Study of Intergranular Corrosion Behaviors of Mn-Increased 5083 Al Alloy with Controlled Precipitation States of Al₆Mn Formed during Homogenization Annealing

Abstract

In this study, as a vital part of the production of Mn-increased 5083 Al alloy, i.e., homogenization annealing before hot rolling, the target states of key Al₆Mn precipitation, including the dispersed, initial coarsening and intensive coarsening states, were designed, and the corresponding precipitates formed via the control of the temperature and holding time in the annealing process. By means of metallographic corrosion and nitric acid mass loss tests (NAMLT) for assessing the intergranular corrosion (IGC) resistance, temperatures ranging from 175 °C to 225 °C were determined to induce a transition from sensitization to stabilization for this innovative 5083. At a temperature of 175 °C for a duration of up to 24 h (2 h, 4 h, 8 h, 16 h, 24 h), the results show that when the soak time is 24 h, the sample with initially coarsened Al₆Mn phases has a lower degree of sensitization (DOS) compared to the samples with Al₆Mn phases in both the dispersed and intensive coarsening states, and its NAMLT is reduced by 11% and 15%, respectively. Subsequently, transmission electron microscopy (TEM) analysis has investigated that for the sample with the best IGC resistance, i.e., that with initially coarsened Al₆Mn phases, plate-like Al₆Mn particles (200~500 nm) can act as heterogenous nucleation sites for β phases, driving their preferential precipitation on Al₆Mn particles and resisting their precipitation along grain boundaries, ultimately improving the IGC resistance of 5083 Al alloy after homogenization annealing.

1. Introduction

The 5xxx Al-Mg alloys are widely utilized in the automotive and shipbuilding industries due to their excellent corrosion resistance, weldability, and formability [1,2,3]. However, prolonged exposure to temperatures between 50 °C and 200 °C significantly increases their susceptibility to IGC, potentially leading to severe stress corrosion cracking (SCC) [4,5,6]. This issue arises primarily due to the precipitation of the β phase, which has an electrochemical potential of 200 mV lower than the Al matrix [7], as tested in a 3.5 wt.% NaCl solution using a saturated calomel electrode (SCE), making it more prone to preferential corrosion at the cathode. Within this temperature range of 50~200 °C, the β phase continuously precipitates along grain boundaries, exacerbating IGC [8,9]. Components such as automotive engines and ship boilers often operate within this temperature range, and extended exposure can result in SCC-induced failures. Stabilization treatment is usually applied to materials to improve the IGC; during stabilization at the specified temperature, the precipitation rate of the β phase decreases, and it preferentially nucleates at the particles along the grain boundaries, which results in the discontinuous distribution of these particles along the grain boundaries, thereby inhibiting IGC [10].

Multiple factors influence materials’ susceptibility to IGC, including their elemental composition, grain orientation, grain size, texture, and dislocations. Research by Carroll [11] and Meng et al. [12] pointed out that adding Zn to Al-Mg alloys can improve the IGC by reducing the amount of β phases. The addition of Zn leads to the formation of the phase in the matrix, i.e., Mg₃₂(Al, Zn)₄₉, known as the τ phase. Compared to the β phase, the τ phase exhibits a lower potential difference than the matrix, leading to a reduced tendency for corrosion. At sensitization temperatures, the τ phase preferentially precipitates from the matrix over the β phase. This effectively reduces the distribution of the β phase at the grain boundaries, thereby enhancing the material’s resistance to intergranular corrosion. Ding et al. [2] found that the β phase typically precipitates at high-angle grain boundaries (HAGB) as well as triple junctions, which can serve as nucleation sites. The metal processing condition also affects the IGC. Qin et al. [13] observed that in 5083 cold-rolled annealed plates, the coarse-grained samples exhibited better IGC resistance, while the fine-grained samples showed superior SCC resistance due to increased grain boundaries, leading to corrosion cracks that propagate along a tortuous path rather than propagating along a straight grain boundary, thus enhancing the alloy’s SCC resistance. Li et al. [14] discovered that the brass texture induced during deformation can effectively improve the material’s IGC behavior, as the large misorientation angles between adjacent grains cause cracks to deflect, thereby enhancing the corrosion resistance. Dislocations can accelerate the diffusion of Mg atoms via pipe diffusion; the diffusion rate of Mg atoms via pipe diffusion is several thousand times greater than their rate of normal precipitation. Consequently, samples with higher dislocation densities have displayed poorer IGC resistance [15].

The research regarding the influences of precipitation, including the type, scale and distribution, on the IGC of Al-Mg alloys is limited; existing research has been limited to the influence of the τ phase and distribution on the precipitation behavior of the β phase, and current studies have not yet identified the impact of other types of precipitates on the IGC of Al-Mg alloys. Meanwhile, this study qualitatively analyzes the sensitization and stabilization transition temperature ranges of the Mn-increased 5083 alloy, examines the precipitation state of the β phase at sensitization temperatures, and investigates the effect of a variety of Al₆Mn precipitations formed during homogenization on the precipitation behavior of the β phase. Mofarrehi et al. [16]. observed that the morphology of Al₆Mn did not change significantly after hot rolling. Therefore, to enhance the experimental precision, homogenization-annealed ingots were selected to avoid the influence of the grain orientation, grain size, texture, and dislocations induced by hot rolling, allowing for a systematic analysis of how different Al₆Mn morphologies affect the IGC performance.

2. Materials and Methods

The 5083 alloy samples used in this experiment were provided by Southwest Aluminum (Group) Co., Ltd. (Kunshan, China), and produced via direct casting. The chemical composition, measured by Spark Optical Emission Spectroscopy (Spark OES, Thermo Fisher Scientific, Waltham, MA, USA), was Al-4.78Mg-0.70Mn-0.05Fe-0.08Cr-0.01Zn-0.01Ti-0.03Si in wt.%. The ingots underwent the annealing process to achieve homogenization at three selected temperatures, i.e., 470 °C, 490 °C and 510 °C, with a holding time ranging from 2 to 24 h. During the homogenization annealing, the samples were heated from room temperature in air, with a heating rate of 30 °C/h below 400 °C and 13 °C/h above 400 °C. The samples were quenched in water upon completion of the homogenization annealing. After annealing, the post-annealed samples were prepared for subsequent TEM observation. The samples were ground and thinned to 70 μm through mechanical polishing on #800 sandpaper, and then subjected to twin-jet electropolishing in a solution of 10% HClO₄ + 90% CH₃CH₂OH at −30 °C with an applied voltage of approximately 20 V. TEM observation was conducted using a JEOL-JEM 2100 transmission electron microscope (Tokyo, Japan) at an operating voltage of 200 kV.

The ingots with a size of 50 mm × 20 mm × 6 mm were subjected to sensitization treatment at three different temperatures, namely 175 °C, 200 °C, and 225 °C, with holding times of 2, 4, 8, 12, 18, and 24 h. The sensitized samples were then etched in a 40% phosphoric acid aqueous solution for 6–8 min, and the precipitation behaviors of the β phases were observed using a metallographic microscope. The ASTM G67 NAMLT [17] test was performed to assess the material’s sensitivity to IGC, and three parallel samples were selected for each set of sensitization parameters. First, the dimensions of the materials were measured to calculate their surface area (S). Then, the materials were cleaned with a 5% NaOH (in wt.%) solution to remove any oil from the surface of the samples. After air drying, their mass was measured as M₀. The samples were then placed in a 70% HNO₃ (in wt.%) solution at a temperature of 35 °C for 24 h. After removal, the surfaces were cleaned with a brush, and after air drying, their mass was measured as M₁. Then, the DOS was calculated using the following formula: DOS = (M₀ − M₁)/S; samples with different DOS values were then selected for further TEM observation.

3. Results and Discussion

During the homogenization annealing experiments on the 5083 alloy, Al₆Mn precipitation was observed, and its shape and size varied with the changes in the homogenization parameters. Figure 1a shows the DSC (Differential Scanning Calorimetry) curve of the ingot, with peaks at 570 and 640 °C corresponding to the solidus and liquidus of the material, respectively. The DSC curve represents the change in heat flow during the exothermic process, thus reflecting the phase transition behavior in the material. The homogenization temperature range was set between 470 and 510 °C. Based on the size and morphology of the Al₆Mn phase under different homogenization parameters, the precipitates can be categorized into three types: dispersed distribution, initial coarsening, and intensive coarsening. Figure 1b displays the morphology of the Al₆Mn precipitates under different homogenization conditions.

The precipitation behavior of Al₆Mn can be analyzed based on Fick’s First Law Equation (1) and the Arrhenius Equation (2), where J is the diffusion flux, D is the diffusion rate, C is the volume concentration of the component, dc/dx is the concentration gradient, D₀ is the initial diffusion coefficient, Q is the activation energy, and R is the gas constant, at approximately 8.314 J/(mol·K). T is the absolute temperature, measured in Kelvin (K).

$$J=D{\displaystyle \frac{dc}{dx}}$$

(1)

$$D=D_0\exp (-{\displaystyle \frac{Q}{RT}})$$

(2)

Fick’s First Law states that the diffusion flux is related to the diffusion coefficient and the concentration gradient, while the Arrhenius Equation indicates that the diffusion coefficient of the material is solely dependent on the temperature (T); the higher the temperature, the greater the diffusion coefficient. Therefore, the diffusion flux can be expressed by Equation (3), and the total amount of diffusion per unit time (N) can be represented by Equation (4):

$$J=D_0\exp (-{\displaystyle \frac{Q}{RT}}){\displaystyle \frac{dc}{dx}}$$

(3)

$$N=J·\mathsf{\Delta }\mathrm{t}=D_0\exp \left(-{\displaystyle \frac{Q}{RT}}\right){\displaystyle \frac{dc}{dx}}$$

(4)

Therefore, when the temperature is constant, the diffusion rate of Mn atoms remains unchanged, with the concentration gradient being the only factor affecting their diffusion flux. During the homogenization annealing process, Mn atoms dissolved in the matrix continuously precipitate out as Al₆Mn, thereby reducing the concentration of dissolved Mn atoms in the matrix and decreasing the concentration gradient.

Figure 2 presents the TEM bright-field images used to illustrate different Al₆Mn precipitation states, i.e., initially coarsened, intensively coarsened and dispersed Al₆Mn phases. First of all, Figure 2c displays a sample with dispersed Al₆Mn particles, with numerous nano-sized Al₆Mn particles in the Al matrix, each with a length not exceeding 80 nm. Additionally, there are minor rhombic and plate-like precipitates ranging from 100 to 250 nm in length. Figure 2a illustrates the initially coarsened Al₆Mn. The dispersed Al₆Mn nanoparticles have disappeared, and the Al matrix exhibits a large amount of rhombic and plate-like precipitates. According to Figure 2d, the rhombic precipitates are measured to be up to 250 nm in length, while the plate-like precipitates range from 200 to 500 nm. Figure 2b displays intensively coarsened Al₆Mn. Here, the density of precipitates in the Al matrix has significantly decreased, with the matrix predominantly featuring rhombic and rod-like precipitates. The rhombic precipitates have grown larger, and their sizes are between 300 and 500 nm, and only a few plate-like precipitates remain. Notably, a number of large rod-like Al₆Mn precipitates appear in the matrix, exceeding 800 nm in length, with some even reaching up to 2.5 μm.

According to our previous research and relevant literature [16,18,19], in the initial stages of homogenization, Mn predominantly exists in the Al matrix as numerous, finely dispersed nanoscale Al₆Mn particles, which are small in size and abundant in number. As the homogenization continues, these Al₆Mn nanoparticles transform into small rhombic and plate-like precipitates, as observed in Figure 2c. With the reduction in the number of Al₆Mn nanoparticles, there are numerous small rhombic and plate-like Al₆Mn precipitates. As the homogenization progresses further, these small rhombic and plate-like Al₆Mn precipitates also undergo coarsening, with the plate-like precipitates exhibiting the most pronounced coarsening. The coarsened rod-like precipitates can grow up to 2.5 μm in length.

To study the corrosion resistance of the designed material, it is essential to understand its corrosion behavior and corresponding mechanisms, as well as to clarify the corresponding sensitization behavior with this composition system. In the Al-Mg alloys, the precipitation of the β phase (Al₃Mg₂) is the primary factor influencing the corrosion performance. Since the β phase is incoherent with the Al matrix, it typically precipitates along grain boundaries. In other words, as the β phases accumulate, they continuously precipitate along the grain boundaries. The potential of the β phase is 200 mV lower than that of the matrix, leading to its preferential corrosion before the matrix, leading to corrosion along the grain boundaries, known as IGC. IGC occurs only when the Mg content exceeds 3%, and in 5xxx alloys, the temperature sensitization range is between 50 °C and 220 °C. The sensitization temperature is related to the content of Mg; in other words, increasing the Mg raises the corresponding sensitization temperature. The Mg content added to the samples is 4.78%;

Table 1. Parameters and sensitization conditions of Al-Mg alloys in different research.

Number	Material and State	Mg Content (wt.%)	Sensitization Conditions
1 [20]	5083 H116	4.4	150 °C–1250 h
2 [21]	5028 H116	3.2–4.8	120 °C–7 day
3 [22]	5083 H	4.2	150 °C–7 day
4 [23]	5456 H116	5.0	100 °C–14 day
5 [24]	5083 H112	4.0	175 °C–100 h
6 [25]	5083 H321	4.8	175 °C–100 h
7 [26]	5182 T6	4.5	150 °C–100 h

lists the sensitization temperatures obtained from the literature for similar compositions.

The practical stabilization temperature of the samples generally differs from the theoretical sensitization temperature by 20~50 °C. When this transition range is reached, the precipitation rate of the β phase slows down until it significantly decreases at the stabilization temperature. In this study, the corrosion behavior was studied within the range of 175~225 °C, and the initial task was to determine the corrosion behavior of the material within this temperature range. Therefore, samples with a dispersed distribution of the Al₆Mn phase were selected for testing in the sensitization and stabilization intervals. Ding et al. [2] conducted NAMLT experiments on Al-Mg alloys and found that during sensitization with a varying Mg content, the precipitation rate of the β phase was significantly higher in the first 24 h compared to later stages. As the holding time extended, the β phase exhibited a network distribution, making it difficult to distinguish under metallography. Therefore, the homogenization time for this study was set to 2~24 h. After sensitization, the samples were etched using a phosphoric acid solution, and the metallographic results are shown in Figure 2.

Figure 3 presents the metallographic images of samples held at 175 °C, 200 °C, and 225 °C for 4, 8, 16, and 24 h. The samples from the first 4 h are omitted due to small differences. As shown in Figure 3a–c, at 175 °C for 8 h, the β phase precipitates are already noticeable at the grain boundaries. Extending the holding time to 16 h results in β phase precipitation not only along the grain boundaries, but also the formation of a network pattern, which becomes fully apparent at 24 h. At 200 °C, the β phase distribution at the grain boundaries is less distinct compared to at 175 °C. However, a network distribution pattern is still observed in samples held for 24 h. At 225 °C, there is no significant change in β phase precipitation, which is limited to a few grain boundaries, and no network distribution is detected.

This indicates that among the three temperature sets, β phase precipitation is most intense in samples held at 175 °C. As the temperature increases, the precipitation rate of the β phase significantly decreases. At 225 °C, β phase precipitation is negligible even with extended holding times, suggesting that this temperature corresponds to the stabilization temperature of the material. In contrast, 175 °C and 200 °C correspond to the sensitization temperature and transition temperature, respectively. Figure 4 presents the NAMLT results for the tested samples at the three temperature sets, showing a correlation between the DOS and temperature. At 175 °C, the samples exhibit the highest DOS, reaching 18.3 mg/cm² after 24 h. At 200 °C and 225 °C after being held for 24 h, the DOSs are 8.4 mg/cm² and 3 mg/cm², respectively. A considerably low DOS at 225 °C (well below 25 mg/cm²) and its stability over time indicate that this temperature corresponds to the stabilization temperature of the samples. It was also noted that at 175 °C, the slope of the DOS curve gradually decreases. According to Fick’s First Law [17], this reduction in intergranular corrosion weight loss is mainly related to the decreased concentration of Mg atoms dissolved in the matrix, and the reduction in the concentration gradient inhibits the precipitation of the β phase.

Figure 5 presents the precipitation behavior of the β phase. Here, Figure 5a shows the sample held at a stabilization temperature of 225 °C for 12 h, where almost no β phase precipitation was observed along the grain boundaries. In contrast, Figure 5b displays the sample held at a sensitization temperature of 175 °C for 24 h, where the β phase is continuously distributed along the grain boundaries. Figure 5d,e illustrate the distribution behavior of the β phase at the triple junctions in a sample that was held at a transition temperature of 200 °C for 24 h. In this case, the β phase is not only continuously distributed along the grain boundaries, but it also exhibits segregation at the triple junctions. This indicates that the β phase exhibits different precipitation behaviors as the temperature increases. At the sensitization temperature, the β phase precipitates in the largest quantity and is continuously distributed along the grain boundaries. As the temperature increases to the transition temperature, the β phase begins to shift from grain boundary distribution to segregation at the triple junctions. There are two reasons for this segregation: firstly, the higher energy of the triple junctions reduces the formation of an energy barrier for the β phase, making it easier to precipitate out in these regions; secondly, the initially precipitated β phase can serve as a heterogeneous nucleation site for the subsequent β phase, leading to further segregation in these areas. In the samples held at the stabilization temperature, the precipitation rate of the β phase is too low, so no β phase precipitation was observed along the grain boundaries. Consequently, the samples that were held at the sensitization temperature exhibit poor IGC. As the temperature increases to the transition temperature, the IGC resistance improves, and the samples at the transition temperature for 24 h show a 57.3% reduction in IGC compared to those at the sensitization temperature. The samples held at the stabilization temperature exhibit the best IGC, with a DOS of 3.7 mg/cm², representing an 85.2% reduction compared to the sensitization temperature.

To investigate the effect of Al₆Mn precipitated during the homogenization on the intergranular corrosion resistance of the material, NAMLT experiments were conducted at the sensitization temperature of 175 °C. The experimental results are shown in Figure 6. It was found that the samples with initially coarsened Al₆Mn exhibited the lowest DOS, with a value of 16.6 mg/cm² after 24 h of holding time. This was followed by the samples with dispersed Al₆Mn, which had a DOS of 18.3 mg/cm² after 24 h. The samples with severely coarsened Al₆Mn showed the highest DOS, with an increase of 15% compared to the initially coarsened samples. The NAMLT experimental results indicate that the amount of β phase precipitation is undoubtedly the lowest in the samples with initially coarsened Al₆Mn, demonstrating the best intergranular corrosion resistance.

Figure 7 presents bright-field images and EDS results of Al₆Mn under TEM, revealing the relationship between β and plate-like Al₆Mn. Figure 7a–c show plate-like Al₆Mn, with lengths ranging from 300 to 500 nm. In Figure 7a,c, white precipitates are observed at the ends and middle of the plate-like Al₆Mn indicated by red arrows. The EDS results indicate that these precipitates contain Mg. Ding et al. [2] studied the 5E83 alloy, and suggests that these white precipitates at the ends are β phase. In this region, the β phase precipitates exist in agglomerated form, attached to the surface of the plate-like Al₆Mn. However, it is noteworthy that not all plate-like Al₆Mn exhibit agglomeration with β phase, as seen in the sample shown in Figure 7b, where no β phase precipitation was observed. Figure 7d shows the phases of rod-like Al₆Mn, the length reached 1.3 μm, but no precipitates were observed on its surface.

Figure 8 shows the bright-field micrographs and EDS results of two rhombic Al₆Mn phases with different sizes. Figure 8a presents Al₆Mn that has undergone initial coarsening and dispersed distribution, with a smaller size of 150 nm. Figure 8b shows the coarsened rhombohedral phase in a sample where Al₆Mn has undergone significant coarsening, reaching a length of 300~400 nm. No precipitates were observed on the rhombic phases, and the EDS results indicate that, apart from a small amount of Fe [27,28], no other elements are present on the rhombic Al₆Mn. This indicates that the rhombic Al₆Mn does not influence the distribution of the β phase when held at 175 °C, thereby ruling out its impact on the sample’s corrosion resistance.

From the analysis of the effects of different Al₆Mn phases on the distribution of the β phase in Figure 7 and Figure 8, it can be seen that the β phases exhibited a preferential distribution only on certain plate-like Al₆Mn precipitates. The β phases were also observed on the fine and coarsened rhombohedral and rod-like phases. However, no β phases were observed on the dispersed, particle-like Al₆Mn, which is typically spherical. Due to its smaller surface area, i.e., spheres usually exhibit the lowest surface energy, it was not expected to provide non-uniform nucleation sites for the precipitation of the β phase. Therefore, the TEM results suggest that only plate-like precipitates can provide non-uniform nucleation sites for the β phase, allowing the β phase to exist within the crystal structure without being coherent with the matrix. Assuming a fixed rate of β phase precipitation, the presence of plate-like Al₆Mn reduces the amount of β phase precipitation at the grain boundaries, thereby improving the material’s resistance to intergranular corrosion.

This suggests that optimizing the homogenization annealing process can effectively improve the precipitation behavior of the β phase at the sensitization temperature. As shown in Figure 1b, which illustrates the relationship between Al₆Mn and the homogenization temperature and holding time, to achieve the formation of plate-like Al₆Mn, the holding time should be 18–24 h at 470 °C, 12–18 h at 490 °C, and no more than 4 h at 510 °C. By adjusting the homogenization process to control the precipitation state of Al₆Mn, plate-like Al₆Mn can be obtained as non-uniform nucleation sites for the β phase. This allows the β phase, which is originally incoherent with the Al matrix, to concentrate on the plate-like Al₆Mn within the matrix, reducing its aggregation at grain boundaries and thereby enhancing the material’s resistance to IGC.

4. Conclusions

The homogenized samples of 5083 alloy were heated and maintained at different temperatures, and the precipitation behaviors of the β phases were observed using OM and TEM. Subsequently, at the determined sensitization temperatures, different states of Al₆Mn were subjected to NAMLT testing to quantitatively analyze the differences in β phase precipitation. The influence of Al₆Mn on β phase precipitation was then observed through TEM. The main conclusions are as follows:

(1)

Sensitization temperature determination: Through metallographic etching, TEM observation, and NAMLT testing, it was determined that 175 °C is the sensitization temperature for the material, 200 °C is the transition temperature, and 225 °C is the stabilization temperature.

(2)

The changes in β phase precipitation behavior: As the temperature increases from sensitization to transition to stabilization, the precipitation behavior of the β phase changes accordingly. At the sensitization temperature, the β phase is continuously distributed along the grain boundaries; at the transition temperature, this continuous distribution weakens, with some β phase forming at the triple grain junctions; at the stabilization temperature, the β phase is almost absent at the grain boundaries.

(3)

The effect of Al₆Mn on the β phase: the TEM observations of sensitized samples containing different Al₆Mn phases showed that the β phase preferentially precipitates on plate-like Al₆Mn, leading to intragranular precipitation, which improves the material’s resistance to intergranular corrosion. Further, the NAMLT testing indicated that plate-like Al₆Mn can reduce the degree of sensitization by 15%.