Recently a considerable amount of research has been devoted to analyzing the effect of adding yttrium and erbium on the structure and properties of pure aluminum Al [1,2,3,4,5,6,7,8,9,10,11,12,13,14] and alloys based on the Al–Mg system [15,16,17,18,19,20,21,22,23]. Small additions of erbium and yttrium improve the mechanical properties of aluminum and its alloys as a result of forming dispersoids of the Al3M type during annealing [2, 4,5,6,7]. In the presence of zirconium and scandium the strengthening effect during cast alloy annealing increases due to forming dispersoids Al3(ZrY) [5, 7, 10, 11], Al3(ErZr) [2, 4], and Al3(ScZrEr) [8]. In alloy Al–0.4% at.% Er the hardness reached during annealing is 400 MPa due to precipitation of dispersed phase Al3Er [2]. On adding 0.08 at.% Zr to this alloy hardness increases to 560 MPa [2]. In alloy of the Al–Zr–Sc–Si system replacement of 0.01% at.% scandium by erbium led to an increase in creep properties after two-stage annealing by a regime: 300 °С, 4 h and 425 °С, 8 h [3]. The authors in [6] showed that the maximum strengthening in alloy Al–0.045 at.% Er is achieved with annealing in the temperature range 300–375 °C. It has been established in [8] that addition of silicon into Al–Sc–Zr–Er leads to strengthening during ingot annealing due to an increase in the density of Al3(ScZrEr) phase precipitates. Addition of yttrium to alloys of the Al–Zr system leads to acceleration of Al3Zr dispersoid separation kinetics and formation of Al3(Zr,Y) particles of reduced size [5, 7, 10, 11]. It has been established in [12] that presence of Fe and Ni impurities in Al–0.2%Zr–0.1%Sc alloy during crystallization leads to partial bonding of scandium in phases (AlFeSiSc) and (AlFeSc) and correspondingly impoverishment of them for aluminium solid solution. Addition of 0.2% yttrium to this alloy (with iron and silicon impurities) leads to formation of Al3Y and (AlYFeSi)/Al10Fe2Y phases, and in this case scandium is entirely dissolved in aluminum solid solution [12]. For alloys of the Al–Mg system additions of erbium and yttrium are effective grain modifiers and strengtheners [15,16,17,18,19,20,21,22,23].

The aim of this work is determination of the effect of iron and silicon impurities on the phase composition of alloys in a cast condition and their mechanical properties in a deformed state with the treatment conditions adopted. Alloys contained simultaneously yttrium and erbium, zirconium, and scandium, and were distinguished by an overall content of iron and silicon impurities: alloy with 0.01$ and alloy 2 with 0.15%.

Aluminum alloys were melted in a Nabertherm electric resistance furnace based on aluminum A99, i.e., alloy 1, and A85, i.e., alloy 2 (Table 1) using Al–1.8%Er, Al–1.3%Y, Al–3.5%Zr, and Al–2.0%Sc master alloys. Master alloys Al–1.8%Er, Al–1.3%Y, Al–3.5%Zr, and Al–2.0%Sc were melted from pure metals using aluminum A99, whereas the most economic method for preparing master alloy is use of erbium and yttrium compounds [24, 25]. Melting and casting was conducted at 750°C. Alloys casting was accomplished in a copper water-cooled mold with a size of 20 × 40 × 100 mm, and the cooling rate during crystallization was ≈ 15 K/min.

Table 1 Test Aluminum Alloy Chemical Composition
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Heat treatment was performed in Nabertherm and SNOL drying cabinets with an accuracy of temperature maintenance of ± 1 °C. Deformation treatment of ingots was carried in a USP-133 rolling mill with a roll diameter of 300 mm to a thickness of 10 mm at 440 °C and 1 mm at room temperature.

Hardness was measured by s standard Vickers method. Tensile testing was carried out using a Zwick/Roell Z250 Allround test machine.

Preparation of microsections for microstructural studies was performed in a Struers Labopol-5 grinding and polishing unit. Microstructural studies and phase identification was conducted in a Neophot 30 light microscope (SM) and a TESCAN VEGA 3LMH scanning electron microscope (SEM) using an X-Max 80 energy dispersion detector.

Experimental Results and Discussion

The microstructure is given in Fig. 1 of test alloys in a cast condition. Within the structure using an SEM apart from aluminum solid solution light inclusions of excess phase were observed, formed during crystallization (Fig. 1a, b). TEM images are provided in Fig. 1c, d of the test alloy cast structure. Light inclusions with assize of ≈ 2 μm (see Fig. 1a, b) are represented by eutectic colonies, within which the size of phases is 200–800 nm. Images in characteristic radiation for elements in distribution maps show (Fig. 2) that within the structure of alloy 1 light phases are enriched with Y, Er, and Zr (see Fig. 2a), and for alloy 2 presence of a phase is noted formed by impurities of iron and silicon (see Fig. 2b). In phases of alloy 2, containing iron and silicon impurities, an increase in Zr, Er, and Y concentration is observed, but in this case yttrium does not form phases with iron (see Fig. 2b). In a cast condition alloy 2 forms multicomponent phases (AlSiFeEr(Zr)) and (Al-SiErY(Zr)). As a result there is a reduction in supersaturation of aluminum solid solution in Zr, Er, and Y. Scandium in both alloys is uniformly distributed within the aluminum solid solution matrix and does not enter into a single phase formed during crystallization.

Fig. 1.
figure 1

Alloy microstructure in cast condition: (a), (b) SEM; (c), (d) TEM; (a), (c) alloy 1; (b), (d) alloy 2.

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Fig. 2.
figure 2

Alloying element distribution maps in test metal phases: (a) alloy 1; (b) alloy 2.

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In cast alloy 2 during annealing a weaker strengthening effect is observed compared with alloy 1, i.e., hardness increases by 7–9 HV (Fig. 3b), and in this case the absolute hardness value does not exceed 35 HV. At the same time in cast alloy 1 a maximum value of hardness (40 HV) is achieved with an annealing temperature of 440 °C (Fig. 3a). The small strengthening effect for alloy 2 is connected with the lower zirconium, erbium, and yttrium concentrations in aluminum solid solution.

Fig. 3.
figure 3

Change in hardness values in relation to metal temperature and annealing time in a cast condition: (a) alloy 1; (b) alloy 2.

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Results are provided in Fig. 4 for a study of the structure by light microscopy of material annealed at 450, 500, and 550 °C for 1 h of sheet 1 mm thick prepared from the test aluminum alloys. For alloy 1 the temperature for the start of recrystallization is in the range 500–550 °C. After annealing at 550 °C for 1 h the structure is entirely recrystallized (see Fig. 4h), whereas at 500 °C deformed grains are revealed (see Fig. 4f). For alloy 2 the temperature for the start of recrystallization somewhat higher, but it is found in the same range since on holding at 500 °C for 1 h a partly recrystallized structure is observed (see Fig. 4g).

Fig. 4.
figure 4

Sheet macrostructure (light microscope) in original (a, b) and annealed at 450 (c, d), 500 (e, f), and 550 °C (h, i) conditions: (a), (c), (e), (g) alloy 2; (b), (d), (f), (h) alloy 1.

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Results are provided in Fig. 5 for measurement of hardness HV for test alloy sheet after annealing, performed at 150 and 200 °C for 0.5, 1, 2, 3, 4, and 8 h. Especially noted is an increase in hardness for both alloys after annealing at 200 °C for 1 h, from 48 HV to 53 HV. Accumulation of defects and internal stresses during subsequent heat treatment could lead to activation of solid solution breakdown and correspondingly an increase in hardness an increase in holding time during annealing, performed at 200 °C, caused formation of a polygonized structure and reduction in hardness to the level of the original condition (with holding for 8 h). Annealing at 150 °C did not lead to such a marked increase in hardness since for breakdown of the solid solution action of a high temperature was required, and also at this temperature to a lesser extent there are recovery ad polygonization processes. In this case material strength should not decrease markedly, and ductility should increase.

Fig. 5.
figure 5

Effect of annealing temperature and time on sheet hardness level: (a) 150 °C; (b) 200 °C.

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Results are given in Table 2 for determination of standard mechanical properties (nominal yield strength (σ0.2), ultimate strength (σf), and relative elongation (δ)) for sheets of aluminum alloys in the original and annealed conditions. The results obtained for determining ultimate strength correlate well with hardness measurement results.

Table 2 Effect of Annealing Regime on Sheet Mechanical Properties
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Tests sheets in the original condition had σ0.2 = 155 MPa, σf = 166 MPa, δ = 8.7% (for alloy 1) and 11.2% (for alloy 2). After annealing at 100 °C for 3 h the ultimate and yield strengths remained at the same level, but relative elongation for alloy 2 decreased by 10%, whereas for alloy 1 it increased to 12.8%. The set of standard mechanical properties for alloys 1 and 2 after annealingat200 °C for 4 h are at the same level: σ0.2 = 151 ± 1 MPa and σf = 164 ± 1 MPa, δ = 12.3 ± 0.1 for both alloys.

Conclusions

  1. 1.

    In alloy of the Al–Y–Er–Zr–Sc system with an overall Fe and Si impurity content of 0.15% presence of phases is noted formed by iron and silicon impurities with an increased concentration of zirconium, erbium, and yttrium, whereas with content in the alloy of 0.01% of these impurities there is no additional phase formation.

  2. 2.

    Alloy of the Al–Y–Er–Zr–Sc with 0.15% (Fe + Si) system in a cast condition shows a weaker effect if strengthening on annealing, which is connected with the lower supersaturation of aluminum solid solution with zirconium, erbium, and yttrium.

  3. 3.

    The temperature for the start of test alloy recrystallization is in the range 500–550 °C, and in this case for alloy with 0.15% (Fe + Si) it is somewhat higher.

  4. 4.

    Mechanical properties of sheet after annealing in the range 100–300 °C are at the same level: the ultimate and yield strengths are 152–157 MPa and 163–168 MPa respectively.