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Article

Microstructure and Properties of Novel Mg-Al-Zn-Mn-Ca-Ni Dissoluble Alloy Fabricated by Industrial Two-Step Extrusion Method

by
Jian Wang
1,
Hongxiang Li
1,2,*,
Jinhui Wang
1,
Yaohong Liu
1 and
Jishan Zhang
1,2,*
1
State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, China
2
Beijing Laboratory of Metallic Materials and Processing for Modern Transportation, University of Science and Technology Beijing, Beijing 100083, China
*
Authors to whom correspondence should be addressed.
Metals 2022, 12(4), 583; https://doi.org/10.3390/met12040583
Submission received: 2 March 2022 / Revised: 21 March 2022 / Accepted: 24 March 2022 / Published: 30 March 2022
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Abstract

:
Dissoluble magnesium alloys for fabrication of fracturing tools have received increasing attention in recent years. However, most of the existing research is focused on the small-sized samples prepared in the laboratory, and there is almost no report on the industrial dissoluble magnesium alloys. In this study, large-scale Mg-Al-Zn-Mn-Ca-Ni alloys with a diameter of 110 mm were prepared by a semi-continuous casting and two-step extrusion method, and the corresponding microstructure and mechanical and corrosion properties were also investigated. It was found that after two-step extrusion, the mainly precipitate phases in the Mg-Al-Zn-Mn-Ca-Ni alloy are bulk-like AlMnNi, strip-like Al3Ni, and granular-like and lamellar-like Mg17Al12 phases. Due to the combined effects of grain refinement and precipitation strengthening, the Mg-Al-Zn-Mn-Ca-Ni alloy obtained excellent mechanical properties after two-step extrusion, and its ultimate tensile strength, yield strength, and elongation were 314.6 MPa, 191.2 MPa, and 13.1%, respectively. Moreover, the corrosion rate of the alloy in 3 wt.% KCl at 93 °C was as high as 97.61 mg·cm−2·h−1. This work provides a high-performance, low-cost, and large-scale alloy product for the fabrication of dissoluble fracturing tools.

1. Introduction

With the development of global industry, the consumption of oil and gas resources is increasing. This puts forward higher requirements for the exploitation of oil and gas resources. In recent years, multi-stage hydraulic fracturing technology has been widely used in the exploitation of oil and gas resources [1,2,3]. Fracturing tools have gradually evolved from drillable fracturing tools to soluble fracturing tools [4,5].
Magnesium alloys have been widely used in the fields of aerospace, automotive, 3C (computers, communications, and consumer electronics), biomedical, etc. due to their excellent properties such as low density, high specific strength, high specific stiffness, and good bio-compatibility [6,7,8]. On the other hand, magnesium alloy is an excellent material for dissoluble fracturing tools due to its good mechanical properties and high corrosion rate in corrosive solutions [9,10,11]. Han et al. studied the effects of different volume fractions of Ni-containing, long-period stacking-ordered (LPSO) phase on mechanical properties and corrosion properties in the Mg-Gd-Ni alloy [12]. They found that the LPSO phase can improve the strength of the alloy. Meanwhile, due to the micro-galvanic reaction of the Ni-containing LPSO phase and Mg matrix, the corrosion rate of the alloys was also accelerated. Wang et al. prepared an as-cast Mg97.5Ni1Y1.5 alloy with a high corrosion rate (6675.32 mm/a at 93 °C) and excellent mechanical properties by introducing a particular Mg12NiY (LPSO) phase [13]. Summarizing the reported dissoluble magnesium alloys so far, it was found that almost all the dissoluble magnesium alloys with high tensile fracture strength and a decent corrosion rate contain a large amount of RE (rare earth) elements, leading to higher cost. On the other hand, some researchers also tried to develop low-cost Mg-Al series dissoluble alloys [9,14,15,16,17,18,19,20,21]. Wang, Chen, Geng et al. developed novel dissoluble magnesium alloys by adding Si, Cu, Gd, and other elements into the Mg-Al alloys [14,15,17]. These alloys all exhibited an excellent corrosion rate, but the plasticity of the alloys was poor (<10%) due to the high contents of Al element (>15 wt.%); the corresponding tensile strength was not measured [14,15,16,17]. In addition, the abovementioned Mg-RE and Mg-Al (high Al content) alloys are not easily deformed; this makes it difficult to meet the needs of industrial application. Thus, it is essential to develop a novel industrial dissoluble magnesium alloy with low cost and high strength–ductility balance.
Grain refinement is an effective method to improve the mechanical properties of magnesium alloys [22]. Equal-channel angular pressing (ECAP) and high-pressure torsion (HPT) have been used to fabricate magnesium alloys with ultra-fine grains (UFG) [23,24]. However, these severe deformation processes are not suitable for large-scale industrial production. In the conventional extrusion process, the main parameters affecting the properties of deformed alloys include extrusion ratio, extrusion temperature, and extrusion speed. High-performance magnesium alloy can also be achieved by controlling these parameters [25,26,27]. Nakata et al. developed a novel Mg-1.3Al-0.3Ca-0.4Mn (wt.%) alloy that can be extruded at a very high speed (24 m/min) [28]. The high number density of monolayer Guinier–Preston zones dispersed in the alloy allows it to exhibit a good combination of strength (tensile 0.2% proof stress of 287 MPa) and elongation (20%) after age hardening. Li et al. prepared a high-strength AZ91D alloy by introducing ultrafine grains, submicron-scale precipitates, and strong texture into the alloy by low-temperature (200–250 °C) extrusion [22]. For fabrication of fracturing tools, providing large-sized (about 100 mm in diameter), dissoluble magnesium alloy bars is obligatory. However, the current studies on the dissoluble magnesium alloys are mostly focused on the laboratory scale; the corresponding research results are difficult to apply in industrial production environments. In industrial production, limited by the extrusion machine, large-sized dissoluble magnesium alloys can only be prepared with a low extrusion ratio, which makes it difficult to achieve high performance. Thus, seeking an effective industrial fabrication method for the high-performance dissoluble magnesium alloys will be necessary. Some researchers have studied the effect of two-step extrusion on the microstructure and properties of Mg-Zn-Y-Nd alloys [29]. The results showed that, after two-step extrusion, the texture intensity was significantly weakened and more uniform and finer grains can be attained by this method. As a result, the ultimate tensile strength (UTS), tensile yield strength (TYS), and elongation of the alloy increased from 226.2 MPa, 159.7 MPa, and 24.5% to 241.8 MPa, 196.5 MPa, and 29.7%, respectively. That is, by using the two-step extrusion method, although the extrusion ratio of each extrusion process is not high, the alloy can obtain a larger final extrusion ratio (the ratio of the ingot cross-sectional area of the two-step extrusion alloy). Then, a magnesium alloy with high mechanical properties can be obtained in a convenient process. However, so far, there is no report on the two-step extrusion method applied in the dissoluble magnesium alloys, although it can provide an excellent alternative for the fabrication of industrial, large-scale, high-performance dissoluble magnesium alloy bars.
For the abovementioned aim, in this study we developed a novel industrial Mg-Al-Zn-Mn-Ca-Ni dissoluble alloy with low cost, high mechanical properties, and a decent corrosion rate prepared by the two-step extrusion method. The corresponding research results including the microstructure and mechanical and corrosion properties were investigated in detail and the mechanisms of the strength–ductility balance and corrosion were also explored.

2. Materials and Methods

The Mg-8.3Al-0.7Zn-0.2Mn-0.09Ca-0.15Ni (wt.%) alloys’ ingots with a diameter of 380 mm were prepared by semi-continuous casting. The casting temperature was 730 °C, and the casting speed was 5 cm/min. The ingots were homogenized at 380 °C for 20 h and then quenched in air. The homogenized ingots were preheated at 360 °C for 30 min. Then, the first-step extrusion was carried out at 360 °C with an extrusion ratio of 2:1 and an extrusion rate of 0.4 m·min−1. The second-step extrusion was performed at 360 °C with an extrusion ratio of 4:1 and an extrusion rate of 0.6 m·min−1. Here, the extrusion ratio is defined as the ratio of the cross-sectional area of the samples before and after extrusion. The alloys after the first-step extrusion and second-step extrusion are called AZN800-A and AZN800-B, respectively. The schematic diagram of the preparation process of the alloys and the photos of the alloy bars after the two-step extrusion are shown in Figure 1.
The specimens were ground and polished for microstructure observation using scanning electron microscope (SEM, Phenom XL, Phenom-World, Eindhoven, The Netherlands) equipped with an energy-dispersive X-ray spectrometer (EDS). The specimens were ground and finally electro-polished in an ACII solution at −20 °C with a 20-V applied potential and a 0.2-A current for electron backscatter diffraction (EBSD) tests.
The dog-bone samples were processed along the extrusion direction (ED) for tensile tests. The tests were performed using a uniaxial tensile test with a strain rate of 1 × 10−3 s−1 at ambient temperature. Three specimens of each alloy were tested for the tensile tests. The morphologies of the fracture surface were observed by SEM.
Immersion tests and electrochemical tests were carried out to evaluate the corrosion behaviors of the alloys. Immersion tests were performed in 3 wt.% KCl at 25 °C and 93 °C, respectively, according to American Society of Testing Materials standards G31-72. After immersion, the samples were immersed in a chromic acid solution of 200 g/L CrO3 + 10 g/L AgNO3 for about 20 min to remove the corrosion products for SEM observation. The electrochemical tests were performed in 3 wt.% KCl solution at 25 °C with a three-electrode system using a CORRTEST CS310H electrochemical workstation (CorrTest, Wuhan, China). The samples with an exposed surface area of 1 cm2, the saturated calomel electrode (SCE), and platinum foil were used as the working electrode, the reference electrode, and counter electrode, respectively. The specimens were immersed for 300 s to ensure the stability of the open circuit potential. Electrochemical impedance spectroscopy (EIS) tests were performed with a frequency range between 100 kHz and 0.1 Hz. The polarization curve tests were carried out in the potential range between −2.0 VSCE and −1.0 VSCE using a scan rate of 1 mV/s. The data of the electrochemical tests were analyzed by the CS studio software (Ver. 5.4, CorrTest, Wuhan, China). All the tests were repeated at least three times. To measure the Volta potential between the precipitate phases and α-Mg matrix, a scanning Kelvin probe force microscopy (SKPFM, Bruker Scientific Instruments Hong Kong Co. Limited, Hong Kong, China) was used. The experiments were conducted at ambient temperature under a relative humidity of 40 ± 5%.

3. Results and Discussion

3.1. Microstructure Characterization

Figure 2 shows the backscattered electron (BSE) images of the extruded alloys after finishing the first and second extrusion, respectively. It can be seen from Figure 2a that the AZN800-A alloy contained three precipitated phases, i.e., a white block-like phase, a white strip-like phase, and a gray irregular-like phase. The white block-like phases were aggregated in the alloy, while the white strip-like phases were uniformly distributed along ED. Both the white block-like phase and the white strip-like phase were broken after extrusion. Moreover, after the second-step extrusion, as shown in Figure 2d, the white block-like phase and white strip-like phase in the AZN800-B alloy were more fragmented. Further observation of the alloys is shown in Figure 2b,e, and it can be seen that there were gray granular-like and lamellar-like phases distributed in the alloys. Additionally, there were more gray granular-like phases in the AZN800-B alloy after the two-step extrusion.
Furthermore, EDS analysis was performed on the extruded alloys to characterize the precipitated phase, and the results are shown in Figure 2c,f and Table 1. From the results, there was no obvious difference in the element content of the precipitate phases in the AZN800-A and AZN800-B alloys, which indicated that the two-step extrusion did not change the type of the precipitate phases. The white block-like phase mainly contained about 60 at.% Al-20 at.% Mn-10 at.% Ni-10 at.% Mg. Generally, in Mg-Al alloys, adding a small amount of Mn element can only form a small-sized Al-Mn phase [30,31]. In this study, the addition of Ni promoted the precipitation and growth of the Al-Mn phase, forming a large-sized AlMnNi phase marked by the white arrow in Figure 2c,f. The white strip-like phase mainly contained about 35 at.% Al-10 at.% Ni-50 at.% Mg, which can be speculated to be the Al3Ni phase, marked by the yellow arrow in Figure 2c,f. The high Mg content may have been due to the signal collected from the matrix. The gray irregular-like phase, composed of about 66 at.% Mg-32 at.% Al-1 at.% Zn, was the Mg17Al12 phase, which commonly exists in Mg-Al alloys [32,33]. According to the above SEM results, both of the AZN800-A and AZN800-B alloys mainly contained three phases besides the Mg matrix, i.e., the AlMnNi phase, Al3Ni phase, and Mg17Al12 phases, but the content, morphology, and distribution of precipitated phase were different for the two-step extrusion process.
Furthermore, the orientation maps, pole figures (PFs), and inverse pole figures (IPFs) are depicted in Figure 3 to examine the texture of the alloys. As shown in Figure 3a, complete dynamic recrystallization occurred in the AZN800-A alloy during the first-step extrusion process, and the average grain size of the AZN800-A alloy was 11.49 μm. After the dynamic recrystallization, the AZN800-A alloy exhibited a typical basal fiber texture with (0002)//ED and the (0002) basal plane was mainly distributed in the transverse direction (TD) and normal direction (ND) [34]. The texture maximum intensity of the alloy was 8.12. A strong <1010>, a relatively weaker <2110>, and an extremely weak <0001> texture components were observed, see Figure 3e. For the AZN800-B alloy, the grains were even finer, with an average grain size of 7.97 μm, after the second-step extrusion. Similar to the AZN800-A alloy, the (0002) basal plane of the AZN800-B alloy was mostly distributed between TD and ND, but it was more dispersed. However, some of the basal planes were tilted to be perpendicular to ED. Meanwhile, the texture maximum intensity increased to 8.74. The IPF (Figure 3f) showed that the texture components may have been a mixed morphology composed of <0001>, <1010>, and <2110> parallel to ED. Particle-stimulated nucleation (PSN) is very important for the recrystallization process of magnesium alloys during hot deformation [35]. In this study, a large number of phases were precipitated in the alloy after the first-step hot extrusion. The PSN of recrystallization was activated, resulting in complete dynamic recrystallization and the formation of a typical basal fiber texture with (0002) plane parallel to ED. Some grains with basal planes perpendicular to ED were also produced in the alloys, which may have been because of the inhibiting effect of Al-rich phases (AlMnNi and Al3Ni) during the recrystallization. Since the alloy underwent complete recrystallization after the first-step extrusion, there was no significant change in the texture of the alloy after the second-step extrusion.

3.2. Mechanical Properties

The engineering tensile stress–strain curves and the fracture morphology of the AZN800-A and AZN800-B alloys are shown in Figure 4 and Figure 5. The UTS, TYS, and elongation of the AZN800-A alloy were 289.6 ± 1.0 MPa, 164.6 ± 4.5 MPa, and 9.0 ± 0.5%, respectively. After the second-step extrusion, the mechanical properties of the AZN800-B alloy were significantly improved. The UTS, TYS, and elongation of the AZN800-B alloy increased to 314.6 ± 0.6 MPa, 191.2 ± 6.3 MPa, and 13.1 ± 1.0%, respectively. From the fracture morphologies of the alloys (Figure 5a,b), it could be seen that the AZN800-A alloy showed a mixed fracture behavior, containing both cleavage planes and dimples. On the other hand, the fracture morphology of the AZN800-B alloy was almost dominated by dimples. From the EDS images of the fracture (Figure 5c,d), large-size, bulk-like AlMnNi, and strip-like Al3Ni phases (marked by black arrows) can be observed in the fracture of the AZN800-A alloy. However, for the AZN800-B alloy, only the fine AlMnNi and Al3Ni phases were observed.
It can be seen from the results of the tensile tests in Figure 4 that the mechanical properties of the alloys were significantly improved by the two-step extrusion. The tensile strength of the AZN800-B alloy was about 45 MPa higher than that of the AZN800-A alloy. The increase in strength was caused by grain refinement and precipitation strengthening [36,37]. After the two-step extrusion, the grain size of the AZN800-B alloy was refined, as shown in Figure 3, and the content of the precipitated Mg17Al12 phase was also increased, as shown in the SEM images (Figure 2). It was reported that, during the tensile experiment, the stress was easily concentrated around the second phase, and the brittle second phase became the initiation site and the path for the propagation of micro-cracks, resulting in failure of the sample [36]. In this study, for the AZN800-A alloy, a large amount of AlMnNi and Al3Ni phases existed in the alloy, which led to the initiation and propagation of micro-cracks and finally resulted in the fracture fail of the samples. Indeed, more brittle phases (marked by black arrows in Figure 5c) were also observed in the fracture. However, the AlMnNi and Al3Ni phases were broken after the second-step extrusion, and only a few brittle phases were observed in the fracture of the AZN800-B alloy (Figure 5d). The deterioration effect of the brittle phases was alleviated and, thus, the plasticity of the AZN800-B alloy increased. In addition, for the extruded magnesium alloys, the texture also affected the mechanical properties of the alloy. At room temperature, the main slip systems of magnesium alloys were basal <a> slip, prismatic <a> slip, and pyramidal <c + a> slip [38,39,40]. The critical resolved shear stress of prismatic <a> slip and pyramidal <c + a> slip was much higher than basal <a> slip at room temperature [41]. Consequently, basal slip was the main deformation mode at room temperature [42]. Both the AZN800-A and AZN800-B alloys showed a typical basal fiber texture ((0002)//ED) after hot extrusion. During the deformation process, the alloys were dominated by base slip. According to the results in Figure 3, the orientation in the ANZ800-A alloy showed more <1010>//ED, while the orientation in theAZN800-B alloy showed more <2110>//ED. More base slip occurred in the AZN800-B alloy during the deformation process. This resulted in the higher plasticity of the AZN800-B alloy. Therefore, the AZN800-B alloy exhibited the optimized balance of strength and plasticity.

3.3. Corrosion Properties

Figure 6 presents the weight loss rate of the alloys after immersion in 3 wt.% KCl at different temperatures. Both of the AZN800-A and AZN800-B alloys exhibited a relatively low corrosion rate (about 6 mg·cm−2·h−1) at 25 °C. When the corrosion condition was changed to 93 °C, the corrosion rate of the alloys was significantly increased. The weight loss rate of the AZN800-A alloy was as high as 100.06 mg·cm−2·h−1. Additionally, the weight loss rate of the AZN800-B alloy was only slightly lower than that of the AZN800-A alloy. The results of the immersion experiments showed that the two-step extrusion did not have an obvious influence on the corrosion rate of the alloys. In contrast, both the AZN800-A and the AZN800-B alloys were all sensitive to the temperature of the corrosive medium.
The polarization curves and related fitting results of the AZN800-A and AZN800-B alloys’ immersion in 3 wt.% KCl at 25 °C are given in Figure 7 and Table 2. The corrosion potential (Ecorr) of the alloys did not change much, with a value of −1.436 ± 0.02 VSCE for AZN800-A and −1.447 ± 0.01 VSCE for AZN800-B. The value of Ecorr represents the thermodynamic tendency of the corrosion process [43,44], and the similar Ecorr value indicates that the two alloys showed the same corrosion tendency. The corrosion current density (Icorr), the slope of the cathodic polarization curve (βc), and the slope of the anodic polarization curve (βa) can all represent the corrosion rate of the alloys [45]. For magnesium alloys, the anodic curves represent the dissolution of magnesium (Mg-2e → Mg2+), and the cathodic polarization curves represent the hydrogen evolution through water reduction (2H2O + 2e → H2↑ + 2OH−) [45]. From the polarization curve shown in Figure 8, it can be seen that the difference of the anodic and cathodic curves between the AZN800-A and AZN800 alloys was relatively small, indicating that the corrosion performance of the two alloys under this test condition was similar.
Furthermore, the EIS results and equivalent circuits of the AZN800-A and AZN800-B alloys measured in 3 wt.% KCl at 25 °C are shown in Figure 8. As shown in Figure 8a, both of the AZN800-A and AZN800-B alloys exhibited three capacitive loops: one in high-frequency range, one in medium-frequency range, and another inductive loop in low-frequency region. These characteristics were also able to be verified in the Bode phase angle plots (Figure 8c). The high-frequency and medium-frequency capacitance loops are often related to the charge transfer resistance of the electric double layer and the formation of oxide film, respectively [46,47,48]. The low-frequency inductance loop may be attributed to the localized or pitting corrosion [49,50]. The AZN800-B alloy had the larger capacitive loop diameter as well as the larger impedance modulus |Z| (Figure 8b), indicating that the corrosion rate of the alloy was slightly slower.
The equivalent circuit fitted with the EIS spectrum is shown in Figure 8d, and the fitting results are listed in Table 3. Rs and Rct represent the solution resistance between the reference electrode and working electrode and the charge transfer resistance, respectively [44]. The value of Rs did not change much during the test. Rf is the film resistance [47]. Constant phase elements, CPE1 and CPE2, are used to replace the double layer and film capacitance, including the capacitance changes caused by various indicated factors [51]. L1, L2, R1, and R2 refer to inductance and electric resistance, respectively, which correspond to pitting corrosion [52]. Due to the different preparation processes, the fitting results of the AZN800-A and AZN800-B alloys were also different, but the changes were not obvious. They revealed that the difference in corrosion rates between the two alloys was not significant, which is in accordance with the abovementioned immersion tests.
The results of the immersion test and electrochemical experiment showed that the corrosion properties of the alloys did not change significantly after the two-step extrusion. Furthermore, the corrosion process of the AZN800-B alloy was explored. Generally, the potential between the second phase and Mg matrix can play an important role in corrosion behaviors [53]. The Volta potential maps of the different precipitate phases in the AZN800-B alloy are shown in Figure 9. From Figure 9, it can be seen that the Volta potential difference between the precipitate phase and the Mg matrix was about 500 mV for AlMnNi, 200 mV for Al3Ni, and 150 mV for Mg17Al12, respectively. Therefore, the high-potential precipitate phases in the alloy can act as the cathode, while the low-potential Mg matrix acts as the anode; the formation of this kind of micro-galvanic reaction between the anode and cathode can play an important role in the corrosion behavior of the AZN800-B alloy.
The surface micrographs of the AZN800-B alloy immersed in 3 wt.% KCl at 25 °C for 2 h and at 93 °C for 10 min are shown in Figure 10. It can be seen from Figure 10a that the surface of the AZN800-B alloys has still some uncorroded areas after immersion at 25 °C for 2 h. After removing the corrosion products (Figure 10b), the corrosion morphology of the AZN800-B consisted of uncorroded areas and corrosion grooves. Furthermore, as shown in the BSE result (Figure 10c), the precipitate phases in the alloys had a larger impact on the corrosion morphology. There were lots of precipitate phases (including AlMnNi, Al3Ni, and granular-like Mg17Al12 phases) at the edge and inside of the uncorroded area. The results showed that the Al-rich precipitate phases, especially AlMnNi and Al3Ni phases, had a certain inhibitory effect on the corrosion of the AZN800-B alloy at 25 °C. From the results in Figure 6, it was found that the corrosion rate of the AZN800-B alloy changed greatly at 93 °C. The corrosion morphology of the alloy had also an obvious change. The surfaces of the AZN800-B (Figure 10d) were almost completely covered by corrosion products. After clearing the corrosion products (Figure 10e), it was seen more clearly that the surface of the alloy was dominated by corrosion pits and only a small number of precipitation-enriched regions remained uncorroded. Different from the corrosion morphology at 25 °C, only AlMnNi and Al3Ni phases were observed in the uncorroded area at 93 °C and almost no Mg17Al12 phase was observed.
Based on the above results, the AZN800-B alloy exhibited still excellent corrosion properties after the two-step extrusion. The factors affecting the dissolution of the AZN800-B alloy were as follows. (1) The addition of Ni element in the alloy: Magnesium alloys are easily corroded due to their highly negative electromotive force (emf = −2.363 V) and the poor protective properties of the surface film [54]. The addition of Ni can further accelerate the corrosion of the alloy, even in very small amounts [55]. (2) The complex influence of the precipitate phases on the corrosion process of the alloy: On the one hand, the high-potential precipitate phases (including AlMnNi, Al3Ni, and Mg17Al12 phases) can form a micro-galvanic reaction with the matrix and accelerate the corrosion of the alloy [21]. On the other hand, the morphology of the precipitate phases has a significant effect on the dissolution of the alloy. The broken AlMnNi, Al3Ni, and lamellar-like Mg17Al12 can significantly accelerate the corrosion process of the matrix and form a large number of corrosion pits. However, the larger-size, bulk-like AlMnNi, strip-like Al3Ni, and granular-like Mg17Al12 may act as barriers during the corrosion process; therefore, many uncorroded areas are formed in the alloy. Under the combined effect of the precipitate phases, the corrosion rate of the alloy at 25 °C was higher than that of traditional magnesium alloys but lower than that of some dissoluble magnesium alloys. It is worth noting that the corrosion rate of the alloy was greatly improved at 93 °C. By observing the morphology of the alloy after being immersed at 93 °C (Figure 10e,f), it was found that the AlMnNi and Al3Ni phases still existed as barriers at the edge in the uncorroded region, while the granular-like Mg12Al12 phase was not observed. This may have been because the micro-galvanic corrosion between the granular-like Mg17Al12 phase and the matrix was aggravated at high temperature and the granular-like Mg17Al12 phase no longer acted as a barrier to inhibit the corrosion of the alloy. Therefore, the corrosion rate of the AZN800-B alloy can be as high as 97.61 mg·cm−2·h−1 at 93 °C.
Table 4 compares the mechanical and corrosion properties of the Mg-Al-Zn-Mn-Ca-Ni alloy developed by our group with other reported typical dissoluble magnesium alloys [12,14,15,16,17,18,44,53,56,57,58]. The UTS, TYS, and ET are the ultimate tensile strength, tensile yield strength, and elongation of tensile test, respectively. The UCS, CYS, and EC are the ultimate compressive strength, compressive yield strength, and elongation of compressive test, respectively. From Table 4, it can be seen that by using the two-step extrusion method, the novel Mg-Al-Zn-Mn-Ca-Ni alloy developed in this study exhibited the outstanding mechanical and corrosion properties, which were much superior to other dissoluble magnesium alloys fabricated on a laboratory scale. To some extent, the comprehensive properties, including mechanical and corrosion, were even comparable with the dissoluble Mg-RE alloys [57]. Meanwhile, considering that no rare earth elements were included, the novel Mg-Al-Zn-Mn-Ca-Ni dissoluble alloy with low cost, high mechanical properties, and a decent corrosion rate prepared by the two-step extrusion process in this study can provide an excellent alternative for the fabrication of dissoluble fracturing tools.

4. Conclusions

In this paper, a novel industrial Mg-Al-Zn-Mn-Ca-Ni dissoluble alloy was successfully prepared by a semi-continuous casting and two-step extrusion method. The alloy achieved excellent comprehensive properties after the two-step extrusion. The conclusions are as follows.
(1)
The mainly precipitate phases in the Mg-Al-Zn-Mn-Ca-Ni alloys were bulk-like AlMnNi, strip-like Al3Ni, and granular-like and lamellar-like Mg17Al12. The types of precipitate phases did not change after the two-step extrusion, but the lager-size precipitate phases were obviously broken and refined after the second-step extrusion. In addition to the change in precipitate phases, the two-step extrusion process also had an effect on the texture of the alloy. The texture distribution of the AZN800-B alloy was more dispersed and the texture intensity was increased.
(2)
The AZN800-B alloy achieved excellent mechanical properties. The ultimate tensile strength, tensile yield strength, and elongation were 314.6 MPa, 191.2 MPa, and 13.1%, respectively.
(3)
Temperature had a great influence on the corrosion susceptibility of the precipitated phase. The corrosion rate of the AZN800-B alloy was 6.04 mg·cm−2·h−1 at 25 °C. However, the AZN800-B alloys exhibited excellent corrosion rates (97.61 mg·cm−2·h−1) at 93 °C.

Author Contributions

Conceptualization, J.W. (Jian Wang) and H.L.; formal analysis, J.W. (Jian Wang) and H.L.; funding acquisition, H.L. and J.Z.; investigation, J.W. (Jian Wang) and H.L.; methodology, J.W. (Jian Wang), J.W. (Jinhui Wang), and Y.L.; validation, H.L. and J.Z.; visualization, J.W. (Jian Wang); writing—original draft, J.W. (Jian Wang); writing—review and editing, H.L. All authors have read and agreed to the published version of the manuscript.

Funding

The Major State Research and Development Program of China (No. 2021YFB3701100, No. SQ2020YFF0405156), the National Natural Science Foundation of China (No. 52171097, No. 51971020), and the Fundamental Research Funds for the Central Universities (No. FRF-IC-20-08) are deeply appreciated for their financial support. The authors also thank the “Dingxinbeike” Project (G20200001105) for the international communication.

Data Availability Statement

Not applicable.

Conflicts of Interest

The authors declare no conflict of interest.

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Figure 1. (a) Schematic diagram of the preparation process of the large-scale dissoluble magnesium alloys. Photos of the extruded alloys: (b) AZN800-A, (c) AZN800-B.
Figure 1. (a) Schematic diagram of the preparation process of the large-scale dissoluble magnesium alloys. Photos of the extruded alloys: (b) AZN800-A, (c) AZN800-B.
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Figure 2. SEM micrographs of the extruded alloys (a–c) AZN800-A, (d–f) AZN800-B.
Figure 2. SEM micrographs of the extruded alloys (a–c) AZN800-A, (d–f) AZN800-B.
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Figure 3. EBSD orientation (IPF-Z) maps and corresponding (0002) pole figures (PFs) and inverse pole figures (IPFs) of the extruded alloys: (a,c,e) AZN800-A, (b,d,f) AZN800-B.
Figure 3. EBSD orientation (IPF-Z) maps and corresponding (0002) pole figures (PFs) and inverse pole figures (IPFs) of the extruded alloys: (a,c,e) AZN800-A, (b,d,f) AZN800-B.
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Figure 4. Engineering tensile stress–strain curves of the AZN800-A and AZN800-B alloys.
Figure 4. Engineering tensile stress–strain curves of the AZN800-A and AZN800-B alloys.
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Figure 5. The tensile fracture topography of the (a,c) AZN800-A and (b,d) AZN800-B alloys.
Figure 5. The tensile fracture topography of the (a,c) AZN800-A and (b,d) AZN800-B alloys.
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Figure 6. Weight loss rate of the AZN800-A and AZN800-B alloys at different temperatures.
Figure 6. Weight loss rate of the AZN800-A and AZN800-B alloys at different temperatures.
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Figure 7. Polarization curves of the AZN800-A and AZN800-B alloys in 3 wt.% KCl at 25 °C.
Figure 7. Polarization curves of the AZN800-A and AZN800-B alloys in 3 wt.% KCl at 25 °C.
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Figure 8. EIS of the AZN800-A and AZN800-B alloys after immersion in 3 wt.% KCl solution: (a) Nyquist plots curves, (b) Bode plots curves, (c) phase angle plots, (d) equivalent circuits.
Figure 8. EIS of the AZN800-A and AZN800-B alloys after immersion in 3 wt.% KCl solution: (a) Nyquist plots curves, (b) Bode plots curves, (c) phase angle plots, (d) equivalent circuits.
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Figure 9. The Volta potential of the precipitate phases in the AZN800-B alloy: (a) The Volta potential of AlMnNi phase; (b) The Volta potential of Al3Ni phase; (c) The Volta potential of Mg17Al12 phase.
Figure 9. The Volta potential of the precipitate phases in the AZN800-B alloy: (a) The Volta potential of AlMnNi phase; (b) The Volta potential of Al3Ni phase; (c) The Volta potential of Mg17Al12 phase.
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Figure 10. Micrographs of the AZN800-B alloys after being immersed in 3 wt.% KCl: (a–c) 2 h at 25 °C, (d–f) 10 min at 93 °C.
Figure 10. Micrographs of the AZN800-B alloys after being immersed in 3 wt.% KCl: (a–c) 2 h at 25 °C, (d–f) 10 min at 93 °C.
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Table
Table 1. The EDS results of the AZN800-A and AZN800-B alloys (at.%).
Table 1. The EDS results of the AZN800-A and AZN800-B alloys (at.%).
PhasesAlloyMgAlMnNiZn
white block-like phaseAZN800-A9.7556.2023.0311.02-
10.1155.5122.0912.29-
AZN800-B10.3056.4921.3011.91-
15.8155.7917.9410.46-
white strip-like phaseAZN800-A49.1838.040.2512.53-
50.2537.630.9811.14-
AZN800-B60.7930.580.238.40-
52.3734.380.5412.71-
gray irregular-like phaseAZN800-A66.5732.23--1.19
66.1132.67--1.22
AZN800-B66.8431.99--1.18
66.3832.30--1.32
Table
Table 2. Tafel fitting results obtained from polarization curves of the AZN800-A and AZN800-B alloys in 3 wt.% KCl at 25 °C.
Table 2. Tafel fitting results obtained from polarization curves of the AZN800-A and AZN800-B alloys in 3 wt.% KCl at 25 °C.
AlloyEcorr (VSCE)Icorr (μA·cm−2)βc (mV)βa (mV)
AZN800-A−1.436 ± 0.0221235.25 ± 66.95296.59 ± 18.9751.00 ± 2.80
AZN800-B−1.447 ± 0.013878.63 ± 53.38273.62 ± 20.4747.89 ± 1.07
Table
Table 3. Fitting results of the EIS spectra of the AZN800-A and AZN800-B alloys.
Table 3. Fitting results of the EIS spectra of the AZN800-A and AZN800-B alloys.
AlloyRs (Ω·cm−2)CPE1 (μF·cm−2)Rct (Ω·cm−2)L1 (H·cm−2)Rf (Ω·cm−2)CPE2 (sn·μΩ−1·cm−2)R1 (Ω·cm−2)L2 (H·cm−2)R2 (Ω·cm−2)
AZN800-A3.7157.85 14.0783.321.7261113.007.8671.21611.64
AZN800-B2.9526.63 40.9415.781.012261.4019.922.10315.69
Table
Table 4. The comparison of mechanical and corrosion properties of the Mg-Al-Zn-Mn-Ca-Ni alloy developed by our group and other reported typical dissoluble magnesium alloys.
Table 4. The comparison of mechanical and corrosion properties of the Mg-Al-Zn-Mn-Ca-Ni alloy developed by our group and other reported typical dissoluble magnesium alloys.
Composition (wt.%)StateUTS (MPa)TYS (MPa)ET (%)UCS (MPa)CYS (MPa)EC (%)Corrosion Rate (mm·day−1)
Mg-8Al-0.8Zn-0.3Mn-0.1Ca-0.2Ni (this study)as-extruded314.6191.213.1---12.66 (3% KCl 93°C)
Mg-17Al-3Zn-5Cu [14]as-cast---4382951.6~11.34 (3% KCl 93 °C) *
Mg-17Al-3Zn-7Cu-1Gd [15]as-cast---4423027.81.51 (3% KCl 93 °C)
Mg-17Al-5Zn-3Si [17]as-cast---3552517.35.39 (3% KCl 93 °C)
Mg-15Al-5Zn-0.25Y [16]as-cast---4172267.20.111 (3% KCl 93 °C)
Mg-2.5Cu-6Al [53]as-cast215.2-10.2378.8-27.35.69 (3% KCl 93 °C)
7Fe/Mg-6Al-1Zn [18]as-extruded---~350~180~9~0.86 (3.5% NaCl 25 °C) *
Mg-4Zn-4Ni [44]as-cast---264.8-20.1~8.52 (3% KCl 25 °C) *
Mg-3Zn-1Y-4Cu [56]as-extruded30224818.23852249.65.52 (3% KCl 93 °C)
Mg-10Gd-3Y-0.2Zr-0.8Ni [57]as-extruded342.9257.815.2596.5296.517.8~3.35 (3% KCl 93 °C)
Mg-6.753Gd-1.66Ni [58]as-cast---340128-5.66 (3% KCl 25 °C)
Mg-16.47Gd-2.05Ni [12]as-cast---266175~10~3.03 (3.5% NaCl 25 °C) *
* The data were calculated based on the results provided in the literature.
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Wang, J.; Li, H.; Wang, J.; Liu, Y.; Zhang, J. Microstructure and Properties of Novel Mg-Al-Zn-Mn-Ca-Ni Dissoluble Alloy Fabricated by Industrial Two-Step Extrusion Method. Metals 2022, 12, 583. https://doi.org/10.3390/met12040583

AMA Style

Wang J, Li H, Wang J, Liu Y, Zhang J. Microstructure and Properties of Novel Mg-Al-Zn-Mn-Ca-Ni Dissoluble Alloy Fabricated by Industrial Two-Step Extrusion Method. Metals. 2022; 12(4):583. https://doi.org/10.3390/met12040583

Chicago/Turabian Style

Wang, Jian, Hongxiang Li, Jinhui Wang, Yaohong Liu, and Jishan Zhang. 2022. "Microstructure and Properties of Novel Mg-Al-Zn-Mn-Ca-Ni Dissoluble Alloy Fabricated by Industrial Two-Step Extrusion Method" Metals 12, no. 4: 583. https://doi.org/10.3390/met12040583

APA Style

Wang, J., Li, H., Wang, J., Liu, Y., & Zhang, J. (2022). Microstructure and Properties of Novel Mg-Al-Zn-Mn-Ca-Ni Dissoluble Alloy Fabricated by Industrial Two-Step Extrusion Method. Metals, 12(4), 583. https://doi.org/10.3390/met12040583

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