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Heat treatment of the new high-strength high-ductility Al–Mg–Si–Mn alloys with Sc, Zr and Cr additions
Materialia 15 (2021) 100981
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Heat treatment of the new high-strength high-ductility Al–Mg–Si–Mn alloys with Sc, Zr and Cr additions O. Trudonoshyn a,b,∗, O. Prach c, P. Randelzhofer a, K. Durst c, С. Körner a a
Friedrich-Alexander-Universität Erlangen-Nürnberg, Martensstraße 5, 91058 Erlangen, Germany Vyatskiy State University, Moscow street 36, 610000 Kirov, Kirov region, Russia c Technische Universität Darmstadt, Alarich-Weiss-Straße 2, 64287 Darmstadt, Germany b
a r t i c l e Keywords: Al–Mg–Si alloys Heat treatment Mechanical properties Scandium addition
i n f o
a b s t r a c t Heat treatment of alloys of the Al–Mg–Si system essentially affects the age-hardening response and, accordingly, their final mechanical properties. However, the specific Mg/Si ratio in casting alloys makes age-hardening by heat treatment negligible. To improve the age-hardening response, the alloys were alloyed with Sc, Cr, Zr, and two types of heat treatment were applied: full T6 (solution treatment (ST), quenching, and artificial aging (AA)); and T5 with AA from the as-cast state. The microstructure after different heat treatment modes was observed by means of SEM and TEM. It was established that the main structural components that affect the mechanical properties of alloys after heat treatment are nanoscale precipitates and a coarse Al7 Cr and Al3 Zr intermetallic phases in the Cr- and Zr-containing alloys. The highest impact in terms of tensile properties was achieved in the Sc-containing alloys after T5 at 325 °C. Alloys with 0.2 wt% Sc exhibit the highest strength after T5 from as-cast state at 325 °C (up to 380 MPa) and highest elongation (up to 20%) after full T6 (ST 520 °C and AA 325 °C).
1. Introduction Al–Mg–Si casting alloys (related to the pseudo-binary section) have relatively high solidus temperatures compared to other commercial Al alloys. A relatively high temperature (up to 575 °C) can be applied to this series of alloys during heat treatment [1–4]. Despite this, the temperature may be limited due to the features of the chosen casting process. Thus, Lumley et al. [5] reported that heat treatment (HT) of the Al–Si high-pressure die casting (HPDC) alloys with the high temperatures of the ST (over 500 °C) can lead to the blistering due to the presence of gases inside the cast. Therefore, the time and temperature of the solution treatment should be controlled taking into account the following [6]: • the dissolution of some equilibrium phases, dispersoids, and precipitates; • the uniform distribution of solute alloying elements; • avoiding the promotion of porosity growth. Since the alloys of the Al–Mg–Si system are promising as hightemperature alloys [7], the transition metals are interesting as additives for further design of high-temperature Al alloys. Cr and Mn are sometimes introduced individually or together to improve 2XX.X and 3XX.X alloys’ elevated temperature properties. Both of them can form the strengthening dispersoids in Al solid solution [8,9]. ∗
Knipling et al. [10] reported that the most promising method for strengthening Al alloys for producing high-temperature alloys is precipitation strengthening by stable nanosized trialuminides with a similar crystal structure to Al. He formulated four principles for the selection of elements for the alloying of the Al alloys to obtain high-temperature alloys: • Solid-state precipitation by coherent with the matrix L12 trialuminides, to achieve a high level of strengthening with low coarsening; • Shallow 𝛼-Al solvus line to achieve the precipitates with the maximum density; • Low solid solubility and diffusivity at the elevated temperatures (during aging) to minimize Al3 M coarsening; • Partition coefficient k~1, to minimize segregation. A shallow Al3 M liquidus boundary is desirable for peritectic systems, to minimize the casting temperature and suppressing crystallization Al3 M as a primary phase. A large number of transition metals can form a thermodynamically stable dispersed phase (Al3 M trialuminides) with a similar crystal structure to Al and with a low lattice parameter mismatch with the Al. The transition elements from groups 3, 4, and 5 form high-symmetry L12 , D022, and D023 structures. The later groups form less symmetrical structures (Fe, Co, Ni, Re, Ir) [10]. It should also be mentioned that most of
Corresponding author at: Friedrich-Alexander-Universität Erlangen-Nürnberg, Martensstraße 5, 91058 Erlangen, Germany. E-mail address: [email protected] (O. Trudonoshyn).
https://doi.org/10.1016/j.mtla.2020.100981 Received 5 August 2020; Accepted 6 December 2020 Available online 8 December 2020 2589-1529/© 2020 Published by Elsevier B.V. on behalf of Acta Materialia Inc.
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Materialia 15 (2021) 100981
Table 1 Nominal composition of the investigated alloys (Al-balance). Chemical composition [wt%] Alloy Z1 Z2 C1 C2 S1 S2 Base alloy
Impurity [wt%]
Mg
Si
Mn
Cr
Zr
Sс
Fe
Ti
Be
Cu/Zn
5.5 5.5 5.5 5.5 5.5 5.5 5.5
2.5 2.5 2.5 2.5 2.5 2.5 2.5
0.7 0.7 0.7 0.7 0.7 0.7 0.7
– – 0.1 0.2 – – –
0.1 0.2 – – – – –
– – – – 0.1 0.2 –
0.2 0.2 0.2 0.2 0.2 0.2 0.2
0.2 0.2 0.2 0.2 0.2 0.2 0.2
0.004 0.004 0.004 0.004 0.004 0.004 0.004
0.15 0.15 0.15 0.15 0.15 0.15 0.15
Table 2 Casting parameters.
the lanthanides, some of the early actinides, and some other elements such as Li (Al3 Li [11,12]) can also form trialuminides. Of all these elements, Sc and Zr deserve special attention [10]:
Preheated temperature, Tp Melt temperature, Tm Mold temperature, Td Piston speed, 𝜐 Final (post) pressure, p
• The Al–Sc system has a shallow solvus curve that promotes precipitation strengthening; Al3 Sc is thermodynamically stable with the L12 structure. However, Sc is very expensive [13]. • The Al–Zr system has one of the lowest diffusion rates. Al3 Zr has one of the lowest lattice parameter mismatches with Al. And it has a reasonable price among the available candidates. However, the L12 structure of Al3 Zr is metastable. Furthermore, the Al–Zr system is peritectic that limits the possible concentration of Zr in solid solution.
250 °C 720 °C 200 °C 2 m/s 250 bar
casting plates with the dimensions of 300 × 150 × 4 mm [24,28]. Table 2 represents the casting parameters. Two types of heat treatment were applied to the studied alloys: • full T6: solution treatment (ST) at 480, 520, and 575 °C followed by quenching into the room temperature (RT) water and further artificial aging (AA) at 225 °C and 325 °C temperatures; • T5, AA from the as-cast state at 125 °C, 175 °C, 225 °C and 325 °C temperatures.
Rheinfelden Alloys [14] recently developed alloy Magsimal-59 Plus alloy with the addition of Zr. They reported that the addition of Zr into the AlMg5Si2Mn alloy promotes the formation of Zr-containing nanosized precipitates. Nevertheless, in the works [15,16], the Zr-containing precipitates after conventional heat treatment modes were not found. Sc-containing precipitates form at relatively high aging temperatures (compared to standard precipitates formed in the Al–Mg–Si system) [17]. This is a disadvantage for the alloying of Al–Mg–Si alloys. Since with temperatures higher than 250 °С 𝛽" and even 𝛽’ precipitates dissolve with further formation of equilibrium 𝛽 phase [18–20]. However, in alloys with the high Mg content, a very small amount of 𝛽" or 𝛽’ is formed during aging at conventional temperatures, or even not at all, and the dominant phase is equilibrium 𝛽 phase that can be found already in as cast state [21,22]. In the previous studies [16,23,24] we developed the new HPDC alloys based on the Al–Mg–Si–Mn alloy, additionally doped by Sc, Cr, Zr and Zn were designed. We found that Cr and Zr slightly improve the yield strength and ultimate tensile strength but lead to degradation in elongation to failure. Sc has the most prominent effect on the microstructure of Al–Mg–Si alloys [16,25] as well as on the mechanical properties, even at concentrations of only 0.1%. Sc acts as a solid solution strengthening element and forms Al3 Sc precipitates. In the present study, we investigated changes in the alloys’ mechanical properties, depending on the heat treatment modes.
A TESCAN Mira3 SEM (operated at 15 kV) with Energy Dispersive Spectrometry (EDS) detector was used for the structural evaluation. The Philips CM20 (operated at 200 kV) was used for TEM investigations. The tensile tests were performed using the Instron 5967. A crosshead speed was set at 1 mm/min (initial strain rate was 8.3 × 10−4 1/s). A rectangular cross-section specimens [24,28] (25 mm gauge length and 4 mm thickness) according to the ASTM E8_E8M_13a were used. The hardness (HB) was measured using a Brinell tester with 306.56 N load, 2.5 mm ball diameter, during 10 s. The average values of hardness and tensile properties for each alloy were taken from 10 measurements. 3. Results 3.1. Selection of heat treatment temper for the studied alloys Fig. 1 shows a comparison of the microstructure (a,b) and the surface (с) for the base Al–5.7Mg–2.6Si–Mn alloy, with different solution treatment temperatures. One can see the temperatures higher than 540°С are inapplicable for the heat treatment of the studied HPDC-parts. However, such low solution temperatures (lower than 520°С) for the studied alloy system do not lead to a maximum of hardness, but in a combination of further overaging leads to significant improvement of ductility levels [29]. Therefore, we choose the several temperatures of the ST followed by AA (T6 regime) for the preliminary experiments. Fig. 2 shows the hardness measurement results across the plate after T6 regime at different ST temperatures for the base alloy. Sample after T6 regime (ST at 570 °C+AA at 175 °C) showed the maximum hardness of 115HV. Hardness values of the samples treated at lower temperatures have not even reached the hardness value of the as-cast state. Table 3 shows the literature data of the applied temperatures for heat treatment of the binary Al-alloys compared to the base alloy. From the thermodynamic calculations [16], we determined the required temperatures of the solution treatment. Table 4 contains the values of solubility of the chosen alloying elements in the current system.
2. Materials and methods Table 1 represents the composition of the studied alloys. The amount of alloying elements was chosen according to the literature data [17,26,27]. To perform thermodynamic calculations, a ThermoCalc software (TC) with the TCAl2:Al-alloys v2.1 database was used to predict the effects of the studied system’s selected elements. Еhe commercial Magsimal59○R ("Rheinfelden Alloys GmbH", Germany) and commercial AlSc2, AlZr15 and AlCr20 master alloys were used as starting materials. The starting materials’ blocks were stored in the preheated furnace for 24 h to avoid moisture before melting. The melt was impelled to minimize microporosity. The density was controlled by a MK 3VT vacuum density tester in combination with MK3000 the density index balance. A HPDC unit with the cold chamber (Frech DAK 450–54 (with 4500 kN mold clamping force) was used to produce 2
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Materialia 15 (2021) 100981
Fig. 1. Microstructure (a,b) and the surface (с) for the base Al–5.7Mg–2.6Si–Mn alloy, with different solution treatment temperatures: a) 520 °C, b) 575 °C. Pores are marked with arrows.
Table 4 The solubility of chosen elements in the Al–5,7Mg,−2,5Si–0,6Mn–X systems based on calculated diagrams. Alloying element Zr Sc Cr
Solubility at: [wt%] 591 °C
540 °C
520 °C
500 °C
480 °C
460 °C
0.07 0.12 0.35
0.05 0.09 0.25
0.04 0.07 0.21
0.03 0.05 0.17
0.02 0.03 0.12
0.01 0.02 0.1
Table 5 Average size of separate Mg2 Si particles as a function of treatment temperature (treatment time – 1.5 h). Temperature °C Diameter, 𝜇m
Table 3 Temperatures for treatment of the alloys doped by chosen elements [°C] [17,19,30–32].
Al–Sc Al–Zr Al–Cr Base alloy
Aging
500– 600
250–350 – – 170–250
480–575
570 0.82
3.2. Full T6 treatment The changes in the structure and properties of Al–Mg–Si–Mn alloys during heat treatment result from several processes that occur simultaneously during heating. The first process is the spheroidization of eutectic fibrous (Fig. 3). A higher solution-treatment temperature leads to a faster eutectic cells decomposition into separate Mg2 Si particles and faster coarsening of these particles. Table 5 shows the average size of these particles depending on the treatment temperature. During the solution treatment, several processes occur in the solid solution. First of all, high-temperature treatment leads to dissolving all strengthening precipitates (including GP-zones, 𝛽’) that form during natural and artificial aging [3,6,15,22]. However, the equilibrium 𝛽-Mg2 Si phase after solution treatment is still present in the microstructure (Fig. 4a,b). The second process is the formation of Mn-containing
Temper Solution treatment
520 0.66
in the a-Al matrix. Further, the alloys can be aged similarly to the base alloy. It does not make sense to perform solutionizing Sc- and Zrcontaining studied alloys. However, Sc-containing alloys can be aged (with the formation of nanoscale Al3 Sc and Al3 (Sc,Zr) precipitates in the Al-matrix) in the range of temperatures 250–500 °C.
Fig. 2. Hardness across the plate after heat treatment and in as-cast state.
Alloying element
480 0.59
Thus, during heat treatment with 520 °С, the new phases in the studied alloys can be dissolved at 0.07 wt% Sc, 0.21 wt% Cr and 0.04 wt% Zr. Based on the above, Cr-containing series of alloys can be heat treated at 520 °C, followed by quenching, to form Cr-containing dispersoids 3
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Materialia 15 (2021) 100981
Fig. 3. Microstructure of the base alloy in a) as-cast state and b) after solution treatment at 480 °C during 1.5 h; c) after solution treatment at 520 °C during 1.5 h; d) after solution treatment at 570 °C during 1.5 h.
Table 6 Tensile properties of the base alloys after full T6 treatment. Temperature [°C] Solution treatment 570 520 480 As-cast
Artificial aging
UTS [MPa]
175 225 175 225 175 225 [16]
291.3 293.7 280.6 242.0 241.6 228.1 299.8
± ± ± ± ± ± ±
2.8 0.5 8.7 12 2.1 1.0 6.3
YS [MPa]
A [%]
235.8 ± 1.1 201.3 ± 9.3 184.8 ± 2.6 130.0 ± 6.6 98.3 ± 0.8 93.7 ± 0.9 163.9 ± 5.6
3.8 ± 0.4 6.0 ± 2.6 9.5 ± 2.3 11.5 ± 1.3 16.8 ± 3.8 18.0 ± 2.9 8.3 ± 0.5
state. The situation is different with treatment at lower temperatures (480–520 °C). Despite the low values of both tensile strength and yield strength, the elongation values are very high and can reach values of 18%. According to the equilibrium phase diagram [16], Al7 Cr is a refractory phase and can be dissolved at relatively high heat-treatment temperatures. In C1 (0.1 wt% Cr) Al7 Cr phase dissolves at 460 °C; in C2 (0.2 wt% Cr) at 520 °C. However, Fig. 7b shows that heat treatment with a temperature of 520 °C even for 3 h is not enough for the dissolution of the Al7 Cr phase. As it was predicted utilizing Thermo-Calc, Al3 Zr and Al3 Sc phases do not dissolve during applied solution treatment. AlMg2 Si eutectic in the studied alloys after heat treatment changes in the same way as the base alloy (eutectic disintegrated into separate spheres). Heat treatment also does not affect the morphology and distribution of Mn-containing intermetallics. All studied additions do not affect the aging behavior of the 𝛽-type precipitates in the studied alloys. Similarly to the base alloy, solution treatment at 520 °C with further aging doesn’t lead to the formation of high density 𝛽-type precipitates (GP-zones and 𝛽’). However, Cr and Zr affects the formation of dispersoids (Fig. 8) inside a-Al dendrites: Cr-containing alloys after solution treatment have (Mn,Cr)-containing dispersoids, while Zr addition leads to the formation Al3 Zr dispersoids. Fig. 9 represents the results of the hardness and tensile tests of studied alloys with different aging temperatures for the time of maximum hardness values (3 h). The higher aging temperature leads to the lower strength (and hardness) and the higher elongation. In general, the behavior of changes is similar to the base alloy.
dispersoids. Fig. 4c and d shows that dispersoids after high-temperature heat treatment form attaching to spherical Mg2 Si particles and the equilibrium 𝛽-Mg2 Si phase. Fig. 5a and b shows that during aging after solution treatment at the temperature of 520 °C, strengthening phases are practically not formed. However, solution treatment at 570 °C with further aging (Fig. 5c and d) leads to the formation of the solid solution’s high-density GP-II zones. Table 6 and Fig. 6 show the mechanical properties of the base alloy after treatment with different temperatures. Results of the hardness measurements (Fig. 6) show that aging after low-temperature solution treatment (e.g., temperature range 480–520 °C) does not lead to a significant increase in the hardness of the base alloy. On the other hand, aging after solution treatment at 570 °C increases the hardness by 25% compared to the solutionized state and by 12% compared to the as-cast
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Materialia 15 (2021) 100981
Fig. 4. TEM images of the base alloy after solution treatment at 520 °C a,b) and at 570 °C c,d) (without aging).
3.3. Artificial aging from the as-cast state
hardness of the Sc-containing series of alloys, especially S2 alloy (with 0.2 wt% of Sc), increases significantly and reaches values about 120HB. Fig. 15 represents the results of tensile tests of studied alloys for different aging temperatures for the time of maximum hardness values (3 h). Both Cr-containing and Zr-containing alloys show similar changes and similar levels of all tensile properties. Maximum values of strength belong to the aging temperature range 175–225 °C. That coincides with the lowest values of elongation. S1 alloy (0.1 wt% Sc) shows approximately similar results of all tensile tests for all aging temperatures. Aging temperature higher than 125 °C leads to a minor increasing of yield strength. S2 alloy shows the highest values of both tensile and yield strength, especially for aging treatment with 325 °C. Fig. 16 shows the results of tensile properties for aging with 325 °C as a function of the time. The aging time increase reduces both tensile and yield strength simultaneously with increasing elongation for all studied alloys.
Fig. 10a shows changes in the hardness of the base alloy with different temperatures and aging times. Thus, from the diagram, it is clear that the base alloy’s heat treatment does not have a significant effect on the hardness. The most noticeable changes were recorded at the aging temperature of 175 °C and time in the range of 2–3 h. However, the increase in hardness in this mode is only 5–7% higher than in the cast state. Fig. 10b shows the results of tensile properties of the base alloy with different temperatures of aging for time with peak hardness (3 h). One can see the increase in tensile strength is insignificant, but yield strength has more prominent changes. Nevertheless, the elongation is decreased. Results of SEM and TEM investigations of the base alloy do not show any significant changes neither in microstructure nor in the structure of the 𝛼-Al dendrites during treatment in a temperature range of 125– 325 °C. The microstructure consists of 𝛼-Al dendrites, Al-Mg2 Si eutectic colonies, and 𝛼-AlMnFeSi intermetallics. The aging in such temperature range and time does not lead to changes in all phases’ morphology in comparison to as-cast state (Fig. 11). The structure of 𝛼-Al dendrites consists of 𝛽’ and stable 𝛽 phases similar to the as-cast state [24] (Fig. 12). Similar to the base alloy, all studied alloys don’t show any changes in the microstructure after heat treatment in the temperature range of 125– 325 °C. However, alloys with Sc addition show increasing nanoscale precipitates’ density after high-temperature aging treatment (325 °C) (Fig. 13). All studied series of alloys show approximately similar values of hardness after aging with temperatures 125–225 °C (Fig. 14). However, the aging with temperature 325 °C shows different behavior. Thus, the
4. Discussion 4.1. Full T6 heat treatment The changes in the properties of Al-Mg-Si-Mn alloys during hightemperature solution treatment (as the first stage of the full T6 treatment) result from the changes of structural components that occur during heating. Since Mg2 Si eutectic is the main strengthening phase in this type of alloys, the process of its disintegration reduces the hardness and strength that coincides with the increase of elongation. Trudonoshyn et al. [15,28] reported a similar tendency for the casting Al-Mg-Si-Mn alloys cast into the permanent mold. Together with eutectic disinte5
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Materialia 15 (2021) 100981
Fig. 5. TEM images of the base alloy after solution treatment at 520 °C a,b) and at 570 °C c,d) and further aging at 175 °C (solution treatment time – 1.5 h; aging time – 3 h).
higher hardness values of the HPDC compared to the regular permanent mold casting alloys [15,28]. Artificial aging after solution treatment causes the formation of the strengthening precipitates again [15] that improve the strengthening properties of the alloys. Nevertheless, different solution treatment temperatures have different effects. Aging after solution treatment at 570 °C has a more noticeable effect on the hardness. This may be due to the formation of a larger number of quench-in vacancies that promote the diffusion of the solute atoms, which in turn, promote the formation of atoms clusters and GP-zones [36–39]. However, such heat treatment is not recommended by the manufacturer [14]. The different situation is with treatment at lower temperatures (480–520 °C). Thus, despite the low values of both tensile strength and yield strength, the elongation values are very high and can reach values of 18%, as confirmed by the results presented in [29]. All studied series of alloys after full T6 treatment have the same microstructural changes as the base alloy. Moreover, solution treatment at 520 °C during 1.5 h (and even 3 h) does not lead to dissolution of the Al3 Zr and Al3 Sc phases in Zr- and Sc-containing alloys respectively (as it was predicted by Thermo-Calc software) and even the Al7 Cr phase in Cr-containing alloys (despite been predicted by Thermo-Calc software). A primary cause of this is the low diffusion rate of Cr. This phase’s dissolution could improve the tensile properties of the alloy, but this did not happen. Moreover, Samuel et al. [40] reported that sludge phases remained stable after solution treatment for 8 h at 515 °C. Solute Cr in 𝛼-Al dendrites during solution treatment promotes the formation of (Mn,Cr)-containing dispersoids, which, however, do not significantly affect the mechanical properties of the casting alloys. In
Fig. 6. Brinell hardness of the base alloys after full T6 treatment with different solution treatment and aging temperatures and different treatment times.
gration, the solution treatment usually leads to the dissolution of the strengthening 𝛽’’-Mg2 Si precipitates that might be formed in the as-cast state of Al-Mg-Si alloys after solidification [33,34]. However, due to the specific Mg/Si ratio of the base alloy, the amount of the strengthening 𝛽’’-Mg2 Si precipitates in as-cast (and even after aging [21,22]) is negligible. Solution treatment of the Al5Mg2SiMn permanent mold casting leads to the formation of 𝛽-AlMnFeSi dispersoids [15,28]. Their lack of coherence with the matrix leads to a reduction in the alloys’ hardness still more (along with decomposition of eutectic lamellas). On the other hand, in the studied HPDC Al-Mg-Si-Mn alloys, we observed 𝛼-AlMnFeSi dispersoids. This type of dispersoids is semi-coherent with the Al-matrix [35] that can slightly affect the hardness of the alloys. This explains the 6
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Materialia 15 (2021) 100981
Fig. 7. SEM images of Z2 (a), C2 (b), and S2 (c) alloys after solution treatment at 520 °C (solution treatment time – 3 h).
Fig. 8. TEM images of the Mn,Cr-containing dispersoids in C2 alloy (a) and Al3 Zr dispersoids in Z2 alloy after solution treatment at 520 °C for 1.5 h.
contrast to the Zr-series of alloys which have the highest values of hardness and strength after full T6 treatment among all studied alloys. According to [31], this effect is caused by the formation of the Al3 Zr dispersoids in 𝛼-Al dendrites’ structure. Except for the high cost, Sc has another weakness as an alloying element for Al-Mg-Si HPDC alloys. Since the Al3 Sc phase does not dissolve at the applied solution treatment temperatures (Table 6), it makes no sense to perform full T6 treatment on the investigated Sc-containing alloys. In contrast to the situation with binary Al–Sc alloys [31], which can be treated using temperatures above the solvus (a precondition for age strengthening is that the strengthening phase must be soluble between solvus and solidus temperatures) and thus, get a significant increase in the mechanical properties. Despite the good mechanical properties of the Zr-containing series, compared to Cr- and Sc-containing series after full T6 treatment, the strength values of all studied alloys are still lower than in the as-cast state, which suggests that the eutectic is the more strengthening phase in the Al-Mg-Si casting alloys.
and time in the range of 2–3 h. However, the increase in hardness in this mode is only 5–7% higher than the as-cast state. Such an insignificant effect is concerned with the fact that the base alloy belongs to alloys with an excess of Mg content. Mg drastically reduces the solubility of Mg2 Si in the 𝛼-Al, which in turn reduces the hardening potential (by the formation of nanoscale precipitates) and also increases the stability of the 𝛽’ and equilibrium 𝛽 phases (which do not have a significant hardening effect) [21,22]. In general, it can be argued that excess of Mg makes the properties of the base alloy very stable at elevated temperatures (which makes the material promising for further research at elevated temperatures). The microstructural investigation confirms the results of the mechanical properties, which shows that the aging treatment in the studied temperature range (125–325 °C) and time (up to 15 h) does at least not change the morphology of the phases (Fig. 11). The structure of 𝛼-Al dendrites mostly consists of 𝛽’ and stable 𝛽 phases similar to the as-cast state [23,24]. However, in some 𝛼-Al dendrites, only small amounts of 𝛽"-precipitates were found after aging at 175 °C (Fig. 12c). Similar to the base alloy, alloys additionally doped by Cr and Zr have changes neither in microstructure nor in precipitation sequence after heat treatment in the temperature range of 225–325 °C. We did not observe any new Cr- or Zr-containing precipitates. Both series of alloys show similar to the base alloy behavior of all studied mechanical properties after aging (Fig. 14, Fig. 15). Despite this, aging from the as-
4.2. Artificial aging from as-cast state The artificial aging of the base alloy at the temperature range 125– 325 °C does not significantly affect the hardness (Fig. 10). The most noticeable changes were recorded at the aging temperature of 175 °C 7
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Materialia 15 (2021) 100981
Fig. 9. Mechanical properties of the studied alloys after full T6 heat treatment (solution treatment at 520 °C during 1.5 h, artificial aging during 3 h).
cast state remains promising for Al-Mg-Si-Mn alloys additionally alloyed by other elements (e.g., Sc) [41]. Aging temperatures for the formation of nano-dispersed Al3 Sc precipitates are significantly higher than the temperatures used to obtain maximum hardness values in Al–Mg–Si alloys by the formation of 𝛽’’-
Mg2 Si precipitates [17]. Such temperatures (250–350°С) lead to overaging and make more stable 𝛽’- and equilibrium 𝛽-Mg2 Si phases (which do not significantly affect the strength). On the other hand, as already mentioned, the base alloy has rather stable properties and structure during aging up to 325°С. Aging of Sc-containing alloys at a temperature
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Fig. 10. Brinell hardness of the base alloy as a function of aging temperature and time (a) Tensile properties of the base alloy as a function of aging temperature (3 h) (b).
Fig. 11. SEM-images of the microstructure of the base alloy: a) as-cast state; b) aging at 175 °C; c) aging at and 325 °C.
Fig. 12. TEM bright-field images of the base alloy after artificial aging: a–c) base alloy after aging at 175 °C; d–f) base alloy after aging at 325 °C.
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Fig. 13. TEM bright-field images of the Sc-containing alloys after artificial aging:a) S1 after 225 °C; b) S1 325 °C; c) S2 225 °C; d) S2 325 °C;.
Fig. 14. Brinell hardness of the studied alloys as a function of aging temperature and time: a) Z1 (0.1 wt% Zr); b) Z2 (0.2 wt% Zr); c) C1 (0.1 wt% Cr); d) C2 (0.2 wt% Cr); e) S1 (0.1 wt% Sc); f) S2 (0.2 wt% Sc). 10
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Fig. 15. Tensile properties of the studied alloys as a function of concentration of alloying elements for peak-hardness aging time: a) Z1; b) Z2; c) C1; d) C2; e) S1; f) S2.
Fig. 16. Tensile properties of the studied alloys as a function of the aging time for 325 °C.
325 °С from the as-cast state is a promising heat treatment. The strength of S2 alloy (with 0,2 wt% Sc) increased up to 380 MPa (11% higher than the as-cast state of S2 and 30% higher than base alloy). The formation of nano-dispersed precipitates Al3Sc can only explain such a significant increase in hardness and strength. Such an increase in properties is in good agreement with the existing studies concerning the alloying of Alalloys by Sc [31,32,41].
implementing the commercial analogs of the Cr-containing and Zrcontaining alloys in the production (Maxxalloy-Ultra, Magsimal Plus). Sc has the most significant influence on the base composition in the as-cast state among all elements. Subsequent artificial aging from the as-cast state increases the strength of the alloys, but it decreases elongation. The reasons for the change in properties in the studied alloys were the formation of new primary phases in the as-cast structure and the formation of nanosized precipitates during the heat treatment. It was found that coarse phases (such as Al7 Cr in Cr-containing alloys and Al3 Zr in Zrcontaining alloys) are brittle and significantly reduce the tensile properties of alloys, especially ductility. The Al3 Sc phase (Sc-containing alloys) did not negatively affect the ductility of the alloys. The primary Al7 Cr, Al3 Zr, Al3 Sc phase doesn’t dissolve even during high-temperature solution treatment. However, Sc that remains in 𝛼-Al during solidification
5. Conclusions This study represents the microstructural evolution and Al-Mg-Si-Mn alloys’ mechanical properties with additional alloying by Cr, Zr, and Sc in different heat treatment tempers. The analysis of all mechanical properties of the studied alloys shows that Cr and Zr have insignifficant effect in all tempers, despite 11
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Materialia 15 (2021) 100981
is sufficient for the formation of a nanoscale strengthening precipitates during aging. As expected, base alloy showed a good response to additional alloying with elements that enhance the precipitation strengthening. Thus, the greatest impact on the mechanical properties was affected by nanoscale Al3 Sc strengthening precipitates formed in 𝛼-Al solid solution in Sc-containing alloys during aging. Although the Cr addition to the base alloy leads to the formation of a dispersed Cr-containing phase in 𝛼-Al, it does not have a significant strengthening effect (similar to Cr-containing wrought alloys).
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Declaration of Competing Interest The authors declare no conflict of interest. Acknowledgments O. Prach and O. Trudonoshyn are grateful to German Academic Exchange Service (DAAD) for the financial support over this research. Supplementary materials Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.mtla.2020.100981. References [1] S. Babaniaris, M. Ramajayam, L. Jiang, T. Langan, T. Dorin, Tailored precipitation route for the effective utilisation of Sc and Zr in an Al-Mg-Si alloy, Materialia 10 (2020) 100656, doi:10.1016/j.mtla.2020.100656. [2] O. Trudonoshyn, O. Prach, V. Boyko, M. Puchnin, K. Mykhalenkov, Design of a new casting alloys containing Li or Ti+Zr and optimization of its heat treatment, in: Proceedings of the 23rd International Conference on Metallurgy and Materials, METAL 2014, 2014, pp. 1399–1404. http://metal2014.tanger.cz/files/proceedings/17/reports/2750.pdf. [3] T. Dorin, M. Ramajayam, S. Babaniaris, L. Jiang, T.J. Langan, Precipitation sequence in Al–Mg–Si–Sc–Zr alloys during isochronal aging, Materialia 8 (2019) 100437, doi:10.1016/j.mtla.2019.100437. [4] O. Trudonoshyn, M. Puchnin, O. Prach, Use of the ABI technique to measure the mechanical properties of aluminium alloys: effect of heat-treatment conditions on the mechanical properties of alloys, Mater. Tehnol. 50 (2016), doi:10.17222/mit.2014.295. [5] R.N. Lumley, R.G. Odonnell, D.R. Gunasegaram, M. Givord, Heat treatment of high-pressure die castings, Metall. Mater. Trans. A 38 (2007) 2564–2574, doi:10.1007/s11661-007-9285-4. [6] F. Yan, W. Yang, S. Ji, Z. Fan, Effect of solutionising and ageing on the microstructure and mechanical properties of a high strength die-cast Al-Mg-Zn-Si alloy, Mater. Chem. Phys. (2015) 1–9, doi:10.1016/j.matchemphys.2015.10.014. [7] C. Li, Y. Wu, H. Li, Y. Wu, X. Liu, Effect of Ni on eutectic structural evolution in hypereutectic Al–Mg 2 Si cast alloys, Mater. Sci. Eng. A 528 (2010) 573–577, doi:10.1016/j.msea.2010.09.056. [8] L. Lodgaard, N. Ryum, Precipitation of chromium containing dispersoids in Al–Mg–Si alloys, Mater. Sci. Technol. 16 (2000) 599–604, doi:10.1179/026708300101508315. [9] L. Lodgaard, N. Ryum, Precipitation of dispersoids containing Mn and / or Cr in Al–Mg–Si alloys, Mater. Sci. Eng. A 283 (2000) 144–152, doi:10.1016/S0921-5093(00)00734-6. [10] K.E. Knipling, D.C. Dunand, D.N. Seidman, Criteria for developing castable, creep-resistant aluminum-based alloys – a review, Z. Metallkd 97 (2006) 246–265. [11] O. Prach, J. Hornik, K. Mykhalenkov, Effect of the addition of Li on the structure and mechanical properties of hypoeutectic Al-Mg2Si alloys, Acta Polytech. 55 (2015) 253–259, doi:10.14311/AP.2015.55.0253. [12] O. Trudonoshyn, Studying the structure of Al-Mg-Si casting alloys doped by lithium, Phys. Met. Metallogr. 121 (2020) 701–707, doi:10.1134/S0031918X2007011X. [13] S. Riva, K.V. Yusenko, N.P. Lavery, D.J. Jarvis, S.G.R. Brown, K.V. Yusenko, N.P. Lavery, D.J. Jarvis, S.G.R. Brown, S. Riva, K.V. Yusenko, N.P. Lavery, D.J. Jarvis, S.G.R. Brown, The Scandium Effect in Multicomponent Alloys the Scandium Effect in Multicomponent Alloys, 6608 (2016). doi:10.1080/09506608.2015.1137692. [14] Rheinfelden Alloys GmbH & Co. KG, Ductile aluminum alloys for automotive structural applications, Rheinfelden (2017) http://rheinfelden-alloys. eu/wp-content/uploads/2015/04/Handbook-Alloys-for-Structural-Applications_ RHEINFELDEN-ALLOYS_2017_EN.pdf.
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Heat treatment of the new high-strength high-ductility Al–Mg–Si–Mn alloys with Sc, Zr and Cr additions O. Trudonoshyn a,b,∗, O. Prach c, P. Randelzhofer a, K. Durst c, С. Körner a a
Friedrich-Alexander-Universität Erlangen-Nürnberg, Martensstraße 5, 91058 Erlangen, Germany Vyatskiy State University, Moscow street 36, 610000 Kirov, Kirov region, Russia c Technische Universität Darmstadt, Alarich-Weiss-Straße 2, 64287 Darmstadt, Germany b
a r t i c l e Keywords: Al–Mg–Si alloys Heat treatment Mechanical properties Scandium addition
i n f o
a b s t r a c t Heat treatment of alloys of the Al–Mg–Si system essentially affects the age-hardening response and, accordingly, their final mechanical properties. However, the specific Mg/Si ratio in casting alloys makes age-hardening by heat treatment negligible. To improve the age-hardening response, the alloys were alloyed with Sc, Cr, Zr, and two types of heat treatment were applied: full T6 (solution treatment (ST), quenching, and artificial aging (AA)); and T5 with AA from the as-cast state. The microstructure after different heat treatment modes was observed by means of SEM and TEM. It was established that the main structural components that affect the mechanical properties of alloys after heat treatment are nanoscale precipitates and a coarse Al7 Cr and Al3 Zr intermetallic phases in the Cr- and Zr-containing alloys. The highest impact in terms of tensile properties was achieved in the Sc-containing alloys after T5 at 325 °C. Alloys with 0.2 wt% Sc exhibit the highest strength after T5 from as-cast state at 325 °C (up to 380 MPa) and highest elongation (up to 20%) after full T6 (ST 520 °C and AA 325 °C).
1. Introduction Al–Mg–Si casting alloys (related to the pseudo-binary section) have relatively high solidus temperatures compared to other commercial Al alloys. A relatively high temperature (up to 575 °C) can be applied to this series of alloys during heat treatment [1–4]. Despite this, the temperature may be limited due to the features of the chosen casting process. Thus, Lumley et al. [5] reported that heat treatment (HT) of the Al–Si high-pressure die casting (HPDC) alloys with the high temperatures of the ST (over 500 °C) can lead to the blistering due to the presence of gases inside the cast. Therefore, the time and temperature of the solution treatment should be controlled taking into account the following [6]: • the dissolution of some equilibrium phases, dispersoids, and precipitates; • the uniform distribution of solute alloying elements; • avoiding the promotion of porosity growth. Since the alloys of the Al–Mg–Si system are promising as hightemperature alloys [7], the transition metals are interesting as additives for further design of high-temperature Al alloys. Cr and Mn are sometimes introduced individually or together to improve 2XX.X and 3XX.X alloys’ elevated temperature properties. Both of them can form the strengthening dispersoids in Al solid solution [8,9]. ∗
Knipling et al. [10] reported that the most promising method for strengthening Al alloys for producing high-temperature alloys is precipitation strengthening by stable nanosized trialuminides with a similar crystal structure to Al. He formulated four principles for the selection of elements for the alloying of the Al alloys to obtain high-temperature alloys: • Solid-state precipitation by coherent with the matrix L12 trialuminides, to achieve a high level of strengthening with low coarsening; • Shallow 𝛼-Al solvus line to achieve the precipitates with the maximum density; • Low solid solubility and diffusivity at the elevated temperatures (during aging) to minimize Al3 M coarsening; • Partition coefficient k~1, to minimize segregation. A shallow Al3 M liquidus boundary is desirable for peritectic systems, to minimize the casting temperature and suppressing crystallization Al3 M as a primary phase. A large number of transition metals can form a thermodynamically stable dispersed phase (Al3 M trialuminides) with a similar crystal structure to Al and with a low lattice parameter mismatch with the Al. The transition elements from groups 3, 4, and 5 form high-symmetry L12 , D022, and D023 structures. The later groups form less symmetrical structures (Fe, Co, Ni, Re, Ir) [10]. It should also be mentioned that most of
Corresponding author at: Friedrich-Alexander-Universität Erlangen-Nürnberg, Martensstraße 5, 91058 Erlangen, Germany. E-mail address: [email protected] (O. Trudonoshyn).
https://doi.org/10.1016/j.mtla.2020.100981 Received 5 August 2020; Accepted 6 December 2020 Available online 8 December 2020 2589-1529/© 2020 Published by Elsevier B.V. on behalf of Acta Materialia Inc.
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Table 1 Nominal composition of the investigated alloys (Al-balance). Chemical composition [wt%] Alloy Z1 Z2 C1 C2 S1 S2 Base alloy
Impurity [wt%]
Mg
Si
Mn
Cr
Zr
Sс
Fe
Ti
Be
Cu/Zn
5.5 5.5 5.5 5.5 5.5 5.5 5.5
2.5 2.5 2.5 2.5 2.5 2.5 2.5
0.7 0.7 0.7 0.7 0.7 0.7 0.7
– – 0.1 0.2 – – –
0.1 0.2 – – – – –
– – – – 0.1 0.2 –
0.2 0.2 0.2 0.2 0.2 0.2 0.2
0.2 0.2 0.2 0.2 0.2 0.2 0.2
0.004 0.004 0.004 0.004 0.004 0.004 0.004
0.15 0.15 0.15 0.15 0.15 0.15 0.15
Table 2 Casting parameters.
the lanthanides, some of the early actinides, and some other elements such as Li (Al3 Li [11,12]) can also form trialuminides. Of all these elements, Sc and Zr deserve special attention [10]:
Preheated temperature, Tp Melt temperature, Tm Mold temperature, Td Piston speed, 𝜐 Final (post) pressure, p
• The Al–Sc system has a shallow solvus curve that promotes precipitation strengthening; Al3 Sc is thermodynamically stable with the L12 structure. However, Sc is very expensive [13]. • The Al–Zr system has one of the lowest diffusion rates. Al3 Zr has one of the lowest lattice parameter mismatches with Al. And it has a reasonable price among the available candidates. However, the L12 structure of Al3 Zr is metastable. Furthermore, the Al–Zr system is peritectic that limits the possible concentration of Zr in solid solution.
250 °C 720 °C 200 °C 2 m/s 250 bar
casting plates with the dimensions of 300 × 150 × 4 mm [24,28]. Table 2 represents the casting parameters. Two types of heat treatment were applied to the studied alloys: • full T6: solution treatment (ST) at 480, 520, and 575 °C followed by quenching into the room temperature (RT) water and further artificial aging (AA) at 225 °C and 325 °C temperatures; • T5, AA from the as-cast state at 125 °C, 175 °C, 225 °C and 325 °C temperatures.
Rheinfelden Alloys [14] recently developed alloy Magsimal-59 Plus alloy with the addition of Zr. They reported that the addition of Zr into the AlMg5Si2Mn alloy promotes the formation of Zr-containing nanosized precipitates. Nevertheless, in the works [15,16], the Zr-containing precipitates after conventional heat treatment modes were not found. Sc-containing precipitates form at relatively high aging temperatures (compared to standard precipitates formed in the Al–Mg–Si system) [17]. This is a disadvantage for the alloying of Al–Mg–Si alloys. Since with temperatures higher than 250 °С 𝛽" and even 𝛽’ precipitates dissolve with further formation of equilibrium 𝛽 phase [18–20]. However, in alloys with the high Mg content, a very small amount of 𝛽" or 𝛽’ is formed during aging at conventional temperatures, or even not at all, and the dominant phase is equilibrium 𝛽 phase that can be found already in as cast state [21,22]. In the previous studies [16,23,24] we developed the new HPDC alloys based on the Al–Mg–Si–Mn alloy, additionally doped by Sc, Cr, Zr and Zn were designed. We found that Cr and Zr slightly improve the yield strength and ultimate tensile strength but lead to degradation in elongation to failure. Sc has the most prominent effect on the microstructure of Al–Mg–Si alloys [16,25] as well as on the mechanical properties, even at concentrations of only 0.1%. Sc acts as a solid solution strengthening element and forms Al3 Sc precipitates. In the present study, we investigated changes in the alloys’ mechanical properties, depending on the heat treatment modes.
A TESCAN Mira3 SEM (operated at 15 kV) with Energy Dispersive Spectrometry (EDS) detector was used for the structural evaluation. The Philips CM20 (operated at 200 kV) was used for TEM investigations. The tensile tests were performed using the Instron 5967. A crosshead speed was set at 1 mm/min (initial strain rate was 8.3 × 10−4 1/s). A rectangular cross-section specimens [24,28] (25 mm gauge length and 4 mm thickness) according to the ASTM E8_E8M_13a were used. The hardness (HB) was measured using a Brinell tester with 306.56 N load, 2.5 mm ball diameter, during 10 s. The average values of hardness and tensile properties for each alloy were taken from 10 measurements. 3. Results 3.1. Selection of heat treatment temper for the studied alloys Fig. 1 shows a comparison of the microstructure (a,b) and the surface (с) for the base Al–5.7Mg–2.6Si–Mn alloy, with different solution treatment temperatures. One can see the temperatures higher than 540°С are inapplicable for the heat treatment of the studied HPDC-parts. However, such low solution temperatures (lower than 520°С) for the studied alloy system do not lead to a maximum of hardness, but in a combination of further overaging leads to significant improvement of ductility levels [29]. Therefore, we choose the several temperatures of the ST followed by AA (T6 regime) for the preliminary experiments. Fig. 2 shows the hardness measurement results across the plate after T6 regime at different ST temperatures for the base alloy. Sample after T6 regime (ST at 570 °C+AA at 175 °C) showed the maximum hardness of 115HV. Hardness values of the samples treated at lower temperatures have not even reached the hardness value of the as-cast state. Table 3 shows the literature data of the applied temperatures for heat treatment of the binary Al-alloys compared to the base alloy. From the thermodynamic calculations [16], we determined the required temperatures of the solution treatment. Table 4 contains the values of solubility of the chosen alloying elements in the current system.
2. Materials and methods Table 1 represents the composition of the studied alloys. The amount of alloying elements was chosen according to the literature data [17,26,27]. To perform thermodynamic calculations, a ThermoCalc software (TC) with the TCAl2:Al-alloys v2.1 database was used to predict the effects of the studied system’s selected elements. Еhe commercial Magsimal59○R ("Rheinfelden Alloys GmbH", Germany) and commercial AlSc2, AlZr15 and AlCr20 master alloys were used as starting materials. The starting materials’ blocks were stored in the preheated furnace for 24 h to avoid moisture before melting. The melt was impelled to minimize microporosity. The density was controlled by a MK 3VT vacuum density tester in combination with MK3000 the density index balance. A HPDC unit with the cold chamber (Frech DAK 450–54 (with 4500 kN mold clamping force) was used to produce 2
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Fig. 1. Microstructure (a,b) and the surface (с) for the base Al–5.7Mg–2.6Si–Mn alloy, with different solution treatment temperatures: a) 520 °C, b) 575 °C. Pores are marked with arrows.
Table 4 The solubility of chosen elements in the Al–5,7Mg,−2,5Si–0,6Mn–X systems based on calculated diagrams. Alloying element Zr Sc Cr
Solubility at: [wt%] 591 °C
540 °C
520 °C
500 °C
480 °C
460 °C
0.07 0.12 0.35
0.05 0.09 0.25
0.04 0.07 0.21
0.03 0.05 0.17
0.02 0.03 0.12
0.01 0.02 0.1
Table 5 Average size of separate Mg2 Si particles as a function of treatment temperature (treatment time – 1.5 h). Temperature °C Diameter, 𝜇m
Table 3 Temperatures for treatment of the alloys doped by chosen elements [°C] [17,19,30–32].
Al–Sc Al–Zr Al–Cr Base alloy
Aging
500– 600
250–350 – – 170–250
480–575
570 0.82
3.2. Full T6 treatment The changes in the structure and properties of Al–Mg–Si–Mn alloys during heat treatment result from several processes that occur simultaneously during heating. The first process is the spheroidization of eutectic fibrous (Fig. 3). A higher solution-treatment temperature leads to a faster eutectic cells decomposition into separate Mg2 Si particles and faster coarsening of these particles. Table 5 shows the average size of these particles depending on the treatment temperature. During the solution treatment, several processes occur in the solid solution. First of all, high-temperature treatment leads to dissolving all strengthening precipitates (including GP-zones, 𝛽’) that form during natural and artificial aging [3,6,15,22]. However, the equilibrium 𝛽-Mg2 Si phase after solution treatment is still present in the microstructure (Fig. 4a,b). The second process is the formation of Mn-containing
Temper Solution treatment
520 0.66
in the a-Al matrix. Further, the alloys can be aged similarly to the base alloy. It does not make sense to perform solutionizing Sc- and Zrcontaining studied alloys. However, Sc-containing alloys can be aged (with the formation of nanoscale Al3 Sc and Al3 (Sc,Zr) precipitates in the Al-matrix) in the range of temperatures 250–500 °C.
Fig. 2. Hardness across the plate after heat treatment and in as-cast state.
Alloying element
480 0.59
Thus, during heat treatment with 520 °С, the new phases in the studied alloys can be dissolved at 0.07 wt% Sc, 0.21 wt% Cr and 0.04 wt% Zr. Based on the above, Cr-containing series of alloys can be heat treated at 520 °C, followed by quenching, to form Cr-containing dispersoids 3
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Materialia 15 (2021) 100981
Fig. 3. Microstructure of the base alloy in a) as-cast state and b) after solution treatment at 480 °C during 1.5 h; c) after solution treatment at 520 °C during 1.5 h; d) after solution treatment at 570 °C during 1.5 h.
Table 6 Tensile properties of the base alloys after full T6 treatment. Temperature [°C] Solution treatment 570 520 480 As-cast
Artificial aging
UTS [MPa]
175 225 175 225 175 225 [16]
291.3 293.7 280.6 242.0 241.6 228.1 299.8
± ± ± ± ± ± ±
2.8 0.5 8.7 12 2.1 1.0 6.3
YS [MPa]
A [%]
235.8 ± 1.1 201.3 ± 9.3 184.8 ± 2.6 130.0 ± 6.6 98.3 ± 0.8 93.7 ± 0.9 163.9 ± 5.6
3.8 ± 0.4 6.0 ± 2.6 9.5 ± 2.3 11.5 ± 1.3 16.8 ± 3.8 18.0 ± 2.9 8.3 ± 0.5
state. The situation is different with treatment at lower temperatures (480–520 °C). Despite the low values of both tensile strength and yield strength, the elongation values are very high and can reach values of 18%. According to the equilibrium phase diagram [16], Al7 Cr is a refractory phase and can be dissolved at relatively high heat-treatment temperatures. In C1 (0.1 wt% Cr) Al7 Cr phase dissolves at 460 °C; in C2 (0.2 wt% Cr) at 520 °C. However, Fig. 7b shows that heat treatment with a temperature of 520 °C even for 3 h is not enough for the dissolution of the Al7 Cr phase. As it was predicted utilizing Thermo-Calc, Al3 Zr and Al3 Sc phases do not dissolve during applied solution treatment. AlMg2 Si eutectic in the studied alloys after heat treatment changes in the same way as the base alloy (eutectic disintegrated into separate spheres). Heat treatment also does not affect the morphology and distribution of Mn-containing intermetallics. All studied additions do not affect the aging behavior of the 𝛽-type precipitates in the studied alloys. Similarly to the base alloy, solution treatment at 520 °C with further aging doesn’t lead to the formation of high density 𝛽-type precipitates (GP-zones and 𝛽’). However, Cr and Zr affects the formation of dispersoids (Fig. 8) inside a-Al dendrites: Cr-containing alloys after solution treatment have (Mn,Cr)-containing dispersoids, while Zr addition leads to the formation Al3 Zr dispersoids. Fig. 9 represents the results of the hardness and tensile tests of studied alloys with different aging temperatures for the time of maximum hardness values (3 h). The higher aging temperature leads to the lower strength (and hardness) and the higher elongation. In general, the behavior of changes is similar to the base alloy.
dispersoids. Fig. 4c and d shows that dispersoids after high-temperature heat treatment form attaching to spherical Mg2 Si particles and the equilibrium 𝛽-Mg2 Si phase. Fig. 5a and b shows that during aging after solution treatment at the temperature of 520 °C, strengthening phases are practically not formed. However, solution treatment at 570 °C with further aging (Fig. 5c and d) leads to the formation of the solid solution’s high-density GP-II zones. Table 6 and Fig. 6 show the mechanical properties of the base alloy after treatment with different temperatures. Results of the hardness measurements (Fig. 6) show that aging after low-temperature solution treatment (e.g., temperature range 480–520 °C) does not lead to a significant increase in the hardness of the base alloy. On the other hand, aging after solution treatment at 570 °C increases the hardness by 25% compared to the solutionized state and by 12% compared to the as-cast
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Fig. 4. TEM images of the base alloy after solution treatment at 520 °C a,b) and at 570 °C c,d) (without aging).
3.3. Artificial aging from the as-cast state
hardness of the Sc-containing series of alloys, especially S2 alloy (with 0.2 wt% of Sc), increases significantly and reaches values about 120HB. Fig. 15 represents the results of tensile tests of studied alloys for different aging temperatures for the time of maximum hardness values (3 h). Both Cr-containing and Zr-containing alloys show similar changes and similar levels of all tensile properties. Maximum values of strength belong to the aging temperature range 175–225 °C. That coincides with the lowest values of elongation. S1 alloy (0.1 wt% Sc) shows approximately similar results of all tensile tests for all aging temperatures. Aging temperature higher than 125 °C leads to a minor increasing of yield strength. S2 alloy shows the highest values of both tensile and yield strength, especially for aging treatment with 325 °C. Fig. 16 shows the results of tensile properties for aging with 325 °C as a function of the time. The aging time increase reduces both tensile and yield strength simultaneously with increasing elongation for all studied alloys.
Fig. 10a shows changes in the hardness of the base alloy with different temperatures and aging times. Thus, from the diagram, it is clear that the base alloy’s heat treatment does not have a significant effect on the hardness. The most noticeable changes were recorded at the aging temperature of 175 °C and time in the range of 2–3 h. However, the increase in hardness in this mode is only 5–7% higher than in the cast state. Fig. 10b shows the results of tensile properties of the base alloy with different temperatures of aging for time with peak hardness (3 h). One can see the increase in tensile strength is insignificant, but yield strength has more prominent changes. Nevertheless, the elongation is decreased. Results of SEM and TEM investigations of the base alloy do not show any significant changes neither in microstructure nor in the structure of the 𝛼-Al dendrites during treatment in a temperature range of 125– 325 °C. The microstructure consists of 𝛼-Al dendrites, Al-Mg2 Si eutectic colonies, and 𝛼-AlMnFeSi intermetallics. The aging in such temperature range and time does not lead to changes in all phases’ morphology in comparison to as-cast state (Fig. 11). The structure of 𝛼-Al dendrites consists of 𝛽’ and stable 𝛽 phases similar to the as-cast state [24] (Fig. 12). Similar to the base alloy, all studied alloys don’t show any changes in the microstructure after heat treatment in the temperature range of 125– 325 °C. However, alloys with Sc addition show increasing nanoscale precipitates’ density after high-temperature aging treatment (325 °C) (Fig. 13). All studied series of alloys show approximately similar values of hardness after aging with temperatures 125–225 °C (Fig. 14). However, the aging with temperature 325 °C shows different behavior. Thus, the
4. Discussion 4.1. Full T6 heat treatment The changes in the properties of Al-Mg-Si-Mn alloys during hightemperature solution treatment (as the first stage of the full T6 treatment) result from the changes of structural components that occur during heating. Since Mg2 Si eutectic is the main strengthening phase in this type of alloys, the process of its disintegration reduces the hardness and strength that coincides with the increase of elongation. Trudonoshyn et al. [15,28] reported a similar tendency for the casting Al-Mg-Si-Mn alloys cast into the permanent mold. Together with eutectic disinte5
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Fig. 5. TEM images of the base alloy after solution treatment at 520 °C a,b) and at 570 °C c,d) and further aging at 175 °C (solution treatment time – 1.5 h; aging time – 3 h).
higher hardness values of the HPDC compared to the regular permanent mold casting alloys [15,28]. Artificial aging after solution treatment causes the formation of the strengthening precipitates again [15] that improve the strengthening properties of the alloys. Nevertheless, different solution treatment temperatures have different effects. Aging after solution treatment at 570 °C has a more noticeable effect on the hardness. This may be due to the formation of a larger number of quench-in vacancies that promote the diffusion of the solute atoms, which in turn, promote the formation of atoms clusters and GP-zones [36–39]. However, such heat treatment is not recommended by the manufacturer [14]. The different situation is with treatment at lower temperatures (480–520 °C). Thus, despite the low values of both tensile strength and yield strength, the elongation values are very high and can reach values of 18%, as confirmed by the results presented in [29]. All studied series of alloys after full T6 treatment have the same microstructural changes as the base alloy. Moreover, solution treatment at 520 °C during 1.5 h (and even 3 h) does not lead to dissolution of the Al3 Zr and Al3 Sc phases in Zr- and Sc-containing alloys respectively (as it was predicted by Thermo-Calc software) and even the Al7 Cr phase in Cr-containing alloys (despite been predicted by Thermo-Calc software). A primary cause of this is the low diffusion rate of Cr. This phase’s dissolution could improve the tensile properties of the alloy, but this did not happen. Moreover, Samuel et al. [40] reported that sludge phases remained stable after solution treatment for 8 h at 515 °C. Solute Cr in 𝛼-Al dendrites during solution treatment promotes the formation of (Mn,Cr)-containing dispersoids, which, however, do not significantly affect the mechanical properties of the casting alloys. In
Fig. 6. Brinell hardness of the base alloys after full T6 treatment with different solution treatment and aging temperatures and different treatment times.
gration, the solution treatment usually leads to the dissolution of the strengthening 𝛽’’-Mg2 Si precipitates that might be formed in the as-cast state of Al-Mg-Si alloys after solidification [33,34]. However, due to the specific Mg/Si ratio of the base alloy, the amount of the strengthening 𝛽’’-Mg2 Si precipitates in as-cast (and even after aging [21,22]) is negligible. Solution treatment of the Al5Mg2SiMn permanent mold casting leads to the formation of 𝛽-AlMnFeSi dispersoids [15,28]. Their lack of coherence with the matrix leads to a reduction in the alloys’ hardness still more (along with decomposition of eutectic lamellas). On the other hand, in the studied HPDC Al-Mg-Si-Mn alloys, we observed 𝛼-AlMnFeSi dispersoids. This type of dispersoids is semi-coherent with the Al-matrix [35] that can slightly affect the hardness of the alloys. This explains the 6
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Fig. 7. SEM images of Z2 (a), C2 (b), and S2 (c) alloys after solution treatment at 520 °C (solution treatment time – 3 h).
Fig. 8. TEM images of the Mn,Cr-containing dispersoids in C2 alloy (a) and Al3 Zr dispersoids in Z2 alloy after solution treatment at 520 °C for 1.5 h.
contrast to the Zr-series of alloys which have the highest values of hardness and strength after full T6 treatment among all studied alloys. According to [31], this effect is caused by the formation of the Al3 Zr dispersoids in 𝛼-Al dendrites’ structure. Except for the high cost, Sc has another weakness as an alloying element for Al-Mg-Si HPDC alloys. Since the Al3 Sc phase does not dissolve at the applied solution treatment temperatures (Table 6), it makes no sense to perform full T6 treatment on the investigated Sc-containing alloys. In contrast to the situation with binary Al–Sc alloys [31], which can be treated using temperatures above the solvus (a precondition for age strengthening is that the strengthening phase must be soluble between solvus and solidus temperatures) and thus, get a significant increase in the mechanical properties. Despite the good mechanical properties of the Zr-containing series, compared to Cr- and Sc-containing series after full T6 treatment, the strength values of all studied alloys are still lower than in the as-cast state, which suggests that the eutectic is the more strengthening phase in the Al-Mg-Si casting alloys.
and time in the range of 2–3 h. However, the increase in hardness in this mode is only 5–7% higher than the as-cast state. Such an insignificant effect is concerned with the fact that the base alloy belongs to alloys with an excess of Mg content. Mg drastically reduces the solubility of Mg2 Si in the 𝛼-Al, which in turn reduces the hardening potential (by the formation of nanoscale precipitates) and also increases the stability of the 𝛽’ and equilibrium 𝛽 phases (which do not have a significant hardening effect) [21,22]. In general, it can be argued that excess of Mg makes the properties of the base alloy very stable at elevated temperatures (which makes the material promising for further research at elevated temperatures). The microstructural investigation confirms the results of the mechanical properties, which shows that the aging treatment in the studied temperature range (125–325 °C) and time (up to 15 h) does at least not change the morphology of the phases (Fig. 11). The structure of 𝛼-Al dendrites mostly consists of 𝛽’ and stable 𝛽 phases similar to the as-cast state [23,24]. However, in some 𝛼-Al dendrites, only small amounts of 𝛽"-precipitates were found after aging at 175 °C (Fig. 12c). Similar to the base alloy, alloys additionally doped by Cr and Zr have changes neither in microstructure nor in precipitation sequence after heat treatment in the temperature range of 225–325 °C. We did not observe any new Cr- or Zr-containing precipitates. Both series of alloys show similar to the base alloy behavior of all studied mechanical properties after aging (Fig. 14, Fig. 15). Despite this, aging from the as-
4.2. Artificial aging from as-cast state The artificial aging of the base alloy at the temperature range 125– 325 °C does not significantly affect the hardness (Fig. 10). The most noticeable changes were recorded at the aging temperature of 175 °C 7
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Fig. 9. Mechanical properties of the studied alloys after full T6 heat treatment (solution treatment at 520 °C during 1.5 h, artificial aging during 3 h).
cast state remains promising for Al-Mg-Si-Mn alloys additionally alloyed by other elements (e.g., Sc) [41]. Aging temperatures for the formation of nano-dispersed Al3 Sc precipitates are significantly higher than the temperatures used to obtain maximum hardness values in Al–Mg–Si alloys by the formation of 𝛽’’-
Mg2 Si precipitates [17]. Such temperatures (250–350°С) lead to overaging and make more stable 𝛽’- and equilibrium 𝛽-Mg2 Si phases (which do not significantly affect the strength). On the other hand, as already mentioned, the base alloy has rather stable properties and structure during aging up to 325°С. Aging of Sc-containing alloys at a temperature
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Fig. 10. Brinell hardness of the base alloy as a function of aging temperature and time (a) Tensile properties of the base alloy as a function of aging temperature (3 h) (b).
Fig. 11. SEM-images of the microstructure of the base alloy: a) as-cast state; b) aging at 175 °C; c) aging at and 325 °C.
Fig. 12. TEM bright-field images of the base alloy after artificial aging: a–c) base alloy after aging at 175 °C; d–f) base alloy after aging at 325 °C.
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Fig. 13. TEM bright-field images of the Sc-containing alloys after artificial aging:a) S1 after 225 °C; b) S1 325 °C; c) S2 225 °C; d) S2 325 °C;.
Fig. 14. Brinell hardness of the studied alloys as a function of aging temperature and time: a) Z1 (0.1 wt% Zr); b) Z2 (0.2 wt% Zr); c) C1 (0.1 wt% Cr); d) C2 (0.2 wt% Cr); e) S1 (0.1 wt% Sc); f) S2 (0.2 wt% Sc). 10
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Fig. 15. Tensile properties of the studied alloys as a function of concentration of alloying elements for peak-hardness aging time: a) Z1; b) Z2; c) C1; d) C2; e) S1; f) S2.
Fig. 16. Tensile properties of the studied alloys as a function of the aging time for 325 °C.
325 °С from the as-cast state is a promising heat treatment. The strength of S2 alloy (with 0,2 wt% Sc) increased up to 380 MPa (11% higher than the as-cast state of S2 and 30% higher than base alloy). The formation of nano-dispersed precipitates Al3Sc can only explain such a significant increase in hardness and strength. Such an increase in properties is in good agreement with the existing studies concerning the alloying of Alalloys by Sc [31,32,41].
implementing the commercial analogs of the Cr-containing and Zrcontaining alloys in the production (Maxxalloy-Ultra, Magsimal Plus). Sc has the most significant influence on the base composition in the as-cast state among all elements. Subsequent artificial aging from the as-cast state increases the strength of the alloys, but it decreases elongation. The reasons for the change in properties in the studied alloys were the formation of new primary phases in the as-cast structure and the formation of nanosized precipitates during the heat treatment. It was found that coarse phases (such as Al7 Cr in Cr-containing alloys and Al3 Zr in Zrcontaining alloys) are brittle and significantly reduce the tensile properties of alloys, especially ductility. The Al3 Sc phase (Sc-containing alloys) did not negatively affect the ductility of the alloys. The primary Al7 Cr, Al3 Zr, Al3 Sc phase doesn’t dissolve even during high-temperature solution treatment. However, Sc that remains in 𝛼-Al during solidification
5. Conclusions This study represents the microstructural evolution and Al-Mg-Si-Mn alloys’ mechanical properties with additional alloying by Cr, Zr, and Sc in different heat treatment tempers. The analysis of all mechanical properties of the studied alloys shows that Cr and Zr have insignifficant effect in all tempers, despite 11
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is sufficient for the formation of a nanoscale strengthening precipitates during aging. As expected, base alloy showed a good response to additional alloying with elements that enhance the precipitation strengthening. Thus, the greatest impact on the mechanical properties was affected by nanoscale Al3 Sc strengthening precipitates formed in 𝛼-Al solid solution in Sc-containing alloys during aging. Although the Cr addition to the base alloy leads to the formation of a dispersed Cr-containing phase in 𝛼-Al, it does not have a significant strengthening effect (similar to Cr-containing wrought alloys).
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Declaration of Competing Interest The authors declare no conflict of interest. Acknowledgments O. Prach and O. Trudonoshyn are grateful to German Academic Exchange Service (DAAD) for the financial support over this research. Supplementary materials Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.mtla.2020.100981. References [1] S. Babaniaris, M. Ramajayam, L. Jiang, T. Langan, T. Dorin, Tailored precipitation route for the effective utilisation of Sc and Zr in an Al-Mg-Si alloy, Materialia 10 (2020) 100656, doi:10.1016/j.mtla.2020.100656. [2] O. Trudonoshyn, O. Prach, V. Boyko, M. Puchnin, K. Mykhalenkov, Design of a new casting alloys containing Li or Ti+Zr and optimization of its heat treatment, in: Proceedings of the 23rd International Conference on Metallurgy and Materials, METAL 2014, 2014, pp. 1399–1404. http://metal2014.tanger.cz/files/proceedings/17/reports/2750.pdf. [3] T. Dorin, M. Ramajayam, S. Babaniaris, L. Jiang, T.J. Langan, Precipitation sequence in Al–Mg–Si–Sc–Zr alloys during isochronal aging, Materialia 8 (2019) 100437, doi:10.1016/j.mtla.2019.100437. [4] O. Trudonoshyn, M. Puchnin, O. Prach, Use of the ABI technique to measure the mechanical properties of aluminium alloys: effect of heat-treatment conditions on the mechanical properties of alloys, Mater. Tehnol. 50 (2016), doi:10.17222/mit.2014.295. [5] R.N. Lumley, R.G. Odonnell, D.R. Gunasegaram, M. Givord, Heat treatment of high-pressure die castings, Metall. Mater. Trans. A 38 (2007) 2564–2574, doi:10.1007/s11661-007-9285-4. [6] F. Yan, W. Yang, S. Ji, Z. Fan, Effect of solutionising and ageing on the microstructure and mechanical properties of a high strength die-cast Al-Mg-Zn-Si alloy, Mater. Chem. Phys. (2015) 1–9, doi:10.1016/j.matchemphys.2015.10.014. [7] C. Li, Y. Wu, H. Li, Y. Wu, X. Liu, Effect of Ni on eutectic structural evolution in hypereutectic Al–Mg 2 Si cast alloys, Mater. Sci. Eng. A 528 (2010) 573–577, doi:10.1016/j.msea.2010.09.056. [8] L. Lodgaard, N. Ryum, Precipitation of chromium containing dispersoids in Al–Mg–Si alloys, Mater. Sci. Technol. 16 (2000) 599–604, doi:10.1179/026708300101508315. [9] L. Lodgaard, N. Ryum, Precipitation of dispersoids containing Mn and / or Cr in Al–Mg–Si alloys, Mater. Sci. Eng. A 283 (2000) 144–152, doi:10.1016/S0921-5093(00)00734-6. [10] K.E. Knipling, D.C. Dunand, D.N. Seidman, Criteria for developing castable, creep-resistant aluminum-based alloys – a review, Z. Metallkd 97 (2006) 246–265. [11] O. Prach, J. Hornik, K. Mykhalenkov, Effect of the addition of Li on the structure and mechanical properties of hypoeutectic Al-Mg2Si alloys, Acta Polytech. 55 (2015) 253–259, doi:10.14311/AP.2015.55.0253. [12] O. Trudonoshyn, Studying the structure of Al-Mg-Si casting alloys doped by lithium, Phys. Met. Metallogr. 121 (2020) 701–707, doi:10.1134/S0031918X2007011X. [13] S. Riva, K.V. Yusenko, N.P. Lavery, D.J. Jarvis, S.G.R. Brown, K.V. Yusenko, N.P. Lavery, D.J. Jarvis, S.G.R. Brown, S. Riva, K.V. Yusenko, N.P. Lavery, D.J. Jarvis, S.G.R. Brown, The Scandium Effect in Multicomponent Alloys the Scandium Effect in Multicomponent Alloys, 6608 (2016). doi:10.1080/09506608.2015.1137692. [14] Rheinfelden Alloys GmbH & Co. KG, Ductile aluminum alloys for automotive structural applications, Rheinfelden (2017) http://rheinfelden-alloys. eu/wp-content/uploads/2015/04/Handbook-Alloys-for-Structural-Applications_ RHEINFELDEN-ALLOYS_2017_EN.pdf.
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