJ. High-strength magnesium alloys for degradable implant applications. Mater Sci Engng, A 2011;528:1047–54. [4] Staiger MP, Pietak AM, Huadmai J, Dias G. Magnesium and its alloys as orthopedic biomaterials: a review. 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Please cite this article in press as: Jafari S et al. Corrosion fatigue of a magnesium alloy in modified simulated body fluid. Engng Fract Mech (2014), http://dx.doi.org/10.1016/j.engfracmech.2014.07.007