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Study of the corrosion mechanism of Mg–Gd based soluble magnesium alloys with different initial texture states

  • Yanchun Zhu , Niulin Wang , Zhibing Chu ORCID logo EMAIL logo , Yong Niu , Liang Ma and Ling Qin
Published/Copyright: August 14, 2023
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Corrosion Reviews
From the journal Corrosion Reviews Volume 41 Issue 6

Abstract

In order to clarify the influence of different initial texture states on the corrosion mechanism of soluble Mg alloy materials, the as cast and after extruded + perforated (EP) deformed Mg–Gd based soluble magnesium alloys are investigated by the microstructure, surface morphology, surface volta potential, immersion test and electrochemical measurement separately. The results indicate that: the rate of corrosion of the as cast Mg–Gd based soluble magnesium alloy can reach 43.85 mg/cm2/h at 93 °C in a 3 wt% KCl solution, while after EP deformation the rate of corrosion is greatly reduced to only 8.37 mg/cm2/h. Combined with the microstructure analysis, it is concluded that the EP deformed destroyed the coarse reticulated second phase in the as cast structure, which reduced the micro-electrocouple corrosion effect of the second phase. Finally, the corrosion mechanism models for different initial texture states are established through the analysis of microstructure and corrosion morphology, respectively. It is found that the microscopic corrosion mechanism of the as cast Mg–Gd based soluble magnesium alloy is mainly intercrystalline corrosion, which is a superposition of micro-electrocouple corrosion and Mg matrix dissolution. While the microscopic corrosion mechanism of the EP deformed is mainly intracrystalline corrosion, which is manifested as pitting corrosion.

Keywords: corrosion mechanism; intercrystalline corrosion; intracrystalline corrosion; Mg–Gd based soluble Mg alloy; surface morphology

1 Introduction

As the lightest class of metals among structural materials, magnesium (Mg) alloys have the advantages of high specific strength, good thermal conductivity, and electromagnetic shielding, and have the prospect of wide application for transportation, communication, and aerospace (Anjum et al. 2022; Peng et al. 2018; Pawar et al. 2017; Sudholz et al. 2009). However, due to its low corrosion resistance, the development and application as a degradable material in the field of oil and gas extraction is gradually emerging (Zhong et al. 2021). Specifically, during the hydraulic fracturing process, sealing tools such as fracturing balls or bridge plugs are used to temporarily seal the oil and gas well. In the past, when the temporary sealing tools were used, they had to be removed, which was very time consuming and environmentally unfriendly. If soluble magnesium alloy is used as a sealing tool, after oil and gas extraction have been completed and used, does not need to be taken out from the ground. Instead, it is dissolved and disabled by direct passage of the corrosion fluid. By dissolving the degradable material, the oil and gas extraction process becomes economical and efficient. This is of considerable importance for the exploitation of unconventional oil and gas fields with low porosity and low permeability (Lian et al. 2015). As a result, attempts are being made to use degradable materials, particularly magnesium alloys, to fabricate sealing tools.

However, as a degradable material, it is expected to show rapid corrosion rates for self-degradation purposes after oil extraction completed. One effective way to increase the rate of degradation of magnesium alloys is through alloying (Liu 2021; Srinivasan et al. 2014; Wang et al. 2022). For example, Cu and Ni, as impurity elements that reduce the corrosion resistance of magnesium alloys, are often added to magnesium alloys to increase the corrosion rate (Oh et al. 2016). Wang et al. (2022) investigated the effect of different Ni contents (0.1, 0.2 and 0.5 wt%) on the corrosion rate of Mg–2Gd alloy, a high plasticity soluble magnesium alloy used for sealing tools in the oilfield. The addition of Ni was found to change the rare earth structure of the alloy to a basic fibrous structure, which leads to an increase in the corrosion rate of the alloy. Niu et al. (2019) prepared a new Mg–4Zn–Ni alloy in order to obtain a fracturing material that could be used for oil extraction. The results showed that the formation of the Mg2Ni phase not only increased the compressive strength of the alloy, but also accelerated the degradation rate of the Mg–Zn series of alloys. Moreover, Zhang et al. added 4 wt% Cu to high strength Mg–3Zn–1Y and found a surprising 38-fold increase in corrosion rate (2017b).

In addition, many scholars have found that the microstructure of magnesium alloys also affects their corrosion behavior. Such as crystal structure, weave orientation, distribution of secondary phases, number of secondary phases, size of grains in the matrix phase, composition and stability of surface films (Lu et al. 2015; Liao et al. 2012b; Sun et al. 2022; Wang et al. 2019). The increased corrosion rate of magnesium alloys containing germanium (Ge) was attributed to more severe micro-electro-coupling corrosion due to an increase in the volume fraction of the second phase (Chen et al. 2019). Similarly, Zhang et al. investigated that the volume fraction of the second phase could be adjusted by controlling the content of Gd, which in turn affected the corrosion rate of soluble magnesium alloys (2017a). More recently, Li et al. (2022) researched the effect of deformation twinning on the corrosion behavior of magnesium alloys. It was found that twinning has a dual effect on corrosion behavior. Among other things, the thickness and nature of the film depends on the crystallographic texture. In most cases, the films formed on the basal planes were thinner compared to the prismatic and pyramidal planes.

In recent years, the long period stacking ordered (LPSO) phase has been receiving a lot of attention in the field of soluble magnesium alloys (Guo et al. 2020; Yamasaki et al. 2011). In terms of corrosion performance, firstly, the LPSO phase has a different electrochemical potential to the Mg matrix and may cause micro-electro-coupling corrosion. Secondly, the morphology and distribution of the LPSO phase also have a significant effect on the corrosion rate (Yamasaki et al. 2011). For example, Liu et al. (2019) investigated the corrosion behavior of LPSO phase in Mg-Gd-Zn-Zr alloys by measuring the potential difference between the LPSO phase and the magnesium matrix, and showed that the low relative potential and volume fraction of the LPSO phase reduced the acceleration of micro-electro-coupling corrosion. Furthermore, Ma et al. (2022) introduced the LPSO phase into the Mg-Gd-Ni alloy matrix and grain boundaries by adjusting the heat treatment process, thereby improving the degradation properties. Based on current research, it has been found that methods such as alloying, crystal orientation, volume fraction of the second phase, potential difference between the matrix phase and the second phase and the distribution of the second phase all have an important influence on the degradation properties of magnesium alloys. However, to our best knowledge, few studies have been reported that focus on how different initial texture states affect the corrosion rate and mechanism of soluble magnesium alloys after EP plastic deformation processes.

Therefore, this paper takes the as cast and after extruded + perforated (EP) deformed Mg–Gd based soluble magnesium alloy as the object of study. On the one hand, the influence of different initial texture states on the corrosion rate was analysed, on the other hand, the corrosion mechanism was clarified to provide theoretical support for the application of soluble magnesium alloys in energy extraction. The study of their corrosion behavior and corrosion mechanisms is intended to provide guidance on the regulation of corrosion rates of soluble temporary plugging tools used in the oil extraction process, such as fracturing balls or bridge plugs.

2 Materials and methods

2.1 Materials

The research material was provided by Chongqing Yuhua New Materials Ltd. Figure 1 shows a flow chart of the production of Mg–Gd based tubes and a physical view of the tube. Firstly, small lumps of pure Mg (99.98 wt%), pure Gd (wt.%), pure Cu powder (99.5 wt%), pure Fe powder (99.5 wt%), and a small amount of Si (99.5 wt%) were added in batches to the preheated crucible, and then heated to melt the metal in the crucible and the molten material was refined once and twice, respectively. The electric melting furnace was then switched off, left to stand and the rested material is cast. After completion the material was loaded into the heat treatment furnace for homogenisation, ramped up to 420 °C and held for 24 h. And then, the homogenised material was extruded into rods for the deformation of the thick-walled dissolvable magnesium alloy tubes. Finally, the use of molten cast soluble magnesium alloy rods were first pretreated, then deformed and hollowed using machining, and finally skinned to obtain a deformed thick-walled soluble magnesium alloy seamless tube. The tube has an outer diameter of 120 mm, an inner diameter of 40 mm and a wall thickness of 40 mm. The final finished physical drawing was shown in Figure 1b; the alloy test results were shown in Table 1.

Figure 1: Mg–Gd based EP deformed tubes: (a) manufacturing process flow diagram; (b) physical diagram.
Figure 1:

Mg–Gd based EP deformed tubes: (a) manufacturing process flow diagram; (b) physical diagram.

Table 1:

Chemical composition Mg–Gd based alloys.

Element Gd Ni Cu Fe Si Mg
Content (wt.%) 2.4 0.1 0.02 0.01 0.01 Balance

2.2 Microstructural characterisation

The morphology and microstructure of the specimens were observed under a scanning electron microscope (SEM, ZEISS OXFORD) and an electron backscatter diffractometer (EBSD, ZEISS OXFORD). The composition of the precipitated second phase was determined by energy dispersive spectroscopy (EDS, ZEISS OXFORD). The Mg–Gd based cast and EP deformed alloys were identified using X-ray diffraction (XRD) at a scanning rate of 4°/min and scanning angles of 10° to 90°. The test specimens with the size of 10 mm × 10 mm × 10 mm were polished to 1200 grit with SiC sandpapers. Among other things, shooting SEM and EDS required electrolytic polishing and etching of the specimen surface with an etching solution (2.1 g C6H3N3O3 + 5 mL CH3COOH + 5 mL H2O+ 35 mL C2H5OH) to facilitate the observation of grain boundaries. The electrolyte was made with a HClO4 to C2H5OH ratio: 1:9, with an electrolysis time of 120 s, a voltage of 20 V and a current of 0.3 A.

2.3 Scanning Kelvin probe force microscopy (SKPFM)

The SKPFM (Bruker Dimension ICON) tests specimens with the size of 10 mm × 10 mm × 2 mm were polished to 5000# SiC sandpapers. Then the surface was mechanically polished with 1 um diamond polishing paste and ultrasonic cleaned with anhydrous ethanol. All tests were carried out at ambient temperature and contact mode was selected, using a PTIR-plated silicon tip with an intrinsic frequency of 75 kHz and a force constant of 2.5 N/m. After the tests were completed, the data was analyzed using the SKPFM image post-processing software: Nanoscope Analysis.

2.4 Immersion test

The immersion test specimen was cylindrical and the size of diameter 40 mm, height 10 mm. The specimens were successively ground to 5000# SiC sandpapers. The test was carried out using a digital thermostatic water bath at a controlled temperature of 93 °C and 25 °C. And 3 wt% KCl was selected as the corrosive solution to soak the specimens for 12 h. After each removal the specimen was immersed in chromic acid (200 g/L CrO3 + 10 g/L AgNO3) solution for 5 min–10 min for washing and then the experimental data related to height (h), diameter (d) and weight (m) were recorded. The dissolution rate was calculated by using the following Equation (1) (Huang et al. 2023).

(1)P=(m1−m2)/S*T

where P is the dissolution rate, (unit: mg/cm2/h), m1 is the pre-test mass, m2 is the post-test mass (unit: mg), S is the mean value of the surface area of the specimen (unit: cm2), T is the time between each measurement of the specimen (unit: h).

2.5 Electrochemical measurement

The size of the electrochemical test specimens was 10 mm × 10 mm × 10 mm. The working electrode (WE) was encased in epoxy resin with only one surface exposed and the back of the specimen is connected to the red terminal of the electrochemical workstation via an Al wire covered with insulation. Furthermore, three parallel specimens were prepared for each specimen to minimise chance errors associated with the experiment. Use platinum (Pt) sheet as counter electrode (CE), size: 10 mm × 10 mm × 0.5 mm, reference electrode: use saturated calomel electrode (SCE). The first test was an open circuit potential (OCP) and the test time was chosen to be 1800 s to obtain a stable open circuit potential. The electrochemical impedance spectroscopy (EIS) test was selected for a frequency range of 100 kHz to 10 mHz, with a maximum amplitude of 5 mV and a scan rate of 1 mV/s. The results were then analysed using ZSimpWin software. The potentiodynamic polarisation (PDP) curves was set at potentials from −0.3 to 0.5 V and a scan rate of 0.01 V/s. The results were analysed using CView software to obtain the corrosion potential (Ecorr) and corrosion current density (Icorr).

3 Results

3.1 Microstructure and surface potential analysis

As shown in Figure 2a, the second phase of the as-cast alloy is coarse, reticulated and distributed along the grain boundaries in the matrix. As shown in Figure 2b, the TD directional structure of the alloy after deformation is equiaxed and the grain boundaries have been refined, which is attributed to the continuous dynamic recrystallisation mechanism (CDRX). As can be seen from the ED directional diagram in Figure 2c, the second phase of the as-cast alloy, which was originally in the form of a coarse mesh, has been broken up by the deformation process, and the broken second phase precipitates in the axial direction. In addition, the shape can be clearly seen as slaty or granular, with a few square precipitated phases. In addition, in order to determine the alloy phase composition, EDS and XRD tests were carried out, respectively. The EDS mapping results in Figure 2e shows that the slate-like bright second phase precipitated in the Mg–Gd based alloy is mainly composed of Mg, Gd, and Ni. Based on the EDS results, it is assumed that the slate-like precipitated phases containing Mg, Gd, and Ni are LPSO phases.

Figure 2: As cast and EP deformed microstructures of Mg–Gd based alloys: (a) as cast alloy texture images; (b–f) EP deformed alloy texture images and corresponding EDS mapping at high magnification.
Figure 2:

As cast and EP deformed microstructures of Mg–Gd based alloys: (a) as cast alloy texture images; (b–f) EP deformed alloy texture images and corresponding EDS mapping at high magnification.

In addition, the cubic precipitated phase in Figure 2d was photographed enlarged and EDS mapped, and the results were shown in Figure 2f. The rare earth element Gd atoms were enriched in the cubic phase and the content of Mg atoms is low, with the cubic phase having a length and width of about 2 μm. The results for the possible phases, obtained by counting the content of the elements in the second phase, were shown in Table 2. As shown in Table 2, the second phase distributed along the grain boundaries of the as cast alloy is Mg5Gd and the particulate phase precipitated within the grain boundaries is the Mg2Gd phase. For EP deformed alloys, the cubic phase has a Gd content of nearly 36.6 at.%, which was a Gd-rich phase. Previous studies have also reported that these cubic phases are precipitates of the Mg–Gd alloy system, possibly the Mg2Gd phase (Zhang et al. 2018; Zhao et al. 2019; Zhang et al. 2018). The mapping results in Table 2-D showed that the precipitated slats contain 83 at.% Mg and 17 at.% Gd in the second phase. Thus, the precipitated phase is probably the Mg5Gd phase. In addition, the mapping results in Table 2-C showed that the bulk precipitated phase contains not only Mg, Gd elements but also some amount of Ni. Therefore, it was speculated that Mg-Gd-Ni ternary LPSO phase was formed in the alloy through EP deformation. To further confirm the EDS test results, XRD analysis of the specimens was carried out and the results were shown in Figure 3. From the test results, it was known that the EP deformed alloy consists mainly of α-Mg, Mg2Gd, Mg5Gd, and Mg-Gd-Ni phases.

Table 2:

EDS scan results.

Specimen Mg (at.%) Gd (at.%) Ni (at.%) Possible phase
A 79.0 21.0 – Mg5Gd
B 63.4 36.6 – Mg2Gd
C 64.9 25.8 9.3 Mg-Gd-Ni
D 83.0 17.0 – Mg5Gd
E 98.7 1.3 – α-Mg
F 64.6 35.4 – Mg2Gd
Figure 3: Mg–Gd based soluble magnesium alloy: XRD results in as cast and EP deformed.
Figure 3:

Mg–Gd based soluble magnesium alloy: XRD results in as cast and EP deformed.

Figure 4 shows the surface morphology and surface local volta potential distribution between the second phase Mg5Gd and the α-Mg matrix phase at the grain boundaries in the as cast Mg–Gd based alloy. As shown in Figure 4a, the distribution of the reticulated Mg5Gd phase along the grain boundaries is slightly lower in height relative to the α-Mg matrix. Moreover, the precipitated material around the second phase has a needle-like distribution and can reach a height of about 400 nm, as shown in Figure 4b. As shown in Figure 4c, the alloy surface is darker and the work of escape is lower at the distribution of the second phase Mg5Gd. This indicates that the second phase Mg5Gd distributed at the grain boundaries has a lower potential value. It also shows that in the early stages of corrosion, the alloy surface first forms micro-electro-couple corrosion at the grain boundaries. In this case, the second phase Mg5Gd distributed along the grain boundaries acts as a micro-anode for the galvanic coupling corrosion, while the Mg matrix acts as the cathodic phase. In addition, in order to derive the magnitude of the corrosion rate between the original castings at the start of the reaction, the rate of the corrosion reaction can be judged from the magnitude of the potential difference between the surface of the matrix phase and the surface of the second phase. It is generally believed that the larger the value of the potential difference between the matrix phase and the second phase in the alloy, the greater the driving force for corrosion and the more likely it is that corrosion dissolution will occur (Liu et al. 2019). As shown in Figure 4d, the potential differences between the reticulated phase Mg5Gd and the α-Mg matrix were 110 mV. This indicated that the addition of the rare earth element Gd forms a rare earth-rich second phase which facilitates the reduction of the potential value of the second phase.

Figure 4: SKPFM test results of surface morphology and surface potential difference between the second phase of the as cast Mg–Gd based alloy and the α-Mg matrix: (a) 2D surface morphology; (b) 3D surface morphology; (c, d) surface potential distribution of Mg5Gd and measurement results.
Figure 4:

SKPFM test results of surface morphology and surface potential difference between the second phase of the as cast Mg–Gd based alloy and the α-Mg matrix: (a) 2D surface morphology; (b) 3D surface morphology; (c, d) surface potential distribution of Mg5Gd and measurement results.

Figure 5 shows the surface morphology and surface local potential distribution between the intracrystalline precipitated phase and the α-Mg matrix phase in the EP deformed Mg–Gd based alloy. As shown in Figure 5a, the second phase precipitated on the surface of the EP deformed alloy showed a certain directionality. Furthermore, the surface of the EP deformed alloy has greater roughness values compared to the surface morphology of the as-cast alloy, as shown in Figure 5b. The surface potential analysis of the slatted Mg-Gd-Ni phase in Figure 5d shows that the second phase Mg-Gd-Ni also has a lower potential value than the α-Mg matrix, reaching a potential value of 100 mV with the α-Mg matrix. Again, micro-electro-coupling corrosion was formed with the matrix during the corrosion process, thus accelerating dissolution.

Figure 5: SKPFM test results of surface morphology and surface potential difference between the second phase of the EP deformed Mg–Gd based alloy and the α-Mg matrix: (a) 2D surface morphology; (b) 3D surface morphology; (c) surface potential distribution; (d) surface potential measurement results.
Figure 5:

SKPFM test results of surface morphology and surface potential difference between the second phase of the EP deformed Mg–Gd based alloy and the α-Mg matrix: (a) 2D surface morphology; (b) 3D surface morphology; (c) surface potential distribution; (d) surface potential measurement results.

3.2 Corrosion morphology analysis and corrosion rate determination

Figure 6 shows the macroscopic morphology of the as cast and EP deformed soluble Mg alloys after corrosion. As can be seen from the figure, the as cast alloy shows a more pronounced reduction in thickness in the same time. Compared to the as cast specimen, huge pitting holes appear on the surface after EP deformation. Analysis of the corrosion morphology revealed that the two differed in the main types of corrosion that occurred. The as cast alloy showed mainly uniform corrosion, while the EP deformed specimen showed mainly pitting corrosion. Figure 6c and d shows the macroscopic morphology of the as cast and EP deformed alloys after 1 h of corrosion at 25 °C and a potassium chloride concentration of 3 wt%. As shown in Figure 6c, when the as cast specimen was first placed in the potassium chloride solution, a large number of bubbles were immediately generated on the surface, and these bubbles were very large and dense. After 1 h of reaction, the specimen is removed and traces of filiform corrosion can be observed on the surface. As shown in Figure 6d, the EP deformed specimen immediately produced a black product film on the surface of the specimen at the beginning of dissolution and a reticulated insoluble material appeared on the surface of the specimen. From the analysis of the corrosion morphology at high and low temperatures, we know that the temperature will have an effect on the formation process of the corrosion product film between the two, thus causing the difference in corrosion behavior between the two.

Figure 6: Macroscopic morphology of Mg–Gd based alloy after 12 h immersion test in 3 wt% KCl solution: (a) 93 °C as cast; (b) 93 °C EP deformed; (c) 25 °C as cast; (d) 25 °C EP deformed.
Figure 6:

Macroscopic morphology of Mg–Gd based alloy after 12 h immersion test in 3 wt% KCl solution: (a) 93 °C as cast; (b) 93 °C EP deformed; (c) 25 °C as cast; (d) 25 °C EP deformed.

In order to quantitatively compare the dissolution rate of as cast and EP deformed alloys, the dissolution rate of the two within 12 h was calculated by using formula (1). Among them, as shown in Figure 7a, the EP deformed and as cast alloy in the first 2 h corrosion rate difference was not very big. It shows that the corrosion of the alloy requires a process of adaptation (erosion of Cl− and generation of product film). In addition, the reaction rate of the as cast alloy increased considerably after 2 h, indicating that the corrosion product film generated in the as cast state is less likely to be deposited on the surface of the magnesium alloy to hinder the corrosion process. As shown in Figure 7b, within the first 2 h of corrosion, the weight gradient change was 0.49 g for the EP deformed alloy and 0.63 g for the as cast alloy. There was no great difference in weight loss between the two. However, when the corrosion time exceeded 2 h, the weight loss of the as cast soluble magnesium alloy increased sharply. When it reaches 12 h, the weight gradient change of the EP deformed alloy is only 0.696 g, while the change of the as cast alloy reaches 3.718 g. Through the above analysis, it shows that in the late stage of corrosion and dissolution reaction, the weight loss rate of as cast magnesium alloy is much larger than that of EP deformed alloy.

Figure 7: Mg–Gd based as cast and EP deformed alloys: (a) dissolution rate curve; (b) graph of weight gradient variation.
Figure 7:

Mg–Gd based as cast and EP deformed alloys: (a) dissolution rate curve; (b) graph of weight gradient variation.

Figure 8 shows the macroscopic morphology and three-dimensional morphology of the surface of the specimen taken at super depth of field after 12 h of corrosion dissolution. As can be seen in Figure 8a and b, pitting occurs in both the as cast and EP deformed alloys. As shown in Figure 8c and d, the pitting pit depth for the as cast alloy is 529.9 μm, whereas the pitting pit depth for the EP deformed alloy can reach 1247.6 μm. Therefore, but in terms of the severity of pitting, the EP deformed alloy pitting is more serious and mainly localised corrosion occurs.

Figure 8: Surface morphology and 3D rendering of specimens after 12 h immersion of Mg–Gd based as cast and EP deformed alloys: (a, c) as cast alloys; (b, d) EP deformed alloys.
Figure 8:

Surface morphology and 3D rendering of specimens after 12 h immersion of Mg–Gd based as cast and EP deformed alloys: (a, c) as cast alloys; (b, d) EP deformed alloys.

3.3 Electrochemical parameter analysis

Figure 9 shows the PDP curves and Bode plots of as cast and EP deformed alloys and Nyquist curves as well as the equivalent circuit diagram of EIS test. The corresponding fitting results were shown in Table 3. As shown in Figure 9a, when the corrosion potential reaches point a, the EP deformed soluble magnesium alloy was the first to reach the self-corrosion potential (Ecorr) and dissolution begins to occur. While the as cast soluble magnesium alloy was still in the cathodic state, no corrosion occurred. Then, when the corrosion potential is scanned to point b, the anode of the EP deformed soluble magnesium alloy has started to corrode, but the as cast alloy has only reached the self-corrosion potential, corrosion has not yet begun. Until, when the corrosion potential reaches point c, both the EP deformed and the as cast soluble magnesium alloy were under the anodic potential. If analysed in terms of Ecorr alone, the corrosion drive of the EP deformed alloy is greater than that of the as cast alloy. However, the self-corrosion current density (Icorr) is more representative of the magnitude of the corrosion rate than the Ecorr analysis. Since magnesium and magnesium alloys often show anomalous hydrogen precipitation behavior during dissolution [anodic precipitation (Liao et al. 2012a)], relevant parameters such as Icorr were calculated by fitting the cathodic Tafel extrapolation method (Shi et al. 2010). It is generally believed that the higher the Icorr, the more likely to be corroded (Filotas et al. 2021; Feng et al. 2022; Ma et al. 2021). According to the fitting results, the Icorr of EP deformed state was 5.6umA/cm2, while the Icorr of as cast magnesium alloy was 131 umA/cm2. Due to the Icorr of the EP deformed soluble magnesium alloy was much smaller than that of the as cast soluble magnesium alloy, indicating that the as cast soluble magnesium alloy has more superior degradation properties.

Figure 9: Mg–Gd based as cast and EP deformed state: (a) polarization curve; (b) Bode plot; (c) Nyquist plot; (d) equivalent circuit diagram.
Figure 9:

Mg–Gd based as cast and EP deformed state: (a) polarization curve; (b) Bode plot; (c) Nyquist plot; (d) equivalent circuit diagram.

Table 3:

Electrochemical parameter fitting of equivalent circuit.

Specimens R s (Ω/cm−2) R ct (Ω/cm−2) Y dl (Ωcm−2sn) n dl
As cast 7.527 ± 0.08 38.16 (0.07 ± 0.03) × 10−5 0.876 ± 0.003
EP deformed 7.743 ± 0.08 748.7 (1.66 ± 0.03) × 10−5 0.908 ± 0.003

It was also evident from the Bode diagram of the alloy in Figure 9b that the impedance modulus of the EP deformed soluble magnesium alloy was much larger than the as cast impedance modulus in the low frequency band. As can be seen from the Nyquist diagram in Figure 9c, both as cast and EP deformed soluble magnesium alloys have a capacitive circuit in the high frequency region. Among them, the high frequency capacitive circuit is related to the double electric layer consisting of alloy and solution, which is controlled by the process of charge transfer. In general, the size of the radius of the capacitive circuit reflects the value of the charge transfer resistance (Rct) summarised by the electrochemical reflection process. And the larger the radius, the greater the Rct value, and the stronger the corrosion resistance (Jiang et al. 2022; Wang et al. 2020). The radius of the capacitive circuit of EP deformed soluble magnesium alloys was about 800, while the radius of the capacitive circuit of as cast magnesium alloys was about 50, indicating that EP deformed soluble magnesium alloys were more resistant to corrosion.

Further analysis fitted the EIS test results into an equivalent circuit as shown in Figure 9d. The electrochemical corrosion reaction of the soluble magnesium alloy is abstracted as a circuit model consisting of three electrochemical components. Where Rs is the resistance between the specimen and the degradation solution (which includes the resistance of the solution and the internal resistance of the specimen itself). Cd refers to electric double layer capacitance, which is derived from the capacitive reactance between the inactive ions in the electrolyte and the surface of the corrosion specimens, without electrochemical reaction, and only changes the distribution state of the charge [with parameters Y and n denoting the non-ideal capacitance parameters, when n = 0, the non-ideal capacitance is equated to a purely resistive element; when n = 1, the non-ideal capacitance is equated to a purely capacitive element (Huang et al. 2023)]. Rct represents the Faraday impedance, which originates from the resistance between the active ions in the electrolyte and the corroded specimen, both in terms of the chemical reaction occurring and the change in the state of charge distribution. From the fitted results Table 3 it is known that the Rs values were essentially the same for the different alloys. The Rct is significantly increased after EP deformation, making the alloy more resistant to corrosion.

4 Discussion

4.1 Study of the corrosion mechanism of cast Mg–Gd based alloys

From the previous analysis of the texture of the Mg–Gd based alloy, it is known that the second phase of the as cast alloy is coarse reticulated, distributed along the grain boundaries, with the main composition being Mg5Gd type intermetallic compounds. And from the SKPFM test analysis, we know that the electrode potential of the second phase of the Mg–Gd based magnesium alloy is more negative than that of the magnesium matrix. So in the process of corrosion and dissolution, the second phase will act as the micro-anode of micro-electro-couple corrosion, while the magnesium matrix acts as the micro-cathode of micro-electro-couple corrosion. In addition, according to the texture morphology analysis of the as cast alloy, a micro-electro-couple corrosion with a large cathode and a small anode can be formed, thus accelerating the corrosion rate of the grain boundaries.

As shown in Figure 10a, micro-electro-couple corrosion is the first to form between the second phase at the grain boundaries and the magnesium matrix during the initial stages of immersion. The corrosion starts to spread gradually into the grain along the grain boundaries (Liao et al. 2012a). In addition, due to the negative differential effect of the magnesium alloy, dissolution of magnesium occurs at the surface of the magnesium matrix, generating Mg(OH)2 onto the surface of the alloy. As shown in Figure 10b, in the middle of the immersion, as the second phase Mg5Gd at the grain boundaries corrodes electrically with its surrounding magnesium matrix, the second phase in the form of a coarse mesh is dissolved off, leaving a mesh of corrosion channels along the grain boundaries. The creation of corrosion channels accelerates the adhesion of Cl− along the grain boundaries, thus accelerating the corrosion of the grain boundaries. In addition, as Cl− can react with the dissolved Mg2+ in solution and the resulting surface film Mg(OH)2 to produce soluble MgCl2, thereby disrupting the surface film covering the alloy surface to accelerate the dissolution of the magnesium matrix. As shown in Figure 10c, by the late stage of the corrosion–dissolution reaction, when the α-Mg around the grain boundaries and the second phase are continuously dissolved off, until the second phase cannot effectively support the Mg matrix, the matrix phase α-Mg will be dislodged in a block, resulting in a rapid weight loss of this soluble magnesium alloy.

Figure 10: Schematic diagram of the as cast Mg–Gd based alloys corrosion mechanism.
Figure 10:

Schematic diagram of the as cast Mg–Gd based alloys corrosion mechanism.

Comprehensive analysis, the as cast alloy has a faster corrosion rate due to three main reasons: Firstly, the as cast Mg–Gd based soluble magnesium alloy has a greatly increased degradation rate compared to the traditional non-addition of rare-earth magnesium alloy, mainly due to the addition of a suitable amount of rare-earth elements Gd, resulting in an increase in the number of rare-earth second phases formed along the grain boundaries in the cast alloy. Secondly, the corrosion dissolution of the as cast soluble magnesium alloy occurs not only by the corrosion dissolution of the matrix magnesium, but also by the formation of micro-electro-couple corrosion between the second phase and the matrix phase. Finally, as the second phase of the as cast alloy is coarse mesh distribution at the grain boundary, making the formation of a large cathode near the grain boundary small anode of the primary cell reaction, accelerating the intergranular corrosion, thus causing the phenomenon of grain shedding.

4.2 Study of the corrosion mechanism of EP deformed Mg–Gd based alloys

The analysis of the microstructure of the EP deformed soluble magnesium alloy shows that the extrusion + perforation deformation not only destroys the coarse reticulated second phase in the as cast texture, but also makes the grain refinement. In metallurgical microscopy shots, it was found that lamellar texture was found within the grains of the EP deformed alloy. EDS and XRD testing of the EP deformed alloy revealed the presence of small amounts of slate-like Mg-Gd-Ni phase and rare earth-rich square phases within the grains of the morphological alloy. These Mg-Gd-Ni phases formed inside the grains are very susceptible to micro-electro-coupling corrosion with the magnesium matrix. Figure 11 shows a schematic diagram of the dissolution mechanism of the EP deformed soluble magnesium alloy.

Figure 11: Schematic diagram of the EP deformed Mg–Gd based alloys mechanism of corrosion.
Figure 11:

Schematic diagram of the EP deformed Mg–Gd based alloys mechanism of corrosion.

As shown in Figure 11a, when the EP deformed Mg–Gd based soluble magnesium alloy is immersed in KCl solution, micro-electro-couple corrosion is first formed on the surface of the specimen inside the grains where there is a potential difference, leading to local defects within the grains. As shown in Figure 11b, with the multiple effects of micro-electro-couple corrosion and the dissolution of Mg within the grain with Cl−, corrosion channels appear within the grain. The creation of channels accelerates the entry of Cl−, forming a pitting core within the grain and accelerating the process of pitting reactions within the grain. Although Mg(OH)2 product film is produced on the surface where Mg matrix dissolution occurs, local corrosion begins to expand rapidly in the region where micro-electro-couple corrosion is formed, as the reaction is more intense and the product film becomes loose and porous, making it difficult to deposit on the specimen surface. As shown in Figure 11c, in the film-free region of the EP deformed alloy, a dual effect of dissolution of the intracrystalline Mg matrix and a small amount of intracrystalline micro-electro-coupling corrosion reaction occurs rapidly, and is accompanied by the continuous dissolution and shedding of intracrystalline α-Mg, which eventually forms large corrosion pits on the surface of the metamorphic soluble Mg alloy. This has the same characteristics as the previous macroscopic morphology of corrosion of EP deformed alloys. In a comprehensive analysis, the dissolution mechanism of the EP deformed soluble Mg alloy is influenced by the initial texture, mainly by intracrystalline corrosion, where the corrosion type is: a superposition process of intracrystalline α-Mg dissolution and microelectric coupling corrosion.

5 Conclusions

The effect of different initial texture states, the microstructure, surface volta potential, surface morphology, immersion tests, and electrochemical tests of Mg–Gd based soluble magnesium alloys were investigated, respectively, and finally a model for the corrosion mechanism was established. The main conclusion are as follows:

  1. Compared to EP deformed soluble magnesium alloys, the Mg5Gd phase of as cast alloys has a continuous coarse network structure distributed between grain boundaries, which can form a large cathode and small anode galvanic couple corrosion with the Mg matrix during corrosion. After EP deformation, the rough Mg5Gd phase of the original casting structure is not only destroyed, but also the grains and grain boundaries are refined, resulting in a weakened micro galvanic couple effect.

  2. At the beginning of the dissolution reaction, the second phase in the Mg–Gd based soluble magnesium alloy will act as the anode for micro-electro-couple corrosion while the magnesium matrix acts as the cathode for micro-electro-couple corrosion, forming an electro-couple corrosion.

  3. The initial texture state not only affects the corrosion behavior of the Mg–Gd based soluble magnesium alloy but also affects its corrosion mechanism. Among them, the corrosion mechanism of the as cast Mg–Gd based soluble magnesium alloy is mainly intercrystalline corrosion, while the corrosion mechanism of the EP deformed is mainly intracrystalline corrosion.

Corresponding author: Zhibing Chu, School of Materials Engineering, Taiyuan University of Science and Technology, Taiyuan, Shanxi Province030024, China, E-mail:

Funding source: The basic research program of Shanxi Province in 2022

Award Identifier / Grant number: 202203021211208

Funding source: The Key Research and Development Program of Shanxi Province

Award Identifier / Grant number: 201903D121088

Funding source: The Applied Basic Research Project of Shanxi Province

Award Identifier / Grant number: 20210302123203

  1. Author contributions: All the authors have accepted responsibility for the entire content of this submitted manuscript and approved submission.

  2. Conflicts of interest: The authors declare no conflicts of interest.

  3. Research funding: The authors gratefully acknowledge the financial support of the basic research program of Shanxi Province in 2022 (free exploration category) (no. 202203021211208), the Key Research and Development Program of Shanxi Province (no. 201903D121088) and the Applied Basic Research Project of Shanxi Province (grant no. 20210302123203).

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Received: 2023-03-17
Accepted: 2023-07-17
Published Online: 2023-08-14
Published in Print: 2023-12-15

© 2023 Walter de Gruyter GmbH, Berlin/Boston

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https://doi.org/10.1515/corrrev-2023-0035
Keywords for this article
corrosion mechanism; intercrystalline corrosion; intracrystalline corrosion; Mg–Gd based soluble Mg alloy; surface morphology

Articles in the same Issue

  1. Frontmatter
  2. Reviews
  3. Organic compounds as corrosion inhibitors for reinforced concrete: a review
  4. The role of microbes in the inhibition of the atmospheric corrosion of steel caused by air pollutants
  5. A review on corrosion and corrosion inhibition behaviors of magnesium alloy in ethylene glycol aqueous solution
  6. Original Articles
  7. Study of the corrosion mechanism of Mg–Gd based soluble magnesium alloys with different initial texture states
  8. Determination of corrosion product film on pure Mg in Cl− environment using XPS etching
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