Effect of Sc and Zr on Al6(Mn,Fe) Phase in Al–Mg–Mn Alloys
2019 Volume 60 Issue 5 Pages 737-742
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2019 Volume 60 Issue 5 Pages 737-742
The effects of Sc and Zr on the shape, size, and number density of Al6(Mn,Fe) phase in Al–5Mg–0.7Mn alloys were investigated by this article. The results showed that Sc and Zr can form strengthening particles of Al3Sc and Al3(Sc,Zr), which have the effect of refining grains and modifying the morphology and distribution of Al6(Mn,Fe) phase. The OM, SEM, and TEM images indicate that with the formation of primary Al3Sc and Al3(Sc,Zr), the shape of Al6(Mn,Fe) became more regular, size of Al6(Mn,Fe) became smaller, and the number density of Al6(Mn,Fe) increased when compared with base alloy. Zr could motivate the precipitation of primary Al3(Sc,Zr) and Al–5Mg–0.7Mn–0.2Sc–0.2Zr had better refining effect than Al–5Mg–0.7Mn–0.4Sc. The grain refinement and finely dispersed Al6(Mn,Fe) phase can improve the mechanical properties such as yield strength and ultimate tensile strength and the corrosion resistance by the results of intergranular corrosion (IGC) also improved.
Fig. 7 OM of the IGC depth and SEM of the IGC surfaces: (a) S0Z0, (b) S2Z0, (c) S4Z0, (d) S2Z2.
Al–Mg alloys, as a kind of structural material, are typical non-heat-treatable aluminium alloys. Al–Mg alloys are widely used in aircrafts and ships due to their excellent tensile strength, specific gravity, elongation, and corrosion resistance.1–3) Mg is the highest alloying element in Al–Mg alloys. Mg has the solution strengthening effect in aluminium alloys and can improve the work hardening rate of the material. However, with the higher Mg content, there are more precipitations at the grain boundaries, which can reduce the corrosion resistance of the alloy. Hence, the Mg content should be limited to 5% or less.4) However, in industrial production, impure elements are inevitable in the aluminium alloys, which have some negative effects on properties of products. In general, iron is an common impurity in Al–Mg alloys and can precipitate as the Al13Fe4 phase.5,6) Al13Fe4 can cut off the matrix and decrease the mechanical properties. The type, number density, and size of Fe-rich phases in the microstructure have an enormous influence on the combination properties of Al-based alloys.7,8)
To decrease the negative effect of iron in Al–Mg alloys, several effective methods have been adopted. The first is to prevent the formation of Al–Fe phase by decreasing the Fe levels.9) The second is to transform the crystal structure. The third is to refine the intermetallics by addition of some elements such as Mn.10–12) In the presence of Fe element, the addition of Mn can form the Al6(Mn,Fe) phase, which can improve the mechanical properties via precipitation and dispersion strengthening and hence, reduce the negative effects.13) However, research has indicated that Al6(Mn,Fe) can act as a local cathode when compared with the matrix, which decreases the corrosion resistance of Al–Mg alloys.14)
Addition of Sc to Al-based alloys have significant grain refinement effect.15–17) It can improve mechanical properties via precipitation hardening.18,19) Mg and Sc reduce the solid solubility of each other in aluminium,20) which can improve the precipitation effect. The effects of Sc elements can be improved further by the addition of Zr which results in the formation of Al3(Sc,Zr) precipitates.21) Mg can reduce the lattice mismatch between the Al matrix and the Al3Sc particles.22,23) Therefore, when both Sc and Zr are added, these precipitates become even more stable at high temperatures.
Sc and Zr are commonly used for improving mechanical properties of Al–Mg alloys. But there has been limited investigation on the effect of Sc and Zr addition on the Al6(Mn,Fe) phase in Al–Mg–Mn alloys.24,25)
In this work, we investigated the effect of Sc and Zr addition on the Al6(Mn,Fe) phase in Al–5Mg–0.7Mn.
Four experimental Al–Mg–Mn alloys with different Sc and Zr contents were prepared. High-purity Al (99.99%), High-purity Mg (99.99%), High-purity Mn (99.5%), Al–2Sc, and Al–4Zr were used for melting. The melt was held at 780°C for 30 min and degassed for 10 min. It was then poured into a Water-cooled steel mold. The average cooling rate is 20°C/s. The materials were homogenized at 460°C for 24 h to improve the as-cast structure. Then, they were rolled from 20 to 4 mm at 400°C and 4 to 2 mm at room temperature. Finally, the annealing was carried out at 300°C. Table 1 show the chemical compositions of the four alloys.
The microstructure of the alloys was observed using an optical microscope (OM) and a Scanning Electron Microscope (SEM) equipped with an energy dispersive X-ray (EDX). Advanced Transmission Electron Microscope (TEM) characterizations were conducted using a JEM-2100F operating at 300 kV. The TEM samples were made by Gatan PIPS 691. The Yield strength (YS), ultimate tensile strength (UTS), and elongation were measured by MTS810. Hardness tests were conducted by Vickers hardness tester.
The intergranular corrosion (IGC) tests were prepared according to the Standard GB/T 7998–2005. The samples were immersed in an acid salt solution (0.5 mol/L NaCl, 0.1 mol/L HCl) for 24 h at (35 ± 2)°C. The maximum corrosion depth was evaluated by optical microscope.
The electrochemical tests were conducted in 3.5% NaCl electrolyte by using CHENHUA 660C. The working electrode is 1 cm2 while the scanning rate of polarization curves was 1 mV/s.
As shown in Fig. 1, S0Z0 had explicit dendritic structure while the other three alloys were typical equiaxed crystals. There were some black second phases in all alloys. Research has shown that they are Al6(Mn,Fe) intermetallics which mostly precipitate on grain boundaries at cast. The average grain size of four alloys were 110 µm, 100 µm, 39 µm, and 37 µm, and the grain refining effect was relatively clearer in S4Z0 and S2Z2. S4Z0 alloy had higher grain refining effect when compared with S2Z0. When the Sc content was high (0.4%), some part of Sc precipitated in the form of primary Al3Sc, which can refine the as-cast microstructure as non-uniform nucleation core. Although the S2Z0 alloy was no longer dendritic structure, when compare with the alloy S0Z0, its grain size was not much reduced. This phenomenon was due to the low content of Sc (0.2%) which could only dissolve in the Al-matrix, hindering the precipitation of primary Al3Sc phase in the as-cast state. There are small difference between the grain sizes of S2Z2 and S4Z0. This is because Sc and Zr formed Al3(Sc,Zr) composite particles in the as-cast state. These kinds of particles refine grains as non-uniform nucleation core during alloy solidification, as in the case of Al3Sc phase.
As-cast microstructure: (a) S0Z0, (b) S2Z0, (c) S4Z0, (d) S2Z2.
Figure 2(a)–(d) show the SEM images of the homogenized structure. There were some huge rod-like white intermetallic phases (Fig. 2(a)) in the base alloy. As the results of EDX (Fig. 2(e)), the phase could be identified as Al6(Mn,Fe). In the case of S2Z0 alloy with little Sc (0.2 mass%), the volume fraction of the Al6(Mn,Fe) phase became lower compared to that for S0Z0 and the morphology also become more regular from the rod to lump. With the grain refinement, the precipitation of second phase became easier. In the case of alloys S4Z0 and S2Z2, the Al6(Mn,Fe) became finely dispersed in grain boundaries due to many primary Al3Sc and Al3(Sc,Zr) precipitating as non-uniform nucleation cores (can be identified in Fig. 2(f)). The addition of Sc and Zr could refine grains and promote the precipitation probability of Al6(Mn,Fe), which made Al6(Mn,Fe) uniformly distributed in the grain boundaries and improved the properties.
SEM images of the samples after annealing: (a) S0Z0, (b) S2Z0, (c) S4Z0, (d) S2Z2; EDX of intermetallic phase: (e) EDX of Al6(Mn,Fe), (f) EDX of Al3Sc.
Figure 3 and 4 show the TEM images of Al–Mg–Mn alloys after annealing treatment. As shown in Fig. 3, the main second phase in the alloy were Al6 (Mn, Fe), Al3Sc, and Al3 (Sc, Zr). Al6 (Mn, Fe) mainly precipitated on the grain boundary while secondary Al3Sc and Al3 (Sc, Zr) mainly precipitated in the inner grain. The Al6(Mn,Fe) were the square shaped phase in Fig. 4 while the Al3Sc and Al3 (Sc, Zr) were the phase in Fig. 3(d).
Grain boundries of the alloys: (a) S0Z0, (b) EDS maps of the area in (a), (c) and (d) S2Z2.
TEM images of the alloys: (a) S0Z0, (b) S2Z0, (c) S4Z0, (d) S2Z2.
Figure 4 displays the TEM image of Al6(Mn,Fe) phase in different alloys. The size and number density of phases were quantified by image analysis on TEM images. In alloy S0Z0, Al6(Mn, Fe) phase was large in size and the number density of it was 2.4 × 1016 m−3. When only 0.2%Sc was added, the size of the Al6(Mn,Fe) phase became smaller and the number density improve to 4.5 × 1016 m−3. As the Sc content was increased to 0.4%, the size further decreased and the number density further increased to 7.2 × 1016 m−3. The simultaneous addition of Sc and Zr increased the effect, the Al6(Mn,Fe) phase in S2Z2 alloy was smaller and more numerous than that in S4Z0 alloy where the number was 11.1 × 1016 m−3. These results are all due to the grain refinement achieved with Sc and Zr elements. S4Z0 and S2Z2 had grains refined due to the presence of primary Al3Sc and Al3(Sc,Zr), which increased the nucleation possibility of Al6(Mn,Fe) phase and promoted the finer dispersion of Al6(Mn,Fe) phase in the alloy. However, the Sc in S2Z0 alloy was dissolved in the matrix; hence, the primary Al3Sc phase cannot precipitate. Only the secondary Al3Sc phase precipitated after heat treatment, which cannot refine the grain size.
The mechanical properties of the rolling annealed state of Al–5Mg–0.7Mn alloys with different Sc and Zr contents are presented in Table 2. As shown in Table 2, the addition of Sc singely or Sc and Zr simultaneously could both significantly improve the mechanical properties. When there are Sc addition, with the increase in Sc levels (from 0% to 0.4%), the yield strength of experimental alloys have significantly improvement (from 186.7 MPa to 351.8 MPa) and ultimate tensile strength also improved (from 361.5 MPa to 432.5 MPa) but elongation significant decrease (from 20.9% to 12.2%). The vicker hardness of Al–5Mg–0.7Mn improved from 100.2 HV to 126.7 HV. After Sc and Zr were added simultaneously, S2Z2 alloy showed higher YS, UTS, and hardness than S4Z0, improving respectively 17.1 MPa, 20.1 MPa, and 10.8 HV when compared with S4Z0. As compared with S4Z0, the elongation of S2Z2 alloy also increased from 12.2% to 15.2%. These results suggest that the addition of Sc and Zr simultaneously could increase the strength and hardness of the alloy while sacrificing less on the deformation ability. The strengthening effect of Sc improved when incorporated with Zr. In order to analyse the property improvement by Sc and Zr, the fractography of experimental alloys with different levels of Sc and Zr additions were also studied.
The factography of tensile samples can contribute to understand the effects of Sc and Zr on experimental alloys. In Fig. 5, there are large quantities of ridges and dimples in all alloys, this shows that all alloys have great plasticity. The dimples in S0Z0 were bigger in size and lower in quantity than in S2Z0 and S2Z2. At the bottom of the dimples, there often existed the Al6(Mn,Fe) phase. The small dimples and cracked intermetallics imply that the fracture of all alloys was a combination of ductile tear of Al matrix and a cleavage fracture of Al6(Mn,Fe). But the situation in the case of S4Z0 alloy (Fig. 5(c)) was different. The fractured surface of S4Z0 contained just a few dimples and it was a partial brittle fracture. These results signify that the brittle fracture mechanism became clearer in the S4Z0 alloys. The addition of only Sc could improve the mechanical properties, but when Sc content was high (0.4%), the risk of brittle fracture increased. With the simultaneous addition of Sc and Zr, the strength improved while the risk of brittle fracture decreased. These results are consistent with the previous experimental results.
SEM images of the fractures: (a) S0Z0, (b) S2Z0, (c) S4Z0, (d) S2Z2.
The polarization curves of the four alloys are shown in Fig. 6, and the corrosion voltages (Ecorr), corrosion current densities (Icorr) and pitting potential (Epit) of the four investigated alloys according to the polarization curves are exhibited in Table 3. With the increase in Sc and Zr content, the shapes of the curves were basically unchanged. As shown in Table 3, the addition of Sc and Zr could decrease Icorr in the sequence of S0Z0 > S4Z0 > S2Z2 > S2Z0, which can partially indicate the corrosion resistance of all the alloys, as in general, there is a positive correlation between the Icorr and corrosion resistance.26)
Polarization curves of the annealed Al–Mg–Mn alloys.
Figure 7 shows the degrees of intergranular corrosion (IGC). According to the maximum IGC depth and the values of the IGC depths,27) we can compare corrosion degrees of different alloys (Table 4). By the SEM, the typical IGC morphologies of all the samples were similar, but the degrees of corrosion varied. The intergranular corrosion depth of the S0Z0 was 41 µm, the whole specimen showed superficial IGC (Fig. 7(a)). With Sc content increased to 0.2%, the IGC depth decreased (23 µm) and the corrosion degree became lower (Fig. 7(b)). Alloy S4Z0 with 0.4% Sc displayed an IGC depth of 23 µm (Fig. 7(c)). When 0.2%Sc and 0.2%Zr were added, the IGC depth of alloy S2Z2 was 10 µm (Fig. 7(d)). The experimental results show that the IGC susceptibility of the four alloys were S0Z0 > S2Z0 > S4Z0 > S2Z2. Sc could improve the corrosion resistance of the experimented alloys. With the increase in Sc content, the IGC corrosion resistance of the experimented alloys became higher, which may relate to the size, shape, and quantity of the Al6(Mn,Fe) precipitates on the grain boundaries.28) As shown on the surfaces presented in Fig. 7(a), the worst IGC corrosion was for S0Z0, as there were several deep corrosion pits which connected into the lines in grain boundaries. Shallower corrosion lines occurred on the original surface of the S2Z0 and S4Z0 alloys, showing higher corrosion resistance than that for S0Z0. The shallowest corrosion line on the surfaces for S2Z2 proves that S2Z2 had the highest corrosion resistance among the four alloys. Combining the results of microstructure and electrochemical texts, we can conclude that the stronger the grain refinement effect, the finer was the distribution of Al6 (Mn, Fe) phase, which improved the corrosion resistance of experimented alloys.
OM of the IGC depth and SEM of the IGC surfaces: (a) S0Z0, (b) S2Z0, (c) S4Z0, (d) S2Z2.
The YS and UTS of the alloys were in the sequence of S2Z2 > S4Z0 > S2Z0 > S0Z0, while the elongation had a sequence of S0Z0 > S2Z0 > S2Z2 > S4Z0. According to the microstructure analysis of the alloy, the strengthening mechanism of Sc and Zr in Al–Mg–Mn alloy was fine grain strengthening, and the main role of refining grains is primary Al3Sc and Al3 (Sc, Zr) phase. They could refine grains as a non-uniform nucleation core, improving the mechanical properties of the alloy. Although the secondary Al3Sc phase could hinder the movement of dislocations and subgrain boundary and eliminate the dendritic structure of the alloy, the grain refining effect was not clear and the strengthening effect was limited. Simultaneously, due to the grain refinement, the Al6 (Mn, Fe) phase was more finely precipitated and dispersed and also had some precipitation strengthening effect.29)
The Al–5Mg–0.7Mn alloy showed decent mechanical and anticorrosion properties and the addition of Sc and Zr could significantly improve the mechanical properties and corrosion resistance of the alloy.
One of the reasons for this phenomenon is that the Sc and Zr could refine the grain of the alloy.30) Different contents of Sc and Zr had different grain refinement effects. Only Al3Sc and Al3(Sc,Zr) precipitated as an inhomogeneous nucleation core in the as-cast state could refine grain significantly. As only 0.2% of Sc was added in the S2Z0 alloy, due to its low content, it could only be dissolved in the matrix; hence, it was impossible to precipitate the primary Al3Sc phase which could refine grains. Sc solid solution in the matrix caused the change in structure of the grain from dendrites to equiaxed crystals along with having a certain degree of strengthening effect. When the Sc content was increased to 0.4%, the primary Al3Sc phase began to precipitate as a non-uniform nucleation core. It could significantly refine grains, and in the process, improving the mechanical properties of the alloy. The addition of Zr to the alloy of low Sc content could promote the precipitation of Al3 (Sc, Zr). Al3 (Sc, Zr) phase could also be the non-uniform nucleation core; therefore, the grain size of S2Z2 was similar to that for S4Z0.
According to the SEM and TEM images of the alloy, the most important second phase in the alloy was Al6 (Mn, Fe), which also had a great influence on the properties of the alloy. Sc and Zr could improve the mechanical properties of the alloy by modifying the morphology, size, and distribution of Al6(Mn,Fe). In the base alloy, the Al6(Mn,Fe) phases had coarse rod-like morphology, which greatly increased the mechanical properties of the alloy. When the content of Sc was low (0.2%), Sc was dissolved in the matrix. Although there was no clear grain refining effect, the morphology of the Al6(Mn,Fe) phase became more regular, which improved the mechanical properties. When the Sc content was increased to 0.4%, the grains were refined due to the presence of primary Al3Sc, resulting in fine and uniform distribution of Al6(Mn,Fe) phases on the grain boundaries. In the case of S2Z2 alloy, Al3 (Sc, Zr) phase was precipitated, whose refining effect was stronger than that for S4Z0; hence, the mechanical properties were better than that of S4Z0. The mechanical properties of the four alloys were in the sequence of S2Z2 > S4Z0 > S2Z0 > S0Z0.
As per the corrosion morphologies, the localized corrosion progress started from the eutectic areas. As per the previous studies,31) it is generally believed that the corrosion of Al–Mg alloy begins with the β-Al3Mg2 phase which easily precipitates in the grain boundaries. However, when Mg content was below 5%, the β-Al3Mg2 phase was difficult to precipitate. After Mn was added, Al6(Mn,Fe) precipitation occurred easier on the grain boundaries. Figure 8(a)–(c) shows the schematic diagram of the IGC of Al–Mg–Mn–Fe. Compared with the matrix α-Al, the Ecorr of Al6(Mn,Fe) was higher, which made Al6(Mn,Fe) into local cathodes with respect to the matrix.32) When corrosion happened, the matrix would dissolute and corrosion rings appeared around the Al6(Mn,Fe) phase. As shown in Fig. 8, after the Al6(Mn,Fe) phase fell out, some corrosion pits remained. In the S0Z0 base alloy, Al6(Mn,Fe) existed in the form of irregular and coarse second phase, increasing the localized corrosion. The addition of Sc and Zr made Al6(Mn,Fe) finer and dispersed, improving the corrosion resistance of the alloys. Meanwhile, the grain size also influenced the corrosion resistance of the Al alloys. Some studies32) have shown that the corrosion resistance of Al alloys increases as the grain size decreases because the fine-grained microstructure retains more reactive surfaces which are related to the oxide deformation. Also, the corrosion performance was better in this microstructure because of segregation improvement. Therefore, the grain refinement also helped in enhancing the corrosion resistance of Al–5Mg–0.7Mn.
Schematic diagram of the IGC of Al–Mg–Mn.
We investigated the mechanical properties, microstructure, and corrosion resistance of Al–Mg–Mn with different contents of Sc and Zr. Two strengthening particles Al6(Mn,Fe) and Al3Sc were researched. The following are the primary conclusions:
This work was supported by the 2015 ShanDong province project of outstanding subject talent group, the Natural Science Foundation of Shandong Province of China (ZR2017MEM005), the project (2017GK2120) supported by the Key Research and Development Program of Hunan Province and the Natural Science Foundation of Hunan Province of China (2018JJ2506).
