1 Introduction

Al–Mg–Si alloys present a high industrial interest since they can be age-hardened, offering improvements in their mechanical strength and a broad spectrum for designing tailor-made production processes according to the desired application, for instance, electrical conductors [1, 2]. The typical microstructure of these alloys is composed of $\alpha ‐\text{Al}$ solid solution with Mg₂Si particles as a dispersed phase. Malekan et al. [3] calculated this system's pseudo-eutectic transformation point and found a range between 583.5 and 594 °C at approximately 13.9% Mg₂Si. The reinforcing phase of Mg₂Si presents a maximum solubility in aluminum of 1.85%, high thermal stability, and melting temperature since Mg–Si bonding energy is quite high. Moreover, its stoichiometric relation Mg: Si is close to 1.73:1 [1].

Al–Mg–Si alloys offer the possibility of precipitation hardening through natural and artificial aging, which sets them in a position against other aluminum alloys. Both the composition of the alloy and the thermal history directly influence the precipitate's final microstructure in terms of size, type, quantity, and distribution of precipitates. 6xxx grade aluminum alloys are regularly subjected to a sequential process of homogenization, extrusion, high cooling rates, and storage at room temperature to end up with artificial aging [4, 5]. The mechanical strength of these aluminum-based alloys is related to the interfacial deformation between the matrix and tiny precipitates, which are generally coherent or semicoherent and metastable concerning the lattice [4]. However, aging comprises some transitions, advancing from a metastable to a stable condition. A general overview of the possible sequential or even simultaneous precipitations is presented in [6]. Nevertheless, Al–Mg–Si alloys' thermomechanical treatments and chemical composition should be carefully controlled to reach optimum mechanical strength and conductivity [2]. The thermomechanical treatment conditions and the chemical composition of Al–Si–Mg alloys must be carefully controlled to obtain the microstructure that guarantees the mechanical properties and electrical conductivity required for a particular application, such as manufacturing umbilical cables [7].

After direct extrusion, a meticulously controlled high cooling rate is crucial. It produces a supersaturated solution of Mg and Si and a high vacancy concentration. For some 6xxx series alloys that have undergone a hot area reduction process by direct extrusion at 450 °C with temperature control in the die, in the outer zone of the extrusion chamber and an inner zone of the extrusion chamber, followed by high-speed cooling, an incubation or natural aging time is necessary before any artificial aging treatment, to guarantee in the decomposing supersaturated solid solution the presence of nucleation sites for the hardening phases, that will precipitate later during the artificial aging treatment. These pretreatments condition the microstructure and properties of the final alloy [4].

Alloying elements play a pivotal role in shaping the characteristics of the final material. For instance, elements such as Mn or Cr could catch Si and Mg of the main composition and induce a higher sensitivity during the cooling process. Other materials classified as impurities and remain after the production process by electrolysis of pure aluminum maintain their quantities at low concentrations, as is the case of Fe [1]. Regarding the main components, the Si content has a more significant effect on the material, although it does not considerably affect the formation of Mg₂Si; in excess, it favors the precipitation of hardening phases during aging much more than Mg or Mg₂Si [1].

Furthermore, alloy composition can affect the supersaturation level and the time it takes for optimal precipitation up to maximum hardness conditions. For low solute concentration, the supersaturation level tends to be down, and the energy level, compared to the equilibrium state, is not entirely different, affecting the precipitation density. Conversely, high solute concentration, without exceeding dissolution limits presented in the phase diagram, induces faster nucleation and growth rate of precipitates, decreasing aging time for hardness peak to appear [8].

On the other hand, according to Reed-Hill [8] and Teichmann et al. [9], there is a direct relationship between aging temperature and the obtained hardness and mechanical strength after this heat treatment. The temperature must be correctly selected so precipitation and diffusion rates are high enough to obtain the highest mechanical strength. In this sense, the temperature should not be so low that solid-state diffusion can hardly advance. Still, not so high and so close to the solvus point that alloy supersaturation and the free energy decrease associated with the precipitation process are too low, and nucleation site density is insignificant [8]. Some researchers have reported the influence of Si in excess on both natural and artificial aging processes of Al–Mg–Si systems, as well as pretreatment and some chemical variations [4, 10–13]. In the case of pre-ageing or pre-treatment, this depletes many excess solutes and vacancies in the alloy, affecting subsequent dynamic aging [14].

In recent studies, the addition of rare earth has been considered an effective microstructure modifier to improve the mechanical properties of aluminum alloys. Primary research with these materials has focused on Al–Si since rare earth elements act as nucleation mechanism modifiers, inducing changes in the form of primary silicon crystals and α-Al matrix. Moreover, it has been found that the simultaneous addition of rare earth elements is more effective than a singular introduction in terms of grain modification to enhance alloy mechanical properties [15, 16]. Moreover, rare earths could induce changes in the precipitation sequence of these alloys. Minimal addition of such materials could change failure mechanisms from transgranular into a mixture of transgranular-intergranular failure after aging, and they could also influence the aging process [17]. In addition, 60–100 °C pre-aging treatments have been recently reported for Mg and Si-rich Al–Mg–Si alloys. In either case, it has been observed that such heat treatment leads to the formation of Mg–Si-vacancies with an Mg/Si rate close to 1, a proportion comparable to elongated precipitates formed during artificial aging. Hence, effective nucleation sites that are more stable against dissolution during artificial aging are obtained, with radii higher than critical values [10–12, 18].

The electrical conductivity of metallic materials can be affected by vacancies, grain boundaries, atomic impurities, elements in solid solutions, dispersoids and precipitates [19, 20]. For 6xxx series aluminum alloys, hardness and conductivity are a function of tempering temperature and are affected by the size and amount of precipitate β(Mg₂Si) that may form [21]. The efforts in the modification and control of the chemical composition and the manufacturing process of aluminum alloys of the 6xxx series are aimed at achieving a positive synergy between mechanical resistance and electrical conductivity; ultimately, it is sought that the most significant amount of the elements present in the alloy is part of a precipitate with a shape and size that does not generate so much dispersion of the conduction electrons [22–24], and that the mechanical properties of the alloy are a consequence of precipitation hardening, grain refinement hardening and strain hardening.

The use of grade 6201 aluminum alloys to produce umbilical cables has sparked significant interest. These alloys, with their inherent aging capacity, offer a promising platform for thermo-mechanical treatments and chemical modifications. While these modifications and production processes are still in the early stages [2], the potential is immense. For instance, the proposal by Karabay [25] to add aluminum-bore master compounds for improving electrical conductivity without inducing a significant decrease in mechanical strength is a beacon of hope for the future of this research.

Based on microstructural and mechanical analyses, this paper considers the influence of the simultaneous addition of rare earth elements and pre-treatments before artificial aging for grade 6201 aluminum alloys. The aging behavior of this alloy has been studied with the standard chemical composition, so this material is currently being industrially processed and artificially aged to produce electrical conductors. Still, mechanical strength and electrical conductivity enhancements are demanded [26]. It is essential to indicate that this article studies the effect of aging of Al–Mg–Si alloys to which a mixture of rare earths has been added on the microstructure and hardness.

2 Materials and methods

For analyzing the influence of rare earth elements (lanthanum and cerium) in addition to an AA6201 on its aging behavior, controlled casting procedures were conducted in an induction furnace of the Materials Engineering School of Universidad del Valle. The chemical composition of raw materials used for either casting process is presented in Table 1. Two chemical variations were considered: combining two rare earths between 0.3% and 0.5% (A-CRE) and a Non-Variated 6201 (A-N). For both cases, a decrease in the presence of iron was demanded. The composition of the analyzed alloys is presented in Table 2. The range of variation in the composition of rare earth elements was determined by reference to the effect of these elements on aluminum alloys reported by different authors [27–30].

Samples were cut into rectangular prisms of 3 × 3 × 1 cm. Figure 1 shows all the heat treatments followed during the current study. The samples were initially homogenized at 560 °C for 5 h. The samples were quenched in water to analyze the influence of the alloy's initial state before aging in the final obtained hardness value; three different pre-treatments were done:

• One day at room temperature.

• Seven days at room temperature.

• One day at room temperature followed by storage between 60 °C and 100 °C for 16 h.

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Overview of the considered thermal treatments

Citation: International Review of Applied Sciences and Engineering 16, 1; 10.1556/1848.2024.00744

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Three random samples were considered for each condition. After each pre-aging situation, samples were artificially aged at 175 °C for up to 8 h. The results of these heat treatments were particularly significant, as they provided crucial insights into the hardness evolution of the alloy.

Our study was meticulously designed to investigate the material behavior after a pre-treatment and the addition of rare earth directly after water quenching. Three Non-Variated (A-N) samples were stored at room temperature, and three more samples containing rare earths (A-CRE) were pretreated at a temperature between 60 °C and 100 °C and then held at room temperature. Each sample group's hardness evolution was considered for up to 90 days. We employed a generalized, randomized, complete block experimental design with repetition to analyze the influence of both chemical variations by adding rare earth elements to the artificial aging behavior of an AA6201 and the different heat treatments mentioned above. Figure 1 and Table 3 provide a comprehensive overview of the different assumed conditions for this research.

Optical Spectrophotometry Analysis was used both during and after casting, with a Spectro Lab Spectrometer with a Hybrid Optic of controlled Si, Mg, and Fe [31]. Chemical analysis was performed using calibration standards for each element under ASTM E1251. Control over the aging precipitation process of the different alloys was done utilizing micro-hardness tests, according to ASTM E92-17, using an Advance Instruments TH-717 Vickers Hardness Tester. Experimental load and time were set to 0.1 kg and 15 s, respectively, and each specimen's mean value of 10 indentations for each sample was considered.

Scanning Electron Microscopy (SEM) was conducted with a JEOL JSM-6490LV, which is equipped with a model 7573 INCAPentaFETx3 detector from Oxford Instruments, owned by the Materials Engineering School of Universidad del Valle, by using 20 keV and using back-scattered electrons for EDS-Analysis, since a better contrast between zones with different chemical composition could be obtained. This test was done to inspect possible microstructural modifications, which could be induced according to changes in precipitation processes due to chemical composition variations and the different thermal history according to the aging behavior analysis.

3 Results and discussions

3.1 Pre-treatment and artificial aging

The temperature for artificial aging, a crucial factor in our study, was meticulously chosen to be 175 °C. This decision was based on a comprehensive review of numerous previous reports, which consistently indicated that the maximum hardness of aluminum alloys is attained at aging temperatures between 160 and 185 °C, with durations spanning from five to eight hours [1, 4, 11]. Importantly, this temperature is significantly distant from the solvus point of the base aluminum alloy, a fact that is further reinforced when silicon in excess is incorporated [32]. The outcomes of these distinct treatments are detailed in the following sections.

Figure 2 vividly illustrates the outcomes of the incubation period and a low-temperature pre-treatment prior to the aging process for samples containing (A-CRE) and not containing (A-N) rare earth. Regardless of the alloy under consideration, a pronounced influence of the variation in the initial incubation period and the pre-treatment they underwent was discernible. This significant impact will be further explored in a dedicated discussion, with a comprehensive comparison of the results.

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Aging behavior of A-N (left) and A-CRE (right) after 1 (1D) and seven days (7D) of incubation time and after a pre-treatment at a temperature between 60 °C and 100 °C for 16h (P)

Citation: International Review of Applied Sciences and Engineering 16, 1; 10.1556/1848.2024.00744

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Central values of the two measured samples are presented for better observation. The green point (pointed with one arrow) and purple point (pointed with two arrows) represent the hardness value in each case, which was measured after three (3) days and nine (9) days of natural aging, respectively. First, N-samples (Fig. 2–Left) showed an increase in hardness as long as the artificial aging (AG) time advanced in both analyzed cases when the incubation period was changed. In this sense, if a different hardness measurement after three days of just solution annealing and water quenching is considered as the initial reference value (green point indicated with an arrow), several increases are observed from the non-artificial state, up to 3 h of heat treatment, being even more noticeable for A-N-7D samples (27% for A-N- 1D and 35% increase for A-N-7D). After that, even though A-N-7D showed higher values at any subsequent aging time, at the final control point of 8h, the hardness value is practically the same for either incubation time, i.e., 88HV. For A-CRE Samples (Fig. 2–Right), it can be seen that there is an even more considerable change in the observed aging behavior, only due to the initial room temperature storage, in comparison to N samples. In this sense, the peak of maximum hardness shifted to later values of time, and a light hardness increase was observed for A-CRE-7D samples instead of A-CRE-1D. For the latter samples, the peak was possibly reached before or close to 3 h of AG, limited to a value between (97.3 ± 3.1) and (101.9 ± 5.4) HV. After that, averaging was the typical state, decreasing approximately 26.3% at six hours.

Conversely, for A-CRE-7D, hardness values raised to (106.2 ± 4.6) and (105.4 ± 4.3) HV; however, after six hours of AG, a decrease close to 30 % was observed at 7 and 8 h. Suppose the hardness value after three (green point indicated with one arrow) and nine days (purple point denoted with two arrows) of natural aging are considered an initial state for A-CRE- 1D and A-CRE-7D, respectively, an initial increase. In contrast with the above-mentioned treatments, pre-aging between 60 and 100 °C for 16 h changed how this alloy reacted against an ulterior aging heat treatment. As a result, a peak maximum is reached for both sample types (A-N and A-CRE). Nevertheless, it appeared just after three hours of artificial aging, where an experimental maximum for A-N-P placed between (101.4 ± 5.8) and (102.1 ± 3.5) HV and for A-CRE-P between (111.1 ± 2.8) and (105.4 ± 10.7) HV. After this time, A-N-P presented a marked decrease of approx. 23.7 %, after six hours, A-CRE-P kept high values for the same time, showing values close to the maximum of A-CRE-7D and following a similar averaging path. In discussing these results, three factors should be considered to explain the observed behavior: the influence of Si excess, thermal history, and the addition of rare earth.

Regarding the above factors, during the natural aging of Al–Mg–Si alloys, the clusters formed approximately acquire the alloy composition in terms of the Mg/Si ratio due to the low diffusion rate at room temperature. Therefore, the closer the composition is to the 1/1 ratio, the closer to the 5/6 ratio, and the β” precipitate will be formed later, leading to a better response after artificial aging [33, 34]. In the present case, such a condition could have been reached, which could explain why no negative effect of excess Si is observed; the silicon held in excess traps the iron present in the alloy to form Al–Fe–Si type precipitates, in addition to reacting with the rare earth to form Al–Si-Rare earth precipitates. All of these precipitates steal Si that should be available for aging, decreasing the probability of formation or the expected density of the precipitates. As a result, the excess Si acts as a counteracting effect to this possible decrease in Si availability. Also, the alloy's final Mg/Si ratio could have been closer to unity. The above may explain the behavior observed in Fig. 2 for the alloy under study with and without the addition of rare earth.

3.4 Natural aging

Figure 3 shows that pre-aging treatment affects the hardening of A-CRE samples at room temperature. A-CRE-N samples presented a hardness increase over time, with a tendency slope of approximately 0.088, which could be measured, starting at 64.88 ± 2.17 HV after three days and rising to 74.11 ± 1.95 HV after 81 days. Conversely, pre-treated samples (A-CRE-PN) showed a different behavior, where a stabilization tendency and no significant hardness change within 79 days of observation were obtained. On the other hand, A-N-PN samples did not present any stabilization. In contrast, they show a faster increasing tendency over time, with a growing slope of approx. 0.135 after 79 days of study. Indeed, the highest values were obtained with this sample.

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Hardness value evolution of samples with rare earths both non-pre-treated (A-CRE-N, left) and pre-treated (A-CRE-PN, center), and non-variated and pre-treated (A-N-PN, right)

Citation: International Review of Applied Sciences and Engineering 16, 1; 10.1556/1848.2024.00744

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3.6 Scanning electron microscopy

To inspect microstructural changes, A-N-1D and A-CRE-1D samples were observed after six hours of artificial aging at 175 °C (see Fig. 4). For clarification, A-CRE-1D samples were analyzed with SEM-EDS; Tables 4 and 5 present the results of the punctual composition of the EDS measurements in Fig. 5. In this sample, three behaviors of the added elements were then observed. The first is related to the precipitation of iron as intermetallic compounds at the grain boundary, as seen in measurements 1, 2, and 3 of Fig. 5.

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Micrographs of A-N-1D (left) and A-CRE-1D (right) after six hours of artificial aging. Blue and red arrows point to two types of precipitates: rounder and coarser

Citation: International Review of Applied Sciences and Engineering 16, 1; 10.1556/1848.2024.00744

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Table 4.

Semi-quantitative chemical composition results with EDS measurements presented in Fig. 5a

Number	Measure	Element (%)
		O	Mg	Al	Si	Fe	La	Ce
1	ASF			66.02	6.07	27.90
2	ASFR			38.79	13.06	1.92	14.18	32.06
3	ASR	1.56	0.54	54.91	5.19		11.55	26.25
4	ASF			77.16	5.68	17.16
5	MAT	3.24	0.67	95.54	0.55
6	MAT	3.34	0.70	95.96

Table 5.

Semi-quantitative chemical composition results with EDS measurements presented in Fig. 5b

Number	Measure	Element (%)
		O	Mg	Al	Si	Fe	La	Ce
1	ASR			33.28	13.41		15.85	37.46
2	ASFR			47.72	11.62	4.83	9.65	26.18
3	ASF			80.59	4.47	14.94
4	ASR			55.80	9.74		10.82	23.64
5	ASR	1.50		61.07	5.23		8.80	23.39
6	ASR	1.79		52.47	8.69		11.58	25.47

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Microstructure and EDS punctual measures of AA6201 alloy with rare earth elements, after one day of incubation period a) and six hours of the artificial aging process at 175 °C (A-CRE-1D) b)

Citation: International Review of Applied Sciences and Engineering 16, 1; 10.1556/1848.2024.00744

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As a result, it could be inferred that three types of precipitates are present: one containing Al–Si–Fe (ASF), another one with Al–Si-rare earth (ASR), and a third one with the whole mixture of Al–Si–Fe-rare earth (ASFR). Considering what was reported by different authors [36] Fe-bearing precipitates are commonly observed with the proposed chemical composition as Al–Fe–Si phases and in a needle-like (plate) disposition (such as punctual analysis #1 and #3 in Fig. 5a and b, respectively) [36]. Moreover, a Chinese-script type Al–Fe–Si precipitate is present (punctual analysis #4 in Fig. 5a), probably corresponding to α-Al-Fe-Si, maintained due to the adsorption of rare earths on its surface and hindering Si diffusion for its transformation to β-Al-Fe-Si. In addition, some brighter precipitates are presented, corresponding to ASR and ASFR, due to the high atomic number of rare earths. Moreover, compared to ASF, these precipitates tend to be shorter and coarser. Furthermore, rounding ASF due to rare earth additions could enhance ductility and hot deformation ability.

Of the precipitates identified, two types are observed at grain boundaries, the first showing a needle-like shape in both micrographs and the second type being rounder and coarser but only present in the A-CRE-1D sample. According to previous reports, needles are typical for Fe-containing precipitates, and the coarser shorter phase could correspond to precipitates influenced by rare earth content, tending to be rounder. Al–Fe–Si compounds in this kind of aluminum alloy are formed during material solidification and do not present a significant diffusion or transformation in the solid state; it would not be expected to dissolve during the solution before quenching since Fe-containing phases can only be partially dissolved during solution treatment and therefore remain present after aging [35], [32]. The presence of the rare earths observed as precipitates at grain boundaries is associated with grain refinement that results in increased strength due to decreased grain boundary mobility and inhibition of simple diffusion of the species present [12].

In addition, a reaction affinity between Si, Fe, and rare earths has also been recently reported. Tang P et al. [16] used a mixture of rare earth for an Al–Si–Fe–Cu aluminum alloy since the simultaneous addition of rare earth elements is more effective in grain refinement and coarse phase modification, tending to be more rounded. In agreement with their results, they reported forming a rare earths-iron-silicon-bearing intermetallic phase, characterized as a high-temperature melting phase, even higher than the α-Al matrix. They, therefore, can induce microstructure transformation [16]. As mentioned, rare earths have a low dissolving capacity in α-Al or eutectic phases with Si.

Moreover, atomic radii and masses of these elements are much larger than those of Al or Si. Therefore, their atoms tend to be distributed at dendrite surfaces, acting as obstacles for the solidification front. This fact leads to a decrease in species in diffusion, that is, a solute accumulation at solidification waves surface and a reduction of constitutional overcooling. Moreover, they have a gathering effect against Fe [16]. It has been reported that Fe concentration decreases proportionally to the addition of rare earths. In this sense, Fe tends to form an intermetallic phase with rare earths. Additionally, rare earth forms an adhered diffusion barrier over the surface of this compound, inhibiting Fe diffusion and presence in α-Al solid state solution [16].

4 Conclusions

The study's findings on the behavior of aluminum alloy 6201 after artificial and natural aging heat treatment are significant. The addition of rare earth elements was found to induce an increase in hardness after artificial aging and a possible decrease in hardening rate during storage at room temperature. These effects can be further enhanced if the material is subjected to pretreatment, underscoring the importance of this research in the field of materials science and engineering.

The behavior of the material during artificial aging heat treatment is a complex process influenced by several factors. Excess silicon, the addition of rare earths, and the initial thermal condition were found to have a significant effect. The study suggests that a higher nucleation site density, a more extended incubation period, and clusters with a closer stoichiometric relation to the upcoming metastable precipitates could have formed. Additionally, a possible influence over precipitation kinetics of a vacancy trapping effect could have been presented due to rare earth addition, further adding to the complexity of the research.

• Natural aging was analyzed in a period close to three months, where a lower increase in the hardening rate and hardness stabilization due to rare earth addition and pre-treatment were observed. Such a behavior could be related to a vacancy-trapping effect of rare earth and cluster formation due to pre-aging. Hence, hardening during natural aging is influenced not only by the heat treatment but also by the composition since, in the first place, there is a decrease in the hardening rate when samples containing rare earth are pre-treated. Second, compared to AA6201 without rare earths but also pre-treated, the CRE-Sample presented a hardness stabilization for close to 3 months.

• Four types of precipitates associated with different processing stages were observed in the alloy under study: Al–Fe–Si intermetallic phases are commonly formed during solidification, being needle-like in two dimensions (β-Al-Fe-Si). Still, rare earths induced a morphology change, turning them shorter, coarser, and Chinese-script-like by keeping α-Al-Fe-Si until room temperature. Furthermore, rare earths could have reacted with Fe to form Al–Fe–Si-rare earth compounds, not just Al–Si-rare earth precipitations.

Funding

The authors thank Minciencias (Colombian Agency), Universidad del Valle, and Centelsa S.A of the Viakable Group for financial support under Contract 0163-2014. We also thank the Universidad de San Buenaventura, Cali (34606041).

Competing interest

The authors do not have any competing interests.