The role of the beta-Mg17Al12 phase on the anomalous hydrogen evolution and anodic dissolution of AZ magnesium alloys

Journal Pre-proof The role of the beta-Mg17 Al12 phase on the anomalous hydrogen evolution and anodic dissolution of AZ magnesium alloys C. Ubeda (Investigation)Writing-Original Draft), G. Garces (Resources)Writing-Review and Editing), P. Adeva (Resources) (Investigation)Writing-Review and Editing), I. Llorente (Investigation)Writing-Review and Editing), G.S. Frankel (Conceptualization)Writing-Review and Editing), S. Fajardo (Conceptualization) (Methodology) (Investigation)Writing-Review and Editing) (Supervision) (Project administration) (Funding acquisition)

PII:

S0010-938X(19)32217-6

DOI:

https://doi.org/10.1016/j.corsci.2019.108384

Reference:

CS 108384

To appear in:

Corrosion Science

Received Date:

29 October 2019

Revised Date:

5 December 2019

Accepted Date:

9 December 2019

Please cite this article as: Ubeda C, Garces G, Adeva P, Llorente I, Frankel GS, Fajardo S, The role of the beta-Mg17 Al12 phase on the anomalous hydrogen evolution and anodic dissolution of AZ magnesium alloys, Corrosion Science (2019), doi: https://doi.org/10.1016/j.corsci.2019.108384

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The role of the beta-Mg17Al12 phase on the anomalous hydrogen evolution and anodic dissolution of AZ magnesium alloys

C. Ubedaa, G. Garcesa, P. Adevaa, I. Llorentea, G.S. Frankelb and S. Fajardoa,c,* a b

National Center for Metallurgical Research (CENIM-CSIC), Madrid 28040, Spain

Fontana Corrosion Center, Department of Materials Science and Engineering, The Ohio

Department of Chemistry, University of La Laguna, La Laguna, Tenerife E-38200, Spain

Highlights

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c

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State University, Columbus, Ohio 43210, USA

Anomalous HE observed for AZ31 and AZ91 Mg alloys

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Passivity exhibited by Mg17Al12 with anomalous HE originating at anodic regions

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Mg17Al12 phase exhibited lower HE rates than AZ Mg alloys during polarization

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Results consistent with anomalous HE originating at actively dissolving Mg sites

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Greater amounts of Mg17Al12 decrease dissolution rates of Mg matrix.

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Abstract

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The role of the Mg17Al12 second phase on the anomalous HE and anodic dissolution of AZ31 and AZ91 Mg alloys was investigated. A passive-like behavior was exhibited by the Mg17Al12 phase with strong anomalous HE originating from local anodic regions where the

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passive film had broken down. The Mg17Al12 exhibited lower HE rates than the AZ Mg alloys during anodic polarization. Surface characterization confirmed that dissolution of Mg17Al12 in the AZ91 alloy only occurred above the Eb for the Mg17Al12. Under anodic polarization, greater amounts of Mg17Al12 are beneficial for AZ Mg alloys durability, decreasing the Mg matrix dissolution rates.

Key Words: Mg alloys; hydrogen evolution; anodic dissolution; beta phase (Mg17Al12); negative difference effect (NDE) *Corresponding author: [email protected], [email protected] (S. Fajardo).

1. Introduction

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Magnesium (Mg) is the lightest structural metal. It has a density of 1.7 g/cm3, which

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is 1.6 and 4.5 times lower than those of aluminum (Al) and steel, respectively [1]. For this reason, Mg alloys are materials of great technological interest for lightweight applications.

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However, Mg shows deficient mechanical properties. For this reason, Mg is normally alloyed

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with other elements. Among the existing Mg alloys, the Mg-Al-Zn (AZ) alloys are the most commonly used. Alloying with Al and zinc (Zn) increases metal hardness, strength and

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castability [2]. However, whilst alloying with Al and Zn increases strength and ductility, a wider use of the AZ Mg alloys has been limited due to their high reactivity in aqueous

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environments compared to other common structural materials. In the AZ Mg alloys, increased concentrations of Al in the alloy promote the enrichment of oxidized species of Al in the surface film, either forming Al2O3 and Al(OH)3 or as a mixture of Mg–Al oxides, such as the spinel MgAl2O4, which enhance corrosion

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resistance [3-5]. However, the low solubility of Al in Mg at room temperature results in the precipitation of second phases when the concentration of Al exceeds ~3 wt.% [2]. In the AZ series, the main Al-containing precipitate is the Mg17Al12 phase (β-phase). The volume fraction of the Mg17Al12 phase in the AZ Mg alloys depends on the time-temperature history of the alloy [2], as well as on the amount of Al in the material. Increased additions of Al above the solubility limit leads to a higher extent of Mg17Al12 phase. The Mg17Al12 phase

exhibits a more noble corrosion potential than Mg, acting as a local cathode during corrosion of the AZ Mg alloys under open circuit conditions [6, 7], as first rationalized by Lunder et al. [8]. Even though the reported effects of Al on the corrosion resistance of the Mg-Al alloys are contradictory [1], a myriad of studies on the role of the Mg17Al12 phase in the corrosion of the AZ Mg alloys have posited that this phase is detrimental for their corrosion resistance at the open circuit potential [9-17]. In brief, greater amounts of the Mg17Al12 phase will result

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in larger regions with the ability to sustain the cathodic reaction, thus enhancing the

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dissolution kinetics of the α-Mg matrix. However, it has also been reported that, in the absence of external polarization and when the volume fraction of Mg17Al12 phase in Al-

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containing alloys is high, a barrier effect is promoted, which inhibits corrosion propagation

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in the alloy [13, 18-20]. This is evidence for the lack of consensus on the effect of the Mg17Al12 phase on the degradation of unpolarized AZ Mg alloys.

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The primary cathodic reaction on Mg alloys under open circuit conditions is the hydrogen evolution (HE) reaction (HER). This is due to their low corrosion potential (Ecorr),

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which is far below the reversible potential for HE (Erev,H). In particular, the AZ Mg alloys exhibit Ecorr values that range from -1.6 to -1.4 VSCE [21]. This creates a large overpotential for the HER, with cathodic reaction rates that are not restricted by the diffusion-limited oxygen reduction reaction. Furthermore, the kinetics of the HER on Mg do not follow the

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expectations of electrochemical kinetics (as exemplified by the Butler–Volmer equation), which describe how the rate associated with the cathodic reaction should decrease with increasing anodic polarization [22]. Instead, the rate of HE increases with increasing anodic polarization of Mg at potentials above the Ecorr. Even though the origins of this anomalous reaction remain unclear, a number of different hypotheses have been offered to explain the enhanced rates of HE on anodically polarized Mg alloys. Briefly summarized, it has been

proposed that this behavior may be due to the catalytic properties of the corrosion products formed during dissolution towards the HE reaction [23-26]; accumulation of impurities more noble than Mg as a consequence of preferential dissolution of the matrix that may act as preferential cathodic sites [27-30]; chemical decomposition of MgH2 formed during Mg polarization [31]; and re-deposition of metal impurities present in the Mg electrode released non-faradaically to the electrolyte after free corrosion [32-34]. However, it has been shown

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that these theories are not able to account for the total anomalous HE rates measured during

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anodic polarization [35-39]. In fact, experimental evidence has been provided that the vast majority of anomalous HE originates at the actively dissolving anodic regions [31, 35-42].

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Furthermore, a mechanism for the anomalous HE reaction on dissolving Mg was recently

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proposed using a mechanistic surface kinetic density functional theory (DFT) model [43]. This model is in good agreement with several experimental observations.

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The AZ Mg alloys also exhibit the anomalous evolution of H2 under anodic polarization [13, 18, 44, 45]. Interestingly, the role of the Mg17Al12 phase on the anomalous

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HE of the AZ Mg alloys has not been previously investigated. This is remarkable considering that this parasitic reaction has significant implications in understanding the anodic dissolution kinetics of the alloys. To date, few studies have investigated the electrochemical response of the Mg17Al12 phase subjected to anodic dissolution, and the work of by Sudholz

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et al. [46] remains the most complete. However, in that paper only the Ecorr and a single potentiodynamic polarization curve of the Mg17Al12 phase was reported. The aim of this paper is to clarify the role of the beta-Mg17Al12 phase on the

anomalous HE and anodic dissolution reactions of the AZ magnesium alloys to fully understand the effect of this intermetallic particle in the durability of the AZ Mg alloys subjected to anodic polarization.

2. Materials and methods Commercially available AZ31 and AZ91 Mg alloys supplied by Magnesium Elecktron Ltd. were used as test specimens. Table 1 shows the chemical composition of the Mg alloys. A master ingot of single-phase intermetallic Mg17Al12 was prepared by induction melting using high purity Mg and Al with a nominal composition of 44 wt.% Al. The cast

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ingot was homogenized at 395 ºC for 24 h and cooled at a slow rate inside the furnace.

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Chemical composition of the resulting Mg17Al12 phase as determined by inductively coupled plasma atomic emission spectroscopy (ICP-AES) is also listed in Table 1.

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Test specimens in the form of plates of the AZ31 and AZ91 Mg alloys, and the

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Mg17Al12 phase were used for potentiodynamic polarization measurements whereas samples for hydrogen collection measurements were cold mounted in epoxy resin.

Electrical

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connection was achieved by attaching a copper wire to the rear part of the metal prior to mounting. The test specimens were ground using silicon carbide emery papers to a 1200 grit

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finish under ethanol prior to testing in all cases. The test solution was 0.1 M NaCl (pH∼6). All solutions were prepared from laboratory grade reagents and with high purity water with resistivity of 18.2 MΩ cm (Millipore™ system). Potentiodynamic polarization measurements in 0.1 M NaCl solution were performed

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using a flat electrochemical cell by scanning in the anodic direction from ‒30 mV to +700 mV vs. the open circuit potential (OCP) at a scan rate of 1 mV/s. Hydrogen volume collection measurements were carried out galvanostatically by passing a constant charge density of 5 C/cm2 at different applied anodic current densities (from 1 to 70 mA cm‒2) using the gravimetric method. This method, originally proposed by Curioni [24] and further developed by Fajardo and Frankel [47] allows for determination of the HE rates with an extremely high

accuracy by measuring the hydrostatic force exerted by the volume of electrolyte displaced by the evolved gas that accumulates in a submerged container. Complete details of the operating principles of the gravimetric method are given elsewhere [47]. All electrochemical measurements were performed using a three-electrode configuration with the Mg based material acting as the working electrode, a silver/silver chloride electrode (SSC) as the reference electrode and a large area Pt mesh as the counter

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electrode. A Gamry Instruments Interface 1000E potentiostat/galvanostat controlled by the

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Gamry Framework software was used to polarize the specimens. In all cases, the OCP of the samples was monitored for 10 minutes until a stable value was observed prior to polarization.

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All experiments were repeated at least in triplicate to check reproducibility.

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Microstructural inspection of the Mg17Al12 phase, AZ31 and AZ91 Mg alloys was performed using optical microscopy (OM), and elemental analysis of the microstructural

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features in the specimens was conducted using a Hitachi S 4800 scanning electron microscope (SEM) equipped with an Oxford Instruments energy dispersive X-ray (EDX)

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microanalyzer. The AZ31 alloy was etched using an acetic-picral reagent (5 g picric acid, 7 mL acetic acid, 10 mL deionized water and 140 mL ethanol), whereas the AZ91 alloy was etched with nital reagent (2 mL nitric acid and 98 mL ethanol) [48]. The surface fraction of the Mg17Al12 phase present in the AZ91 Mg alloy was estimated using image analysis

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(ImageJ® software). To obtain reliable results, surface analysis was carried out in 5 optical micrographs randomly selected corresponding to different regions of the same specimen. Even though this method presents some limitations in discriminating the possible different phases present in the surface, it represents a reliable tool for the estimation of the surface fraction of the Mg17Al12 phase. The images and video of the surface appearance of the Mg17Al12 phase during potentiodynamic polarization were acquired using a digital stereo

microscope (Motic DM-143-FBGG). X-ray diffraction (XRD) was performed using a Siemens D-5000 diffractometer equipped with monochromatized Cu Kα radiation. Patterns were recorded from 10 to 120 2θ degrees, in the step-scanning mode, with a 0.03° (2θ) step and a counting time of 4 s/step. XPS measurements were carried out using a VG Microtech model MT 500 spectrophotometer equipped with a non-monochromatic MgKα1.2 anode Xray source operating at 300W. A pressure below 10‒9 torr was maintained during data

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collection. Adventitious C was used as binding energy (BE) reference, with 285.0 eV for

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C1s.

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3. Results and discussion

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3.1. Microstructural characterization

Fig. 1 shows SEM micrographs of the AZ31 and AZ91 Mg alloys. Specimens were

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previously etched to reveal the presence of microconstituents, in particular the Mg17Al12 phase. As observed in Fig. 1, the microstructure of the AZ31 Mg alloy consisted of α-Mg

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grains with an average grain size of about 20 μm with no evidence of the Mg17Al12 phase in the alloy. In contrast, the AZ91 Mg alloy showed a dendritic α-Mg matrix microstructure with a considerable amount of Mg17Al12 phase. As previously explained, this difference is due to the low solubility of Al in Mg, where alloys containing above ~3 wt.% Al form

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intermetallic secondary phases (mainly Mg17Al12 in the Mg-Al series) [2]. Although not shown in Fig. 1, a small amount of Mn-contaning particles (likely Al-

Mn and Al-Mn-Fe intermetallic phases) were also observed in the AZ31 and AZ91 Mg alloys, exhibiting elongated or rounded regular shapes and a diameter of about 5 µm. Finally, the AZ91 Mg alloy showed small colonies of lamellar particles (indicated in Fig. 1 as α + β),

which have been considered to be the partially divorced eutectic α-Mg/Mg17Al12 resulting from non-equilibrium solidification during casting [48-51]. 17121712.

The surface fraction of the Mg17Al12 present in the AZ91 Mg alloy was

estimated on a previously etched AZ91 Mg alloy specimen to aid in Mg17Al12 phase identification. Fig. 2 shows a typical optical micrograph used for the analysis. According to the microstructural characterization shown in Fig. 1, the Mg17Al12 phase is associated with

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the dendritic regions on the surface. The calculated area fraction of the Mg17Al12 phase was

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14 ± 2 %. This value is slightly lower that that reported in a previous work by Majhi and Mondal [52] and Mingo et al. [48], where area fraction of the Mg17Al12 phase was determined

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to be 16.7 % and 19.8 ± 1.3 %, respectively. Nevertheless, the calculated area fraction of the

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Mg17Al12 phase in this study is of the same order and can be considered reliable. Fig. 3a shows the XRD pattern of the intermetallic Mg17Al12 phase. All peaks

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corresponded to the Mg17Al12 phase and no secondary phases were detected by XRD. Furthermore, the EDS elemental maps depicted in Fig. 3b show that Mg and Al were

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homogenously distributed over the scanned area, confirming that the processed intermetallic material was single phase.

3.2. Potentiodynamic polarization measurements

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Potentiodynamic polarization curves were performed on the Mg17Al12 phase and the

AZ31 and AZ91 Mg alloys to assess the electrochemical kinetics both under cathodic and anodic conditions. For simplicity, one typical curve will be presented for each material. Fig. 4 shows the potentiodynamic polarization curves for the Mg17Al12 phase, AZ31 and AZ91 Mg alloys in 0.1 M NaCl solution. AZ31 and AZ91 Mg alloys exhibited similar corrosion potential (Ecorr) values, although the Ecorr values of AZ91 were slightly lower. At potentials

below the Ecorr, the AZ31 sample exhibited higher cathodic kinetics than the AZ91 alloy, which is in agreement with previous works [45]. Interestingly, the cathodic current densities measured on the Mg17Al12 phase suggest that this intermetallic phase exhibits higher catalytic properties towards the HER than the two AZ Mg alloys studied. Detailed investigation of the cathodic kinetics for these materials is out of the scope of this study, but it will be part of future work.

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The anodic polarization curve of AZ91 Mg alloy exhibited what might be interpreted

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as a small passive region followed by current increase, whereas the AZ31 specimen exhibited an increase in current immediately upon polarization above its lower Ecorr. As a result, the

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dissolution rates of AZ91 were lower than AZ31 at a given potential. This behavior is

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expected due to the higher concentration of Al in the alloy, which promotes a more protective surface film thus lower anodic kinetics. More importantly, the Mg17Al12 phase was

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spontaneously passive, with the current density on the order of a few A/cm2 over a potential range of about 400 mV above Ecorr. After a critical potential value was reached, a sharp

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increase of the current density was observed. This is normally attributed to the disruption of the protective passive film and indicates the onset of localized corrosion. Visual inspection of the surface during polarization indicated the concomitant appearance of pit-like as well as filiform-like attack, as will be shown below. The breakdown potential (Eb) for the Mg17Al12

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phase was around ‒0.950 VSSC.

To confirm the passive nature of the surface on the Mg17Al12 phase, EDS element

point analysis was carried out on a line scan from the bulk metal to the outermost region of the surface on a crossed-sectioned Mg17Al12 specimen. Prior to cross-sectioning the sample, the film was potentiostatically grown on the intermetallic surface by applying a constant

anodic potential at ‒1.2 VSSC in 0.1 M NaCl solution for 30 min. Potentiostatic polarization was performed to aid in identification and analysis of the film which, otherwise, may be too thin to be clearly detected. Please note that at this potential value the surface exhibited a passive-like behavior (see Fig. 4), thus only the growth of the surface film is expected. Fig. 5a shows the BSE-SEM micrograph of the crossed-sectioned surface where 10 EDS measurements were consecutively carried out at a constant separation of 1 µm (marked with

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a white cross in Fig. 5a. This procedure was performed in triplicate to obtain reliable results.

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Fig. 5b shows the Al/Mg atomic ratio as a function of the distance from the surface obtained by EDS analysis according to the protocol explained above. To aid in comparison, the

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stoichiometric ratio expected for the Mg17Al12 phase is also shown. It is evident from Fig. 5b

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that the surface was enriched in aluminum. This aluminum enriched region was about 2 µm thick and exhibited increased concentrations of Al towards the outermost region of the

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surface. The enrichment of Al is consistent with Mg being preferentially dissolved during anodic polarization due to its higher reactivity.

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The passive layer formed on the Mg17Al12 phase after potentiostatic polarization was also studied using XPS. Fig 6 shows the O1s, Mg 2p and Al 2p high resolution XPS spectra for the AZ31 and AZ91 alloys, and Mg17Al12 specimens after potentiostatic polarization at ‒ 1.2 VSSC in 0.1 M NaCl solution for 30 min. The O1s signal can be fitted to two components

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at a binding energy of 529.9 eV, associated with the presence of metallic oxides, and 531.8 eV, associated with the presence of oxygen in hydroxide form. The high resolution spectrum of Mg2p shows a main component at a binding energy of 50.5 eV, which is typically found in magnesium in its oxidized form Mg2+ (MgO and Mg(OH)2) and a component located at 48.3 eV associated with the presence of magnesium in its metallic state. Finally, the Al2p spectrum consists of two components that include the peak at 74.7 eV corresponding to

oxidized Al3+ (oxide and/or hydroxide) and a less intense peak at 72.6 eV corresponding to metallic Al. The O1s peak main component located at 531.8 eV (hydroxides) suggests that the film was mainly composed of an Al-Mg hydroxide. This is in perfect agreement with a previous work where it was shown that during immersion of the Mg17Al12 phase in aqueous environments, a surface film forms of about 10 nm thick and consisting of AlMg2.5(OH)8 [53]. This surface film provides the protective properties of Al and Mg in the neutral and

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alkaline pH ranges, respectively. The increased concentrations of Al towards the outermost

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region of the surface and the reported existence of a stable Al-Mg hydroxide layer are consistent with the passive-like behavior shown by the intermetallic Mg17Al12 phase in the

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potentiodynamic curves (Fig. 4).

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Interestingly, strong anomalous HE was observed during potentiodynamic polarization of the Mg17Al12 phase. Streams of H2 were observed at discrete sites on the

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surface after the breakdown potential was reached, likely the sites of pitting corrosion. Furthermore, the HE rate at the pits increased with increasing anodic polarization. Fig. 7

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shows a series of optical images of the dissolving Mg17Al12 phase electrode surface during an anodic potentiodynamic polarization experiment. These optical images labeled A to D are associated with points on the polarization curve that were selected because they correspond to representative stages during anodic polarization: A, the Ecorr; B, passive region; C, small

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polarization above Eb; and D, large polarization above Eb. Recording of the experiment was carried out at a 20X magnification. The surface was polished to 1 µm to assist the capture of images without undesirable light reflections, and covered with tape so that only a circular surface area with a diameter of about 6 mm was exposed to the electrolyte. The experiment started at the OCP after the sample reached a steady state potential. H2 bubbles formed on the surface immediately after the electrolyte covered the specimen

(prior to polarization) due to free corrosion of the metal because HE is the main cathodic reaction during Mg corrosion. Furthermore, a visible corrosion film partially covered the surface of the Mg17Al12 phase at the OCP. After anodic polarization progressed to the passive region (point B), the corrosion film developed with a more visible accumulation of corrosion products, as indicated by an enhanced contrast with the unfilmed region. However, no significant changes were observed in the amount of surface covered with this corrosion

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product film, or in the amount of H2 bubbles. Most of the H2 bubbles originated during the

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very early stages of immersion and remained unchanged until the Eb was reached.

At potentials above Eb, the current increased sharply (point C) due to localized attack.

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This was accompanied with the appearance of streams of H2 gas that emerged from fixed

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locations that were presumably the sites of pits (optical image C). Finally, as the applied potential increased further, greater amounts of H2 were produced at the local anodes (optical

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image D). Analogously, Figs. 8 and 9 show the optical images of the dissolving AZ31 and AZ91 Mg alloys, respectively, during anodic potentiodynamic polarization. The AZ31

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exhibited local anodic regions shortly after the onset of polarization with simultaneous anomalous HE originating from the same sites. In contrast, the AZ91 did not show any HE until potentials above the small passive region (see optical image C in Fig. 9). In that case, after disruption of the protective film, local dissolution with concomitant streams H2 gas were

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observed. For both Mg alloys, the rate of HE increased with polarization. This behavior can be explained using the interpretations by Fajardo and Frankel [35, 37, 38], who have shown that the anomalous HE exhibited by dissolving Mg occurs primarily at the actively dissolving regions. The strong anomalous HE exhibited by the Mg17Al12 phase has important implications on the dissolution kinetics of the AZ Mg alloys that contain this intermetallic second phase, as will be explained in detail below.

Fig. 10 shows the surface appearance of the AZ31 and AZ91 Mg alloys and the Mg17Al12 after the potentiodynamic polarization measurements shown in Figs. 7-9. Differences are readily observed for each material. The AZ31 Mg alloy exhibited a heavily corroded surface with the typical black corrosion products normally observed in Mg and Mg alloys that extended radially on the specimen. On the other hand, the AZ91 Mg alloy showed a significantly less corroded surface, with a number of localized corroded regions distributed

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on the sample. Note that a slight crystallographic contrast is observed in the exposed surface,

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suggesting the presence of two different phases. The corroded regions appear to follow the tracks determined by the interdendritic darker phase. According to Fig. 1 and due to its higher

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reactivity, it is likely that this darker phase corresponds to the α-Mg matrix. It is reasonable

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to assume that a greater aluminum content in the material leads to a higher resistance to localized corrosion, associated with the formation of a protective layer on the metal surface.

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These results are in agreement with the lower anodic kinetics shown by the AZ91 Mg alloy (see Fig. 4) and were previously observed in a recent study using the same materials [45].

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Furthermore, a slightly darker color was visually detected in all the AZ91 exposed surface, which is consistent with the presence of a surface film. Finally, clear localized attack was also evident for the Mg17Al12 specimen, which is consistent with the potentiodynamic polarization curve shown in Fig. 4. In this case, pit-like and elongated corroded regions were

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observed. Interestingly, some filiform tracks that interconnected these pit-like regions were also detected. In summary, whilst the AZ31 Mg alloy exhibited a more uniform corrosion attack, localized corroded regions consistent with the presence of a protective surface film were observed on the AZ91 and the Mg17Al12 phase. A closer look at the corrosion products shown in Fig. 10 is provided in Fig. 11, where SEM micrographs of the surfaces of the AZ31, AZ91 and Mg17Al12 phase are shown. The

AZ31 Mg alloy exhibited a dense corrosion product layer with abundant cracks. In contrast, the SEM image of a filament track on the AZ91 showed a more uniform and continuous layer, with fewer cracks and a more compact appearance. Finally, corroded areas in the Mg17Al12 phase exhibited a completely different morphology, with no presence of detectable deposition of corrosion products. Instead, a clear localized attack that revealed a very high density of pits was observed. Evident orientation of the remaining metallic structure after

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polarization indicates that dissolution followed crystallographic pitting. The presence of

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crystallographic pits is consistent with the notion that the solution inside the pits was not saturated (likely due to concomitant HE) [36], explaining the absence of corrosion product

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deposition. Elemental analysis obtained by EDX on a number of sites on the corroded areas

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is listed in Table 2. It is observed that corrosion products on the ZA31 and AZ91 Mg alloys consisted mainly of Mg and O at a Mg:O ratio of 1:2, which is consistent with the presence

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of Mg(OH)2. A small signal of Cl, probably from the entrapment of chloride ions from the electrolyte in the corrosion product film, was detected for the AZ31 and AZ91. No

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enrichment of Al in the corrosion product was observed for any of the Mg alloys studied. Interestingly, the EDX analysis carried out on attacked regions of the Mg17Al12 phase showed that the concentration of O was lower, and that, besides oxygen, only Mg and Al were detected. The ratio of Mg:Al in the corroded areas was 1.3 ± 1, which is nearly identical to

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the nominal stoichiometric ratio of Mg and Al for the Mg17Al12 phase (~1.4). This confirms that after breakdown of the passive film, the matrix was exposed to the electrolyte and dissolved locally with simultaneous HE, mixing the pit and bulk environments and hindering deposition of corrosion products [36].

3.3. Hydrogen collection measurements

The rates associated with the anomalous HE on each of the materials were studied using the gravimetric method for hydrogen evolution. All experiments were performed galvanostatically by applying different anodic current densities to a constant net charge density of 5 C/cm2. Figs. 12-14 show the volume of H2 collected during the application of the anodic current densities as a function of time for the AZ31, AZ91 and Mg17Al12 phase specimens in 0.1 M NaCl solution, respectively. All experiments were repeated at least in

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triplicate to ensure reproducibility. A linear-like behavior was shown for every applied

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current density, which indicates that the rate of HE during polarizations was constant in all cases. It is evident from Figs. 12-14 that strong anomalous HE manifested in all the materials

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studied. While this behavior has been described in previous works for the AZ31 and AZ91

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Mg alloys [6, 45], it is the first time that this phenomenon is reported for the intermetallic Mg17Al12 phase.

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It is possible to calculate the rate associated with anomalous HE by the linear fit of the HE curves in Figs. 12-14, which can be expressed in terms of current density values after

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conversion using Faraday’s Law [47]. Fig. 15 shows the calculated HE current densities increased with increasing applied net anodic current density in 0.1 M of NaCl solution. Furthermore, the anomalous HE rate decreased with increasing Al content in the materials. This behavior is in agreement with a previous work on the anomalous HE on different AZ

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Mg alloys [45], and provides solid evidence that the amount of hydrogen evolved during dissolution of these alloys is related to the Al content in the metal. Fig. 15 also shows that under galvanostatic polarization, the Mg17Al12 phase exhibits strong anomalous HE, but at rates lower than those shown by the AZ Mg alloys. The charge density related to hydrogen evolution normalized to the net charge density applied is shown in Fig. 16 as a function of applied anodic current. This charge ratio was almost independent of current density but

decreased with increasing aluminum concentration. The HE ratios of about 40 and 30% for AZ31 and AZ91, respectively are lower than those previously reported for Mg of different purities, where the HE charge ratio was about 50% [35]. The Mg17Al12 phase exhibited the lowest values for the HE charge, confirming that this intermetallic phase is less prone to suffer from this parasitic reaction than the AZ Mg series. It was hypothesized previously [45] that, at the typical potentials observed during

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anodic polarization of the AZ Mg alloys (above the Ecorr for the Mg17Al12 phase), hydrogen

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may either diffuse into the intermetallic particle [6] or contribute to the net applied current density due dissolution of the Mg17Al12 phase, decreasing the α-Mg dissolution and

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anomalous HE current densities. However, no evidence was provided. Furthermore, the

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anomalous HE exhibited by the Mg17Al12 phase was unknown. From the results in this paper, a complete description of the role of the intermetallic Mg17Al12 on the anomalous HE and

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anodic dissolution of AZ magnesium alloys can be realized.

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3.4. Role of the Mg17Al12 phase on the anomalous HE and anodic dissolution of AZ Mg alloys Previous observations have determined that the anomalous hydrogen evolution exhibited by Mg and Mg alloys is primarily related to the regions on the surface undergoing active dissolution [35-39]. Furthermore, a surface kinetic DFT model was recently presented

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clarify the mechanism of anomalous HE on anodically polarized Mg, where it was shown that Mg dissolution and anomalous HE proceed at the same active sites on the reacting surface [43]. This theory is consistent with the results shown for the intermetallic Mg17Al12 phase, where anomalous HE only occurred after breakdown of the passive film (highly enriched in aluminum). In this case, Mg was exposed to the electrolyte leading to anomalous HE within the actively dissolving pits.

In the case of the AZ31 Mg alloy, where no Mg17Al12 phase was detected, it is predicted that polarization applied is exclusively associated with α-Mg dissolution. However, for the AZ91 Mg alloy, the presence of Mg17Al12 would result in a lower surface fraction of α-Mg subjected to anodic polarization. Considering that below the Eb for the Mg17Al12 phase no other phase than the α-Mg matrix promotes HE, lower rates of anomalous HE are expected until breakdown of the passive film on the Mg17Al12 phase occurs. Furthermore, despite its

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passive-like behavior, part of the net anodic charge passing will be associated with

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dissolution of the intermetallic particle at potentials below Eb (where no HE occurs). Even though the rate of dissolution of the Mg17Al12 phase is slower than that of the α-Mg matrix,

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greater amounts of the intermetallic particle would reduce the reactive Mg area. As a

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consequence, a lower fraction of the net anodic charge will be associated with α-Mg dissolution, implying lower rates of HE. This description was hypothesized previously [45],

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but this work provides solid evidence for the first time that confirms the theory. Once the passive film on Mg17Al12 phase breaks down, it is expected that anomalous

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HE will occur on both the α-Mg matrix and the intermetallic particle. This should result in increased rates of HE above the Eb for the Mg17Al12. To confirm this hypothesis, the HE current density values shown in Fig. 15 for the AZ91 are analyzed in detail in Fig. 17. In Fig. 17, the right y-axis shows the steady state potential values measured during

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galvanostatic polarization. These values were not corrected for ohmic potential drop. For informative purposes, the Eb for Mg17Al12 determined in the potentiodynamic polarization curves is also depicted as a horizontal line at –0.950 VSSC. Note that the Eb separates regions where anomalous HE is predicted to originate from distinct regions of the surface. For this reason, the HE current density values will be analyzed as two separate sets of data.

Below that critical potential, anomalous H2 is expected to evolve only from α-Mg. Consequently, a linear fit of the HE current densities from 1 to 10 mA/cm2, indicates that the anomalous HE on the AZ91 matrix was 28 ± 2% of the net applied current density. However, above the Eb for Mg17Al12 both the α-Mg matrix and the Mg17Al12 phase are expected to contribute to the total HE determined gravimetrically. A linear fit of the HE current densities from 30 to 70 mA/cm2 indicated that anomalous HE increased to 33.17 ± 0.03% of the net

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applied current density. This represents an increase in the HE ratio of approximately 18%.

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This increase in the HE current density is in good agreement with the surface fraction of the Mg17Al12 phase in the AZ91 Mg alloy calculated in this work and previously reported by

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others [48, 52].

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This interpretation is further validated in Figs. 18 and 19, which show optical micrographs of the same region of the AZ91 Mg alloy before (referred to as Pre) and after

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(referred to as Post) galvanostatic anodic polarization at 10 mA/cm2 and 30 mA/cm2 in 0.1 M NaCl solution, respectively, to a total net charge density passed of 5 C/cm2. The

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experimental conditions were nominally identical to those in the gravimetric H2 collection measurements. It should be mentioned that specimens were etched using 2% nital before polarization to reveal the Mg17Al12 phase. These current density values were selected because they represent the situation where no corrosion attack (10 mA/cm2) and active dissolution

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(30 mA/cm2) should occur on the Mg17Al12 phase. Figs. 18 and 19 show that black regions developed on the AZ91 surface during the

application of both anodic current densities. This localized corrosion attack was restricted to the matrix and is associated with dissolution of α-Mg. However, it is observed in the red and yellow insets that during application of 10 mA/cm2 current density, no dissolution of the Mg17Al12 phases occurred (see Fig. 18), whereas evident signs of corrosion on the

intermetallic particle appeared at an applied current of 30 mA/cm2 (see Fig. 19). Consequently, there was no contribution of the Mg17Al12 phase to the total anomalous HE on the AZ91 Mg alloy below the critical potential. However, above the critical potential, the increased rates of HE were due to dissolution of both the α-Mg and the intermetallic Mg17Al12 phase. Figs. 18 and 19 provide solid experimental evidence that validates the proposed role of the Mg17Al12 phase on the anomalous HE of the beta phase-containing AZ Mg alloys.

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Finally, it should be commented that α-Mg dissolution kinetics are expected to be

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simultaneously influenced by the net anodic current density applied and the current densities associated with dissolution of the Mg17Al12 phase the anomalous HE reaction, according to:





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iMg  inet  i  iHE ,Mg  iHE ,

(1)

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where iMg and iβ are the current densities associated with dissolution of the α-Mg and Mg17Al12 phase, respectively; inet is the net anodic applied current density; and iHE,Mg and iHE,β

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are the HE current densities associated with anomalous HE on the α-Mg and Mg17Al12 phase, respectively, expressed as absolute value. In light of the results presented in this work, the

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following scenarios can be considered:

a) Below the Eb for the Mg17Al12 phase, iβ can be considered constant due to its passive behavior. Furthermore, according to Fig. 7, iHE,β = 0. Therefore, iMg will only be influenced by inet and iHE,Mg, magnitudes that increase with increased amounts of

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polarization. Please note that the situation is similar to that exhibited by pure Mg or the non-containing Mg17Al12 phase AZ Mg alloys (i.e. AZ31), where Mg dissolution kinetics will be the sum of the applied current density and the current density associated with the anomalous HE on the with the α-Mg matrix. However, in the case of the AZ91 (and consequently the rest of the Mg17Al12 phase-containing AZ Mg alloys) the term

associated with dissolution of the Mg17Al12 phase on the right side of Eqn. 1 indicates that greater amounts of this intermetallic particle due to increased Al contents in the alloy (i.e. higher iβ values) will effectively reduce α-Mg dissolution kinetics. In addition, increased concentrations of Al in the alloy are likely to promote a more protective surface film thus a decreased number of actively dissolving sites for anomalous HE on the α-Mg matrix.

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b) Above the Eb for the Mg17Al12 phase, all the terms in Eqn. 1 are expected to increase. In

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this case, both the α-Mg matrix and the Mg17Al12 phase will simultaneously dissolve and exhibit anomalous HE. However, while both phases will experience the parasitic HE

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reaction, iHE,β < iHE,Mg as shown in Fig. 15. Consequently, the lower rates of anomalous

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HE occurring on the Mg17Al12 phase than on the α-Mg matrix, will hinder Mg dissolution kinetics on the AZ91 with respect to the AZ Mg alloys that do not contain Mg17Al12 phase

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(i.e. AZ31). Furthermore, iβ values are predicted to increase significantly after breakdown of the passive film on the Mg17Al12 phase, thus reducing iMg. according to Eqn. 1. This is

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in agreement with the differences in the surface appearance of the AZ31 and AZ91 Mg alloys shown in Figs. 10 and 11.

This analysis introduces a new point of view regarding the effect of the Mg17Al12 phase on the durability of the Al-containing alloys. As discussed in the Introduction, it has

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been traditionally proposed that greater amounts Mg17Al12 phase due to increased concentrations of aluminum in the AZ alloys influence dramatically the corrosion resistance of these materials. Due to the more noble nature of the Mg17Al12 phase with respect to the αMg matrix, a greater surface fraction of the intermetallic Mg17Al12 phase results in a greater number of preferential cathodic sites under open circuit conditions, hence in increased corrosion rates. However, this is not the case under real operating conditions for these

materials. In fact, Mg alloys are almost in all cases galvanically coupled to a more noble metal, not to mention the particular cases where this is the desired use for the alloy (e.g. as sacrificial anode). The results in this paper have clarified the role of the beta-Mg17Al12 phase on the anomalous HE and anodic dissolution reactions of the AZ magnesium alloys, and have provided solid evidence that under anodic polarization greater amounts of Mg17Al12 phase

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are beneficial for AZ Mg alloys durability, decreasing the dissolution rates associated with

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the Mg matrix.

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4. Conclusions

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The role of the intermetallic Mg17Al12 second phase on the anomalous HE and anodic dissolution of AZ magnesium alloys was studied using surface characterization,

can be concluded:

Increases rates of HE were observed for the AZ31 and AZ91 Mg alloys during anodic

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

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potentiodynamic polarization and gravimetric H2 collection measurements. The following

polarization. Actively dissolving regions with concomitant HE were immediately formed during anodic potentiodynamic polarization of the AZ31 Mg alloy. A small passive region was exhibited by the AZ91. Anomalous HE originated from the anodic sites after

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breakdown of the protective film.



Passive-like behavior was exhibited by the Mg17Al12 phase with strong anomalous HE only observed at the local anodic dissolving regions following breakdown. This is the first time that enhanced rates of HE are reported for the Mg17Al12 phase during anodic polarization.



These results are in agreement with the notion that anomalous HE on anodically polarized Mg (and Mg alloys) is primarily associated with the regions undergoing active anodic dissolution.



The appearance of the previously polarized surfaces was consistent with the potentiodynamic polarization experiments. The amount of corroded regions was directly

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related with the concentration of Al in the material. Whilst the AZ31 Mg alloys exhibited a heavily corroded surface, only limited localized sites that suffered corrosion attack were



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observed for the AZ91 and Mg17Al12 phase.

The rate of anomalous HE on the AZ Mg alloys was inversely proportional to the content

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of Al in the material. The Mg17Al12 phase exhibited lower rates of HE than the AZ Mg

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alloys during anodic polarization. The charge associated with anomalous HE were determined to be around 40, 30 and 20% of the net charge passed for the AZ31, AZ91



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and Mg17Al12 phase, respectively.

Surface characterization of previously polarized AZ91 Mg surfaces at different anodic

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current densities confirmed that dissolution of the Mg17Al12 phase only occurred above the Eb for the Mg17Al12 phase. The increase in the HE current density was in good agreement with the surface fraction of the Mg17Al12 phase in the AZ91 Mg alloy. The dissolution kinetics of the Mg17Al12 phase-containing AZ Mg alloys is highly dependent on the amount of this intermetallic second phase in the material. At low

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

polarizations (below the Eb for the Mg17Al12 phase) greater amounts of this intermetallic particle due to increased Al contents in the alloy will effectively reduce α-Mg dissolution kinetics. At polarizations above the Eb for the Mg17Al12 phase, the lower rates of

anomalous HE and the increasing dissolution current densities occurring on the Mg17Al12 phase will hinder the α-Mg dissolution kinetics. 

Under anodic polarization, greater amounts of Mg17Al12 phase are beneficial for AZ Mg alloys durability, decreasing the dissolution rates associated with the Mg matrix.

Data availability

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The raw/processed data required to reproduce these findings cannot be shared at this

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time as the data also forms part of an ongoing study.

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Conflict of interest

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There are no conflicts to declare.

C. Ubeda: Investigation, Writing-Original Draft;

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G. Garces: Resources, Writing-Review & Editing; P. Adeva: Resources, Investigation, Writing-Review & Editing; I. Llorente: Investigation, Writing-Review & Editing; G.S. Frankel: Conceptualization, Writing-Review & Editing;

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S. Fajardo: Conceptualization, Methodology, Investigation, Writing-Review & Editing, Supervision, Project administration, Funding acquisition.

Declaration of interests The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

Acknowledgements S. Fajardo expresses gratitude to the State Research Agency (Ministry of Science, Technology and Universities of Spain), the Spanish National Research Council (CSIC) and the European Regional Development Fund (ERDF) for the financial support under the Project MAT2015-74420-JIN (AEI/FEDER/UE). The authors thank the CENIM XRD and Electron

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Microscopy Laboratories. Ms. Nuria Benavente is acknowledged for her help with

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metallographic preparations.

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on the Mg–Al intermetallics: Al3Mg2 and Mg17Al12, Corros. Sci. 51 (2009) 1115-1127.

AZ31

50 µm

β‒Mg17Al12

β

-p

50 µm

α

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α‒Mg

α+β

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AZ91

20 µm

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Figure 1. Secondary electrons SEM micrographs of the AZ31 and AZ91 Mg alloys.

100 µm

Figure 2. Typical optical micrograph of the AZ91 Mg alloy etched in 2% nital reagent used for Mg17Al12 surface fraction determination using image analysis.

a)

20

50

40

30

60

b) SE

80

70

90

-p

2 / degrees

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Intensity / a.u.

Mg17Al12

Mg

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re

Al

5 µm

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Figure 3. (a) XRD pattern and (b) EDS elemental map distribution of the homogenized intermetallic Mg17Al12 phase.

10

-4

10

-5

10

-6

10

-7

10

-8

of

-3

ro

10

AZ31 AZ91 Mg17Al12

-p

-2

-1.6

-1.4

re

2

Current density / A cm

10

-1.2

-1.0

-0.8

-0.6

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Potential / VSSC

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Figure 4. Potentiodynamic polarization measurements for the AZ31 and AZ91 alloys, and Mg17Al12 specimens in 0.1 M NaCl solution. Experiments were performed scanning upwards starting from ‒30 mV versus the OCP. A potential scan rate of 1 mV/s was used. Typical data from replicated experiments are presented. Current density is presented as absolute value.

a)

Film

Resin

b)

Resin

Metal

0.95

Al / Mg atomic ratio

0.90

Metal EDS Point Analysis

0.85 0.80 0.75 0.70 0.65

Al 12   0.71 Mg 17

5 µm 0.60 0

1

2

3

4

5

6

7

8

9

10

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Distance from the surface / m

Experimental 2

Experimental 0 Mg 2+ Mg Background Fitting

Mg 2p

Al 2p

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Intensity / a.u.

O  OH Background Fitting

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536 535 534 533 532 531 530 529 528 527

Binding energy / eV

55

54

53

52

51

Experimental 0 Al 3+ Al Background Fitting

Intensity / a.u.

O 1s

Intensity / a.u.

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Figure 5. (a) Back-scattered electrons SEM micrograph showing the EDS elemental point analysis performed on the Mg17Al12 phase at different distances from the surface; (b) Al/Mg ratio as a function of distance from the surface determined from the EDS elemental point analysis. Measurements performed after potentiostatic polarization at ‒1.2 VSSC in 0.1 M NaCl solution for 30 min. Mean values from replicated measurements are presented. Error bars are standard deviation.

50

49

48

Binding energy / eV

47

46

45

80 79 78 77 76 75 74 73 72 71 70 69

Binding energy / eV

Figure 6. High resolution XPS O1s, Mg2p and Al2p spectra for the AZ31 and AZ91 alloys, and Mg17Al12 specimens after potentiostatic polarization at ‒1.2 VSSC in 0.1 M NaCl solution for 30 min.

Mg17Al12 10

-3

-4

10

-5

10

-6

10

-7

10

-8

C

B

of

10

A -1.4

-1.2

-1.0

-0.8

-0.6

-p

Potential / VSSC

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Current density / A cm

2

D

B

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re

A

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1 mm

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C

D

Streams of H2

Figure 7. Potentiodynamic polarization curve of the Mg17Al12 phase in 0.1 M NaCl solution (above). Surface appearance of the Mg17Al12 phase during the same polarization experiment at different applied potentials, labelled from A to D (below).

10

-2

10

-3

10

-4

C

10

-5

10

-6

A 10

AZ31

-7

-1.6

-1.4

-1.2

-1.0

Potential / VSSC

B

Stream of H2

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-p

A

-0.8

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B

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Current density / A cm

2

E D

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C

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1 mm

E

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D

Corrosion front Streams of H2

Figure 8. Potentiodynamic polarization curve of the AZ31 Mg alloy in 0.1 M NaCl solution (above). Surface appearance of the AZ31 during the same polarization experiment at different applied potentials, labelled from A to E (below).

10

-2

10

-3

10

-4

10

-5

10

-6

10

-7

10

-8

10

-9

E D

C

B

A

-1.6

-1.4

AZ91 -1.2

-1.0

Potential / VSSC

A

re

-p

B

-0.8

of

-1

ro

2

Current density / A cm

10

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C

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1 mm

E

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D

Corrosion front Streams of H2

Figure 9. Potentiodynamic polarization curve of the AZ91 Mg alloy in 0.1 M NaCl solution (above). Surface appearance of the AZ91 during the same polarization experiment at different applied potentials, labelled from A to E (below).

AZ31

AZ91 2 mm

2 mm

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2 mm

Mg17Al12

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re

-p

ro

Figure 10. Surface images of the AZ31 and AZ91 Mg alloys, and the Mg17Al12 phase after the potentiodynamic polarization measurements in 0.1 M NaCl solution (Figs. 7-9).

AZ31

AZ91

50 µm

50 µm

-p

ro

of

Mg17Al12

3 µm

re

50 µm

Jo

ur na

lP

Figure 11. SEM micrographs of the corroded regions on the AZ31 and AZ91 Mg alloys, and the Mg17Al12 phase, shown in Fig. 10.

0.25 3 mA cm

2

1 mA cm

0.20

2

0.15

of

0.10 0.05

ro

H2 volume collected / mL cm

2

0.30

0.00 1000

2000

3000

4000

5000

-p

0

0.30

2

50 mA cm

0.25

2

7 mA cm

2

10 mA cm

lP

30 mA cm

0.20

Jo

2

2

5 mA cm

ur na

H2 volume collected / mL cm

2

70 mA cm

re

Time / s

2

0.15

3 mA cm

0.10

2

0.05

1 mA cm

2

0.00 0

200

400

600

800

1000

Time / s

Figure 12. Volume of H2 determined from gravimetric measurements as a function of time for the AZ31 Mg alloy in 0.1 M NaCl solution under the application of different anodic 40

current densities. A net charge density of 5 C cm‒2 was passed. Mean values from replicated experiments are presented. Error bars are standard deviation.

0.20

3 mA cm

2

1 mA cm

2

of

0.15

ro

0.10

0.05

0.00 1000

2000

3000

re

0

-p

H2 volume collected / mL cm

2

0.25

4000

5000

Time / s

lP

0.25

0.20

2

50 mA cm

10 mA cm

Jo

2

2

2

30 mA cm

2

5 mA cm

ur na

H2 volume collected / mL cm

2

70 mA cm

7 mA cm

2

0.15

0.10

3 mA cm

2

0.05

1 mA cm

0.00 0

200

400

600

Time / s

41

800

1000

2

Jo

ur na

lP

re

-p

ro

of

Figure 13. Volume of H2 determined from gravimetric measurements as a function of time for the AZ91 Mg alloy in 0.1 M NaCl solution under the application of different anodic current densities. A net charge density of 5 C cm‒2 was passed. Mean values from replicated experiments are presented. Error bars are standard deviation.

42

0.12 0.10

3 mA cm

2

1 mA cm

2

0.08

of

0.06 0.04 0.02

ro

H2 volume collected / mL cm

2

0.14

0

1000

2000

-p

0.00 3000

4000

5000

Time / s 70 mA cm

0.10

50 mA cm

2

lP

0.12

2

30 mA cm

10 mA cm

2

7 mA cm

Jo

2

2

ur na

H2 volume collected / mL cm

2

0.14

re

0.16

0.08

5 mA cm

2

0.06 0.04

3 mA cm

2

0.02

1 mA cm

2

0.00 0

200

400

600

800

1000

Time / s

Figure 14. Volume of H2 determined from gravimetric measurements as a function of time for the Mg17Al12 phase in 0.1 M NaCl solution under the application of different anodic 43

30 AZ31

25

of

20 AZ91

ro

15 10

-p

Mg17Al12

0 0

10

20

re

5

30

40

lP

H2 evolution current density / mA cm

2

current densities. A net charge density of 5 C cm‒2 was passed. Mean values from replicated experiments are presented. Error bars are standard deviation.

50

60

70

ur na

Applied current density / mA cm

80

2

Jo

Figure 15. Current density values associated with HE for the AZ31 and AZ91 Mg alloys, and the Mg17Al12 phase as a function of the applied anodic current density in 0.1 M NaCl solution. Current densities were calculated from the HE rates determined using the gravimetric method (Figs. 12-14). Current density is presented as absolute value. Mean values from replicated experiments are presented. Error bars are standard deviation.

44

1.0 AZ31 AZ91 Mg17Al12

of ro

0.6

2

QH / Qnet

0.8

-p

0.4

0.0 1

lP

re

0.2

10

ur na

Applied current density / mA cm

100 2

Jo

Figure 16. Charge associated with HE normalized to the net charge passed for the AZ31 and AZ91 Mg alloys, and the Mg17Al12 phase as a function of the applied anodic current density in 0.1 M NaCl solution. The net charge density passed was 5 C cm‒2. Mean values from replicated experiments are presented. Error bars are standard deviation.

45

1.5

2

1.0

of

1

0.5

0 2

4

6

8

0.0

ro

0

10

10

0.0003

-p

Slope = 0.3317 2 R=1

Eb -phase

5 Slope = 0.28 2 R = 0.9886

0

10

20

30

-0.5 -1.0 -1.5

0.02

40

ur na

0

re

15

Potential measured / VSSC

20

3

lP

H2 evolution current density / mA cm

2

25

-2.0 50

60

70

Applied current density / mA cm

80

2

Jo

Figure 17. Current density values associated with HE for the AZ91 Mg alloy as a function of the applied anodic current density in 0.1 M NaCl solution from Fig. 13. Right hand axis shows the potentials measured during galvanostatic polarization. Potential values are not corrected for ohmic potential drop. Current density is presented as absolute value. Mean values from replicated experiments are presented. Error bars are standard deviation.

46

Post

50 µm

-p

50 µm

Post

Pre

Post

20 µm

lP

re

Pre

ro

of

Pre

20 µm

Jo

10 mA/cm2

ur na

Figure 18. Optical micrograph of the AZ91 Mg alloy before (pre) and after (post) galvanostatic anodic polarization at 10 mA/cm2 to a total net charge density passed of 5 C/cm2 in 0.1 M NaCl solution. The surface was etched in 2% nital reagent prior to polarization to reveal the Mg17Al12 phase. The insets show two selected regions at a higher magnification.

47

Post

50 µm

Post

Pre

Post

lP

re

Pre

-p

50 µm

ro

of

Pre

20 µm

20 µm

Jo

ur na

Figure 19. Optical micrograph of the AZ91 Mg alloy before (pre) and after (post) galvanostatic anodic polarization at 30 mA/cm2 to a total net charge density passed of 5 C/cm2 in 0.1 M NaCl solution. The surface was etched in 2% nital reagent prior to polarization to reveal the Mg17Al12 phase. The insets show two selected regions at a higher magnification. White arrows locate the attacked regions on the Mg17Al12 phase after anodic polarization.

48

Table 1. Chemical composition (wt. %) of the Mg alloys and the Mg17Al12 phase used. Material

Mg

Al

Zn

Mn

Si

Cu

Fe

Ni

Ca

Zr

Others

AZ31

Bal.

3.1

0.73

0.25

0.02

<0.001

0.005

<0.001

0.0014

<0.001

<0.30

AZ91

Bal.

9.4

0.82

0.20

<0.05 <0.025 <0.004 <0.001 <0.001 <0.001

<0.30

0.003 0.032

‒

0.004

0.009

‒

‒

‒

Jo

ur na

lP

re

-p

ro

of

Mg17Al12 56.07 43.93

49

‒

Table 2. EDX element analysis composition (at. %) on the corroded areas of the Mg alloys and the Mg17Al12 phase used. Mean values and standard deviations from different measurements are reported. Element (at.%) Material Mg

Al

Cl

AZ31

59 ± 4

35 ± 4

2.4 ± 0.3

2.1 ± 0.4

AZ91

65 ± 5

27 ± 2

0.09 ± 0.07

8±3

Mg17Al12

10 ± 3

50 ± 2

40 ± 2

of

O

Jo

ur na

lP

re

-p

ro

‒

50