Corrosion Science 91 (2015) 29–36

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Corrosion Science

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Assessing the degradation performance of ultrahigh-purity magnesiumin vitro and in vivo

http://dx.doi.org/10.1016/j.corsci.2014.09.0080010-938X/� 2014 Elsevier Ltd. All rights reserved.

⇑ Corresponding author. Tel.: +41 44 632 2565; fax: +41 44 633 1421.E-mail address: joerg.loeffler@mat.ethz.ch (J.F. Löffler).

Joëlle Hofstetter a, Elisabeth Martinelli b, Annelie M. Weinberg b, Minh Becker a, Bernhard Mingler c,Peter J. Uggowitzer a, Jörg F. Löffler a,⇑a Laboratory of Metal Physics and Technology, Department of Materials, ETH Zurich, 8093 Zurich, Switzerlandb Department of Orthopaedics and Orthopaedic Surgery, Medical University Graz, 8036 Graz, Austriac Health & Environment Department, AIT Austrian Institute of Technology GmbH, Biomedical Systems, 2700 Wr. Neustadt, Austria

a r t i c l e i n f o

Article history:Received 29 April 2014Accepted 10 September 2014Available online 18 September 2014

Keywords:A. MagnesiumB. SEM

a b s t r a c t

The biodegradation of ultrahigh-purity Mg (XHP Mg, Fe � 2 ppm) was tested in vitro in NaHCO3/CO2-buffered simulated body fluid and in vivo in the femur of rats. The in vitro degradation rate, which wasevaluated using an also here described newly designed hydrogen evolution testing setup, is with�10 ± 3 lm/y very low and shows very good agreement with in vivo data. The results for XHP Mg are alsocompared with in vitro tests on high-purity Mg (Fe � 37 ppm) in the as-cast and annealed states. The lesspure specimens exhibit significantly higher degradation rates due to the formation of Fe-containingprecipitates during casting and annealing.

� 2014 Elsevier Ltd. All rights reserved.

1. Introduction

Magnesium (Mg) offers great potential as biodegradable andbiocompatible material for medical applications [1–6]. Its corro-sion performance, however, is often insufficient and thus limitsits use. Within the living body, the generation of gaseous H2 andhydroxides (OH�) upon Mg degradation can influence the healingprocess of an injured tissue due to the formation of hydrogen pock-ets and localised increase in pH [1,7,8]. Therefore, slow and homo-geneous degradation of Mg implants is required.

The degradation behaviour of pure Mg is influenced by impurityelements such as Fe, Ni, Cu and Co. These impurities are detrimen-tal to corrosion resistance because of their low solid-solubility ina-Mg and due to their tendency to form cathodic sites for micro-galvanic corrosion in corrosive media [7,8]. The corrosion rate ofMg increases drastically when the concentration of Fe in Mg cast-ings exceeds its tolerance limit of 180 ppm. This drastic accelera-tion is due to the precipitation of an Fe-enriched body centeredcubic (BCC) Mg–Fe phase, which forms from the melt before finalsolidification [9,10]. Liu et al. [11] demonstrated that the corrosionrate in pure Mg increases with increasing Fe content. They alsoobserved a significant increase in degradation after heat treatment,which is a necessary processing step for the production ofMg-wrought alloys. The tolerance limit was found to be much

smaller in heat-treated pure Mg than in the as-cast material, dueto the fact that above a limit of 5–10 ppm Fe the BCC Fe-rich phasemay precipitate upon annealing [11]. Although high-purity (HP)Mg already exhibits very low amounts of impurities (see Table 1)and corrodes more slowly than conventional pure (CP) Mg, itsalloys still corrode too fast for specific medical applications. Toexplore a solution, ultrahigh-purity (XHP) Mg [12] was producedin-house via a distillation process to significantly lower the impu-rity level (see Table 1). One of the investigation’s main goals was totest the degradation performance of this new XHP Mg in vitro andin vivo, and thus to understand its degradation in more detail.

The in vitro degradation of the XHP Mg was analysed by thehydrogen (H2) evolution method [13] and compared to HP Mg.Current H2-evolution setups, however, show several limitationsas to their accuracy and ability to determine the corrosion rate,especially when testing very slow-corroding high-purity Mg alloys[13–15]. Therefore we designed within this work a new testingdevice, which enables precise determination of the evolved H2 vol-ume, including a regulation of the pH via NaHCO3/CO2-buffering. Inthis context, we present solutions to the drawbacks of currentH2-evolution setups, whose limitations were also recentlydescribed by Kirkland et al. [16]. To test the reliability of our setup,we also compare the in vitro degradation rate of XHP Mg with thein vivo degradation of XHP Mg pins implanted in the femurs of ratsfor 12 weeks.

In summary, the research aims of this study are to analyse andcompare the in vitro and in vivo performance of ultrahigh-purity

30 J. Hofstetter et al. / Corrosion Science 91 (2015) 29–36

(XHP) and high-purity (HP) Mg to understand Mg corrosion inmore detail. To achieve these goals, the in vitro degradation wasstudied with a newly designed hydrogen evolution testing setup.

2. Experimental methods

2.1. Materials and methods

HP Mg (99.99%, CHEMCO Germany) and three batches of XHPMg were used in the as-cast, annealed and as-extruded states tostudy the influence of impurity contents on in vitro degradationperformance. The degradation rate of extruded XHP Mg was alsoevaluated in vivo. The chemical compositions of the tested materi-als are given in Table 1.

The XHP Mg batches were produced via a purification process ina vacuum distillation apparatus [12]. In order to avoid contamina-tion, high-purity graphite crucibles with an ash content of about20 ppm were used for the production of the XHP material. The dis-tillation procedure resulted in a very low Fe content of 0.2–2.2 ppm(see Table 1), which is more than one order of magnitude less thanthat of the HP material.

The samples for a first series of experiments, used to evaluatethe corrosion properties in various conditions, were producedusing a simple casting and machining procedure. HP Mg and XHPMg (batch XHP#1) were cast to rods (10 mm diameter; length�100 mm) in a vacuum-induction furnace under argon atmo-sphere. From these rods, discs of 3.6 mm thickness and 10 mmdiameter were produced, generating a surface area of 2.7 cm2. AnXHP Mg billet in the as-distilled condition (diameter 50 mm,length 120 mm; batch XHP#3) was also extruded to a rod of6 mm diameter at 300 �C (extrusion ratio 1:69). From these rodstwo types of specimens were machined: samples for the in vitrotests with 6 mm diameter and 2 cm length, generating a surfacearea of 4.3 cm2; and pins for the in vivo tests with 1.6 mm diameterand 8 mm length, with a surface area of 0.44 cm2.

It is well known that the influence of trace elements on the cor-rosion behaviour of Mg and Mg alloys depends strongly on thethermal history. For example, as-cast HP Mg samples with45 ppm Fe exhibit a significant increase in corrosion rate after anannealing treatment at 550 �C [11], presumably caused by the pre-cipitation of Fe-rich BCC particles which act as cathodic sites. HPand XHP Mg samples were thus annealed for 48 h at 400 �C toinvestigate whether the corrosion resistance of Mg can beimproved via the proper adjustment and control of impurity levelsand heat treatments [9]. For these experiments, disc-shaped spec-imens (thickness 3.6 mm, diameter 10 mm) were machined fromthe HP Mg and XHP#1 batches and heat-treated in a resistance fur-nace under normal atmosphere (designated as HP400 andXHP400).

A second series of experiments was performed (i) to control theaccuracy of the setup; (ii) to test the reproducibility of the resultsfor XHP Mg; and (iii) to study the possible effect of different spec-imen geometries. Recently, Cao et al. [17] reported on the corrosionof our XHP Mg (XHP#2) in 3.5% NaCl solution. They observed local-ised corrosion and higher corrosion rates in some specimens. For

Table 1Impurity levels (in ppm) of high-purity (HP) and ultrahigh-purity (XHP) Mg samples. BatXHP#2 was characterised in its as-distilled condition. Batch XHP#3 was used in the as-ex

Mg quality Fe Si Mn Cu Ni

HP 37 11 9 <1 <1XHP#1 2.2 0.3 2.7 <1 <1XHP#2 1.4 0.9 2.2 <1 <1XHP#3 0.2 0.1 1.2 <1 <1

reasons of comparison, the specimens for our second series of deg-radation tests were taken from the same distillation batch(XHP#2). In these experiments the material was in the as-distilledcondition, i.e. no further casting or extrusion procedure was per-formed. In addition to the disc geometries deployed in the first ser-ies of experiments, cuboidal blocks in the as-distilled conditionwith a cross-section of 10 � 10 mm2 and a height of 5 mm wereused, generating a surface area of 4.0 cm2.

Prior to immersion testing, all sample surfaces were ground andpolished manually to remove defects and oxide layers. Grindingwas performed on abrasive SiC paper of granularities 2500 and4000; polishing was performed with a polishing cloth (1 lm dia-mond paste) and rinsed with isopropanol. Between the grindingand polishing steps, the specimens were cleaned in an ultrasonicbath using isopropanol and dried with hot air. The specimens forthe in vivo tests were additionally cleaned in a cascade of pure eth-anol in an ultrasonic bath, and again dried with hot air.

Elemental distribution maps of the intermetallic particles wererecorded using a Hitachi SU-70 scanning electron microscope(SEM, Schottky-type field emission gun) equipped with an X-Maxenergy dispersive X-ray (EDX) detector (Oxford instruments).

2.2. In vitro testing

The degradation performance of the pure Mg samples was eval-uated by immersing the polished discs in simulated body fluid(SBF) at 37 �C. SBF is an aqueous solution which simulates theion concentration and pH of human blood [18]. In this work theSBF was buffered with carbon dioxide (CO2, 99.90 vol.%) and thepH was kept constant at 7.45 ± 0.06. The ion concentration of theSBF used for the immersion tests is reported in Ref. [2], and wasused here without adding the buffers Tris or HEPES. In addition,no stock solutions were used, and the solid salt was added directlyto distilled water and heated to 37 �C. This was advantageousbecause as the pH was already in the range of about 7.5 after prep-aration neither diluted acid nor alkali had to be added. Approxi-mately 100 ml SBF per cm2 sample surface was used, which wellexceeded the minimum value required by ASTM G 31-72 (2004)[19], and three samples per batch and condition were measuredand evaluated. Many polymers are permeable to hydrogen [20],and thus glass equipment was used nearly exclusively. Abidinet al. [14] suggested saran wrap to control the atmosphere abovethe solution. In this work, however, several layers of large andsmall hollow polypropylene spheres were used, which significantlyreduced heat and water loss. With these spheres, only slight occa-sional refilling of the container with distilled water was necessaryto maintain the solution’s initial water volume.

Recently several authors have reported that the NaHCO3/CO2

buffer is a more appropriate pH buffer for simulating the in vivo sit-uation than the buffers Tris (tris(hydroxymethyl)aminomethane)or HEPES (N-2-hydroxyethylpiperazine-N0-2-ethanesulfonic acid)[14,21]. This statement was also confirmed by one of our previousstudies, in which a WZ21 alloy (Mg–2Y–1Zn–0.25Ca–0.15Mn, inwt.%) was tested in vitro with a prototype of the setup presentedhere and compared with its in vivo data [22,23]. In comparison to

ches HP and XHP#1 were tested in their as-cast and annealed states, whereas batchtruded state in vitro and in vivo.

Al Series no. of experiment Materials condition

17 1 as-cast, annealed1 1 as-cast, annealed

<1 2 as-distilled<1 1, in vivo as-extruded

Fig. 2. The Mg samples are positioned within a funnel on a sample holder (black)and beneath a burette. The increase in pH due to hydrogen evolution is controlledby a regulated CO2 inlet: gaseous CO2 enters a glass–ceramic frit and acidified SBF isproduced, which then passes through a ceramic filter into the glass container.

J. Hofstetter et al. / Corrosion Science 91 (2015) 29–36 31

other buffer systems, our NaHCO3/CO2 setup produced a good cor-relation with in vivo results [22,23].

The combination of this NaHCO3/CO2-buffer system with XHPMg generates very low hydrogen evolution, which requires a test-ing assembly of high precision. Therefore the prototype setup ofour previous study [22] was upgraded to overcome the many lim-itations reported in literature and to correct fluctuations in atmo-spheric pressure and gas solubility. The new setup stabilizes thepH at around 7.45 due to a regulated CO2 inlet and thus approachesin vivo conditions, and also enables precise evaluation of theevolved hydrogen gas.

Fig. 1 shows the setup, which comprises several devices such asthe glass container with the burettes and funnels; the polymeric(polycarbonate) plate; the heater with the aluminium plate andthe stir bar; the solenoid valve connected to a pH electrode andCO2 controller (threshold switch); and the frit to distribute theCO2 gas. We now describe these single components in more detail.

As shown in Fig. 2, each Mg sample is placed within a glass fun-nel (diameter of top = 5 cm) within a glass container. The samplesare fixed on POM MultiClips (Struers) sample holders [15]. Theseholders allow free flow of the solution around the sample. The bur-ettes are fixed to the clips with a Teflon band and with an O-ring tothe funnels to avoid any drifting of the samples. An intermixture ofthe solution in the vicinity of the sample is also guaranteed due tothe specially prepared funnels and the use of a stir bar at the bot-tom of the container. For this purpose, the two opposite sides of thefunnels have semi-circular openings (see semi-circles in Fig. 2 withopenings of approximately 0.5 cm in height and 3 cm in width).Glass burettes with a scaling of 0.1 ml are placed on the funnelsto collect the evolved hydrogen during the experiment. These bur-ettes are beneficial because a precise amount of SBF can be drawnout with the help of a stopcock to precisely regulate and evaluatethe level of solution (or hydrogen gas). The impermeability of thisassembly is ensured via an O-ring at the funnel-burette intersec-tion and by greasing the stopcock before every application.

This glass arrangement is positioned on a polycarbonate plate,which is fully covered with SBF. This specific polymer prevents apinning of bubbles, which may be formed at the beginning of orduring the experiment. Below this plate, a magnetic stir bar

Fig. 1. Schematic illustration of the immersion testing setup designed for slowlydegrading Mg samples in NaHCO3/CO2-buffered simulated body fluid (SBF). The Mgsamples are positioned beneath the burettes, which are fixed with O-rings andTeflon bands to the sample holders. A reference burette controls pressurefluctuations and gas solubility. The frit, the polycarbonate (PC) plate and the heaterwith a stir bar facilitate the homogenous CO2 distribution and keep the temperatureat 37 �C. The solenoid valve, the threshold switch, and the pH electrode control theCO2 inlet. The temperature is held constant using a thermocouple, and hollowpolypropylene spheres reduce the water evaporation.

agitates the solution in the glass container and keeps the temper-ature homogeneous throughout the setup. This is of crucial impor-tance for the degradation rate, as also stated in literature [24]. ThepH of the testing medium is measured using a pH electrode(Mettler Toledo, InLabExpertPro pH); see Fig. 1. The pH is held con-stant via a programmable threshold switch, which automaticallyopens a solenoid valve as soon as the pH reaches a lower thresholdvalue, and feeds gaseous CO2 into the solution. This feeding iscarried out until the pH reaches an upper threshold value. Thehomogeneous, fine distribution of the gas is controlled by aglass–ceramic frit (DURAN filter crucible with a maximum poresize of 10–16 lm); see Fig. 2. In contrast to other systems[14,17], this frit ensures a fine dispersion of the gas within thesolution. It also has the advantage that the dispersion of the gastakes place within the frit. This is possible because at the beginningof the experiment a small amount of the testing medium is auto-matically sucked into the frit. With the first opening of the valve,this fraction of solution is then acidified due to the direct inlet ofthe CO2 gas. As the experiment proceeds, this acidified solutionpasses through the frit into the testing solution regulating the pHin the glass container. This acidified solution makes it possible topremix the solution with CO2 and thus avoid any large gas bubblesnear the samples. Gas bubbles next to the samples are also avoidedby appropriate positioning of the frit just above the magnetic stirbar. The premixed solution is dragged down through a hole inthe polymeric plate above the stir bar and distributed evenly inthe glass container. A small lateral gap between the plate and theglass container on all four sides ensures the circulation of thedragged solution and hence a constant pH within the whole testingmedium.

Hollow polypropylene spheres of different diameters (10 and20 mm) cover the solution in the glass container to reduce evapo-ration of the liquid and to minimise heat loss. A homogeneous heatdistribution (of ±0.1 �C) in the entire container is achieved by theproper positioning of an aluminium (Al) plate between the glassand the heater (C-MAG HS7 IKAMAG). The only large temperaturedifference, of around 5 �C, can be detected between the bottom andthe top of the liquid column in the burette. The sizes of the glasscontainers can be varied and thus different quantities of batchesand specimens can be measured simultaneously. In this study, aglass container with five batches was used. In addition, referenceburettes were used for each measurement (see Fig. 1). These refer-ence burettes contain no samples and are taken as references in thecontext of fluctuation of atmospheric pressure and gas solubility.

Tests in the framework of this study have shown that it is cru-cial to consider the atmospheric pressure during the experiments,

Table 2Calculated volume content and calculated composition of the Fe-containing BCCphase and Mg2Si in as-cast HP, as-cast XHP, annealed HP400, and annealed XHP400.The as-cast nonequilibrium state was calculated according to Scheil.

Samples VBCC (mol%) VMg2Si (mol%) Element contentin BCC (wt.%)

HP (as-cast) 7 � 10�4 5 � 10�4 82Fe, 17Si, 1MnXHP#1 (as-cast) – – –HP400 (annealed) 2.2 � 10�3 1 � 10�4 83Fe, 16Si, 1MnXHP400 (annealed) 5 � 10�5 – 99.9Fe

32 J. Hofstetter et al. / Corrosion Science 91 (2015) 29–36

especially when the corrosion rate is small. Kirkland et al. reporteda correlation between the volume of the evolved hydrogen, the sur-rounding atmospheric pressure (the altitude at which the experi-ment takes place) and the samples’ mass loss [16]. In ourexperiments, the mass loss of the samples was negligible and onlythe influence of the atmospheric pressure within a burette at aconstant altitude was considered. The effect of atmospheric pres-sure on the solution level in the burettes and the influence of gassolubility are described in more detail in the discussion section.

2.3. In vivo testing

The animal experiments were conducted with ethical respectfor animals and were authorised by the Austrian Ministry of Sci-ence and Research (BMWF-66.010/0087-II/3b/2011). Six maleSprague–Dawley rats with body weights of 140–160 g and 5 weeksof age were used. Under general anaesthesia, each rat had twoidentical pins implanted into its femoral bones. Thus, 12 pins intotal were implanted. The surgical procedure and postoperativetreatment are described by Kraus et al. [23]. Micro CT scans wereperformed at four prearranged time points: 1 week, 4 weeks,8 weeks and 12 weeks after pin implantation. During the lCTexaminations the animals were anaesthetized using volatile isoflu-rane (Forane�, Abbot AG, Baar, Switzerland). The rats were scannedusing a Siemens Inveon Acquisition Workplace 1.2.2.2. Scans wereperformed at 70 kV voltage, 500 lA current, and 1000 ms exposuretime. The volume of each implant was measured by 3-D morpho-metric analysis and the pin volume was quantified using the soft-ware program ‘Mimics’ (Version 15.0, Materialise, Leuven,Belgium); for more details see Ref. [23]. To gain information aboutthe pins’ long-term behaviour none of the rodents were euthanisedfor weight-loss measurements and histological investigations. Thelong-term animal study is still running, and its final outcome willbe presented in a separate publication.

3. Results

3.1. Microstructure

For the first series of experiments, HP Mg and XHP#1 Mg werecast to rods to obtain samples of similar microstructure. Theseexhibit a typical cast microstructure with a grain size of approx.500 lm. The annealing treatment at 400 �C for 48 h did not resultin significant changes in the microstructural features observable byoptical microscopy. However, by means of scanning electronmicroscopy a few intermetallic particles with a size of typically150 nm were detected in the annealed HP400 specimens. Theintermetallic particles were sporadically distributed within thesample and contained Fe, Si and Mn at a ratio of about 84:15:1

Fig. 3. (a) HP400: EDX Fe Ka mapping of an Fe–Mn–Si particle w

(measured by EDX analysis; see Fig. 3a). This value fits very wellwith the calculated ratio given in Table 2. The extruded pins exhib-ited a completely recrystallized structure with equiaxed grains of�60 lm in size.

For the second series of experiments, the as-distilled conditionof XHP#2 Mg was used for microstructural investigations. Here thesamples exhibited a homogeneous microstructure with giantgrains of more than 1 mm in size. Very sporadically, small poresof diameter �7.5 ± 2.5 lm were detected (Fig. 3b).

3.2. Immersion tests

All Mg samples were immersed for around two weeks inNaHCO3/CO2-buffered SBF at 37 �C and a pH of 7.45 ± 0.06. Fig. 4shows the mean hydrogen evolution as a function of immersiontime for HP and XHP samples. Each data point was corrected forgas solubility and atmospheric pressure effects.

3.2.1. Mg of different purities and processing conditionsThe data for HP specimens show a steep increase of mean H2

evolution at the beginning of the experiment, followed by a steadyincrease (Fig. 4). The samples depict a few tiny local corrosionattacks after two weeks (Fig. 5). The curves of the XHP samples(as-cast XHP#1 and as-extruded XHP#3) also feature an initialsteep increase, but with a further slower rising. These samplesshow no local attacks, as was observed for HP Mg. The XHP sam-ples released approximately half the amount of hydrogen com-pared to the HP samples for the same immersion time (Fig. 4).

The annealed HP400 samples not only released more hydrogenthan their untreated counterparts HP (Fig. 4), but also exhibitedmore severe local attack (Fig. 5d). The gas amount for XHP400was nearly equal to that for XHP and no significant change in thesurface morphology was observed.

The average degradation rate, PH (mm/y), evaluated from thehydrogen evolution rate, VH (ml/cm2/d), using PH = 2.088VH [17],is calculated as roughly 10 ± 3 lm/y for XHP Mg, 28 ± 2 lm/y forHP Mg, and 39 ± 3 lm/y for HP400 Mg for an immersion time of20–320 h.

ith a size of �150 nm. (b) Micropore in as-distilled XHP#2.

Fig. 4. Mean hydrogen evolution as a function of immersion time for HP and XHPsamples in NaHCO3/CO2-buffered SBF at 37 �C. For the sake of clarity only datapoints after 14 days of exposure are shown for XHP#2 (2nd series; o discs, h cubes).The specimens HP400 and XHP400 were annealed for 48 h at 400 �C. Three samplesper batch and condition were tested and evaluated. The error bars for HP illustratethe typical data scatter. The volume loss results for the in vivo tests were transferredto equivalent hydrogen evolution and displayed as average in vivo degradation ratevia a straight line, including upper and lower limits (dashed lines).

J. Hofstetter et al. / Corrosion Science 91 (2015) 29–36 33

3.2.2. Accuracy and reproducibilityFig. 4 also shows the hydrogen evolution data for the second

series of experiments with XHP#2 in the as-distilled state and dif-ferent specimen geometries. For the sake of clarity, only datapoints after 14 days of exposure are displayed. While these dataindicate that XHP samples degrade at comparable rates indepen-dent of slight changes in trace element content (see Table 1), pro-duction procedure (‘as-cast’, ‘annealed’ and ‘as-extruded’ versus‘as-distilled’ condition), and sample shape (discs and rods versuscubes), there was a significant exception: one XHP#2 disc samplereleased large amounts of H2 and reached a hydrogen level ofapproximately 2.4 ml/cm2 after two weeks of immersion (notshown in Fig. 4). This is about 20 times more than the averageH2 release of discs and cubic samples. Localised corrosion attackin the form of a huge ‘volcano’ [25] was observed for this specimen(Fig. 5e).

3.3. In vivo results

Fig. 6 illustrates the results of the lCT scans of the in vivo deg-radation after 1, 4, 8, and 12 weeks post operationem. No significant

Fig. 5. Surface appearance after 2 weeks of immersion in NaHCO3/CO2-buffered SBF: no lHP and (d) HP400; and (e) formation of a large volcano on one XHP#2 specimen (outlie

degradation of the XHP pins is visible by eye. The pins’ surfaceremained smooth until week 12, indicating very homogenous deg-radation. An evaluation of the extent of degradation in terms of thepins’ volume loss (using ‘Mimics’ software) provided the datashown in Fig. 6e. The degradation increased nearly linearly overtime. As for the in vitro immersion results, the degradation ratewas very low. From the volume loss an average in vivo degradationrate of PIV = 13 ± 3 lm/y for the 3-month period can be calculated,which agrees very well with the in vitro studies.

4. Discussion

4.1. Mg of different purities and conditions

The hydrogen amount released by the XHP Mg samples withextremely low amounts of impurities is clearly smaller than thatreleased by HP Mg. Lee et al. [26] analysed the effect of impuri-ties on the corrosion behaviour of pure Mg. They stated that thecorrosion behaviour of pure Mg is dependent on the contentratio of two specific impurity elements (Fe/Mn), rather thanthe absolute value of each. This fact was also observed earlierin Mg–Al alloys [27]. Our Fe/Mn ratio is �4 for HP and <1 forXHP. In this sense, the results of our study accord qualitativelywith Ref. [26].

The annealed samples HP400 released larger hydrogen amountsthan their untreated counterpart HP. In comparison to theuntreated samples they also exhibited more local corrosion andthus more ‘H2 volcanos’ [25] on the surface. On the other hand,the XHP specimens showed no local corrosion attack, whereas afew volcanos were found on HP (Fig. 5). Liu et al. [11] investigatedthe tolerance limit for Fe in pure Mg. For the annealed state(550 �C/24 h) they proposed a value of 5–10 ppm Fe, above whichthe precipitation of an Fe-rich BCC phase was predicted. This phaseacts as an effective cathodic site and thus high corrosion rates areexpected for Fe-contents above this tolerance limit. Our experi-mental results verify this suggestion.

The two conditions ‘as-cast’ and ‘annealed’ differ in their ther-modynamic states; for the as-cast versions HP and XHP a nonequi-librium condition can be assumed, whereas HP400 and XHP400 areclose to equilibrium. By means of thermodynamic simulations wecalculated the constitution of the different states using the Pandatsoftware package with the database PanMg8 [28]. For the as-caststate, a nonequilibrium approximation according to Scheil–Gulliver was used (see Table 2). Fig. 7 displays the equilibriumresults for the annealed state and Table 2 summarises theequilibrium and nonequilibrium data.

ocalised attack in (a) XHP and (c) XHP400 samples; slight and moderate attack in (b)r).

Fig. 6. lCT reconstructions illustrating the in vivo degradation process of 12 identical XHP Mg pins in 2-D slices: (a) 1 week; (b) 4 weeks; (c) 8 weeks; and (d) 12 weeks afterimplantation into femoral bones of Sprague–Dawley rats. (e) Pin volume loss illustrating in vivo degradation. The insert in (e) illustrates the position of the pin, which isperpendicular to the longitudinal axis of the femoral diaphysis.

34 J. Hofstetter et al. / Corrosion Science 91 (2015) 29–36

The calculations for solidification in nonequilibrium accordingto Scheil illustrate a significant difference (see Table 2): Fe-containing BCC is formed in as-cast HP, but not in as-cast XHP.Accordingly, the degradation rate of HP is higher than that ofXHP. It is interesting to note that the Fe content of HP Mg is alsobelow the tolerance limit of 180 ppm for cast Mg [11] for whichthe BCC phase precipitates from the melt only for near-equilibriumcondition. In our case, the solidification rate was rather high andScheil approximation seems reasonable.

For the annealed specimens the equilibrium state is relevant. InHP400 with 37 ppm Fe, the iron-containing BCC phase precipitatesupon annealing at 400 �C by an amount of 2.2 � 10�3 mol% (Table 2and Fig. 7); note that the calculated ratio of Fe:Si:Mn correspondsvery well to the measured ratio (Section 3.1). Compared to the as-cast condition the volume content of cathodic sites is significantlyhigher and consequently the degradation rate is higher for HP400than for HP. The annealed alloy XHP400 with 2.2 ppm Fe alsodegrades no faster than the as-cast counterpart XHP (Fig. 4).Apparently, the very low amounts (5 � 10�5 mol%) of BCC precipi-tates in XHP400 do not affect the corrosion behaviour. Conse-quently, the Fe-tolerance limit for heat-treated alloys liessomewhere between 2 ppm and 37 ppm, which confirms thetolerance limit of 5–10 ppm for wrought alloys predicted by Liuet al. [11].

Fig. 7. Isopleth graph for Mg–11 ppm Si–9 ppm Mn–Fe calculated using the Pandatsoftware package [28]. The content of the BCC phase at 400 �C in XHP Mg iscalculated for 0.3 ppm Si and 3 ppm Mn.

Three discs and three cubic XHP#2-Mg specimens were testedto control the accuracy of the setup and the reproducibility ofour distilled material (second series of experiments; see Sec-tion 2.1). The experiments yielded almost no differences in hydro-gen evolution among the various sample geometries. Obviouslygeometry or sample shape do not affect the corrosion rate, andthe setup enables reproducible results for XHP samples. However,there was one disc sample with exceptionally pronouncedhydrogen evolution and localised corrosion attack, which wasobservable as a large volcano on the surface (Fig. 5e). This phenom-enon was not only observed in NaHCO3/CO2-buffered SBF here, butalso in 3.5% NaCl solution as reported by Cao et al. [17]. Theseauthors suggested localised differences in chemistry (impurityparticles or other particles) or slight differences of the samplesurfaces to provoke this gas increase. However, they also statedthat no chemical inhomogeneity was identified [17]. In thepresent work we propose the presence of micropores as the causeof such localised attack. As described in Section 3.1 (Fig. 3b), bymeans of optical metallography we were able to detect suchmicropores, which may have formed during the solidificationprocess. The shape and quantity of such porosity depends on manyfactors, such as the processing method and its design, thermal con-ditions and feeding parameters, and is usually found incastings [29]. Inside the pores (or surface pits) an autocatalytic cor-rosion cell may have formed [30–32], producing a seriouslocalised corrosion breakdown. This explanation is also supportedby the fact that no experimental outliers were observed ineither the pore-free wrought states of XHP#3, or in XHP Mg-alloys[33].

4.2. Comparison of in vitro and in vivo degradation of XHP Mg

The degradation rate of XHP Mg evaluated in vitro by immersiontests in NaHCO3/CO2-buffered SBF compares well with in vivo stud-ies where lCT was used to analyse the volume loss of pinsimplanted into the femurs of rats: In both cases comparable, veryslow homogeneous degradation was observed with calculatedaverage rates of PH � 10 ± 3 lm/y (in vitro; see Fig. 4) andPIV = 13 ± 3 lm/y (in vivo; see Fig. 6). Similarly, good agreementbetween in vitro and in vivo degradation was also obtained forthe alloy WZ21, which was tested with a preliminary version ofthe present testing setup [22]. Thus our investigations indicateclearly that in vitro degradation measured in NaHCO3/CO2-bufferedSBF correlates well with degradation results measured in vivo,when one attends to the sophisticated experimental details. Our

Fig. 8. Pressure and decrease in column height of the solution (plotted as ‘volume’)

J. Hofstetter et al. / Corrosion Science 91 (2015) 29–36 35

results regarding the buffer system agree with those of Kirklandet al. [21], who suggested that carbonate buffers under partialCO2 atmosphere are more appropriate for mimicking biocorrosionrates than are zwitterionic organic compounds. This is due to thefact that zwitterion-based buffers and carbonate buffers form dif-ferent corrosion layers which affect the behaviour of the hydratedoxide layer present on the Mg surface in aqueous solution [21].

As reported in Ref. [33], the consistency of in vivo and in vitroresults is not valid for every alloying system, because corrosionbehaviour is affected not only by the composition of the electrolytebut also strongly by the interaction of the alloying system with thespecific electrolyte. It can also be assumed that the in vivo situationcould be mimicked even better if proteins were added to thein vitro CO2-buffered system. The goal of future studies is thereforeto test the interactions between different alloying systems andelectrolytes, and to compare them with in vivo data.

as a function of immersion time. The representative pressures of the ‘empty’ and‘saturated’ experiments oscillate in a pressure range of 1011–1021 mbar during theimmersion time. An accumulation of the gases (H2 and CO2) shows a similarincrease in volume as the single gas components. A ‘saturated’ solution does notexhibit such an increase and is only dependent on the atmospheric pressure.

4.3. Immersion testing setup

4.3.1. Maintaining a constant pH valueAs mentioned earlier, Kirkland et al. [16] list the significant

advantages and limitations of current hydrogen evolution setupsof various research groups. Hydrogen measurements are mainlybeneficial because of their real-time character, but intermixing ofthe solution remains a problem (see Refs. [14,16,34]). Bringingthe funnel into contact with the beakers’ base represents a ‘closed’system and the pH thus increases locally; a pH-buffering systemwhich intermixes with the solution is therefore crucial in main-taining a constant pH level throughout the setup. Otherwise therapid OH� release during the Mg corrosion shifts the pH of themedium outside of the physiological range and thus falsifies thein vitro data. In this study, several new units to improve the immer-sion testing setup (see Section 2.2) were assembled to circumventthis and other limitations.

In general changing of the testing solution regularly [2] is time-consuming, and every change reduces the accuracy of pH con-stancy and gas volume. The pH value of Hank’s solution is reportedto have a particularly strong effect on the corrosion rate of Mg [35].For example, a decrease in pH from 7.4 to 6.8 generates a morethan tenfold increase in hydrogen evolution [35]. In our studythe pH was kept at around 7.45, even though it is known that anin vivo Mg implant always generates a local increase in pH duringthe initial period of implantation [36]. However, keeping the pHconstant during the whole experiment makes in vitro testing muchsimpler and eventually more predictive, although it does not rep-resent the biological situation precisely.

4.3.2. Towards predicting in vivo resultsThe new setup makes possible very precise analysis of the

evolved gas. Our observations using this setup have shown thatnot only does atmospheric pressure play an important role, butthat other parameters also influence the fluctuations in the bur-ettes, in particular at the beginning of immersion. Despite reportsthat the solubility of H2 is very low at ambient temperature andatmospheric pressure [37], we observed a gas solubility effect withour measurement setup. A few articles mention that the solubilityof H2 and CO2 is also dependent on the salt concentration and thetypes of salt [38]. The pH also has to be taken into consideration.With the dependence of gas solubility on the mixed salt solutionconcentration, the pressure, the temperature, and the pH, the sys-tem is clearly complex. In this study we therefore conducted fivesimple immersion tests to quantitatively investigate the solubilityof H2 and CO2 while taking into account atmospheric pressure. For

this purpose, three reference burettes and a container weredeployed with two litres of SBF at 37 �C.

The experiments were first performed with an ‘empty’ system(no Mg samples and no CO2 inlet). The results show ‘normal’ atmo-spheric pressure fluctuations (denoted as ‘empty’ in Fig. 8). In asecond experiment, marked as ‘Mg’ system, a piece of pure Mg(99.95%) was positioned beside the burettes and the pH was keptconstant with a diluted HCl solution. No decrease in the height ofthe solution column would be expected because the Mg sampleis not directly beneath the burettes. However, a rapid decrease inthe column heights in the burettes was observed within the firstfew hours. This decrease in column height is labelled as ‘volume’in ml in Fig. 8. A third experiment tested the ‘CO2’ system: noMg was in the solution and the pH was kept constant with gaseousCO2. Even though no H2 was produced, a decrease in column heightin the burettes was visible in the initial stage. In a fourth experi-ment, a ‘Mg–CO2’ system was tested by placing a Mg sample withapproximately the same mass and surface area as used in the sec-ond experiment beside the burettes and maintaining a constant pHlevel with CO2. In this case a doubling of the decrease in columnheight within the same period of time would be expected, but thiswas not observed. This case illustrates the manifold interactions ofthe various species in the solution. After reaching equilibrium, thecolumn heights in the burettes are dependent on atmosphericpressure only. This was confirmed by a fifth experiment, denotedas ‘saturated’, where SBF was saturated over five days with CO2

and Mg before performing a ‘Mg–CO2’ experiment. In comparisonwith non-saturated SBF, volume changes in the burettes were onlyobserved to result from atmospheric pressure fluctuations, but notbecause of gas solubility.

Therefore there are two ways to evaluate the data correctly: onthe one hand, the curves must be corrected by using reference bur-ettes which measure the initial volume decrease and the atmo-spheric pressure fluctuation; on the other, the solution may besaturated with CO2 and H2 prior to immersion testing. The formeris of course much simpler and requires only a subtraction of thereference data. This method was thus applied in this study. In sum-mary, the setup presented is much more precise than one made ofburettes or cylinders as a ‘closed’ system (beakers) with large pHfluctuations. Based on this, the new setup can predict in vivoresults with a level of reliability which has so far not been obtainedin other studies.

36 J. Hofstetter et al. / Corrosion Science 91 (2015) 29–36

5. Conclusions

Measurements of degradation rates of as-extruded ultrahigh-purity (XHP) Mg with Fe impurities of 0.2–2.2 ppm measuredin vitro in NaHCO3/CO2-buffered SBF agree well with in vivo testsperformed by implanting XHP Mg pins in rat femurs. In both casesvery low and homogeneous degradation at an average rate ofroughly 10 lm per year was evaluated. This very slow and homo-geneous biodegradation can be explained by very low amounts ofimpurities, such as Fe, Si, Mn, Cu, Ni, and Al (see Table 1), which actas cathodic sites in galvanic corrosion. In contrast, as-cast high-purity (HP) Mg (Fe � 37 ppm) shows significantly faster in vitrodegradation than as-cast XHP Mg due to the presence of Fe-richprecipitates. Annealed samples (at 400 �C/48 h: HP400 andXHP400) reveal the existence of an Fe tolerance limit of2–37 ppm Fe for Mg wrought alloys, in good agreement withthermodynamic predictions [11]. Very few specimens in theas-distilled condition show higher corrosion rates due to thepresence of microporosity.

The novel testing setup presented here enables the precise mea-surement of H2 evolution in slowly degrading Mg samples. Thetemporary CO2-inlet approach has various advantages over previ-ous setups: it keeps the pH value at around 7.45 and changes ofsolution can be avoided. The in vivo situation can be mimickedmuch better in the new setup than in Tris- or HEPES-buffered sys-tems [21,22]. Constant stirring minimises local pH increase aroundthe samples and keeps the pH globally constant. The setup exhibitsa high degree of reproducibility and its accuracy can be easilyimproved by using reference burettes to incorporate daily atmo-spheric pressure fluctuations and dissolution changes for both H2

and CO2 in the SBF at the initial phase of the immersion tests.

Acknowledgements

The authors appreciate support via the K-project OptiBioMat(Development and optimization of biocompatible metallic materi-als), FFG – COMET program, Austria, and the Laura Bassi Centre ofExpertise Bioresorbable Implants for Children (BRIC), FFG, Austria.We are also grateful to Christian Wegmann and Manuel Spielhoferfor fruitful discussions and experimental assistance.

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