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Article

Effect of Fe-Bearing Phases on the Mechanical Properties and Fracture Mechanism of Al–2wt.%Cu–1.5wt.%Mn (Mg,Zn) Non-Heat Treatable Sheet Alloy

by
Nikolay Belov
1,
Torgom Akopyan
1,*,
Kirill Tsydenov
1,2,
Stanislav Cherkasov
1 and
Natalia Avxentieva
1
1
Department of Metal Forming, National University of Science and Technology MISiS, 4 Leninsky Pr., Moscow 119049, Russia
2
Laboratory of Mechanics of Gradient Nanomaterials, Nosov Magnitogorsk State Technical University, 38 Lenin pr., Magnitogorsk 455000, Russia
*
Author to whom correspondence should be addressed.
Metals 2023, 13(11), 1911; https://doi.org/10.3390/met13111911
Submission received: 24 October 2023 / Revised: 17 November 2023 / Accepted: 17 November 2023 / Published: 20 November 2023
(This article belongs to the Special Issue Feature Papers in Metal Failure Analysis)
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Abstract

:
The effects of Fe-bearing phases on the structure, mechanical properties, and fracture mechanism of a non-heat-treatable model sheet alloy (wt.%: Al–2%Cu–1.5%Mn(-Mg,Zn)), designed for Al20Cu2Mn3 dispersoids, was investigated. This involved a combination of thermodynamic modeling in the Thermo-Calc program and experimental studies of structure and mechanical properties. It has been shown that the addition of 0.5 and 0.4% iron and silicon leads to the formation of eutectic inclusions in the Al15(Mn,Fe)3Si2 phase. In addition to the Fe- bearing inclusions, the formation of the eutectic Al2Cu and Al2CuMg phases can be expected in the as-cast structure of the experimental alloys. Despite their relatively high fraction of eutectic particles, non-homogenized alloy ingots demonstrated sufficiently high deformation processability during the hot (400 °C) and cold rolling, which made it possible to obtain high-quality sheet alloys (with reduction degrees of 80 and 75%, respectively). The results of the tensile tests revealed that, after cold rolling, the addition of 1% Mg significantly increased the tensile and yield strengths, whereas the effect of 1% Zn was negligible. At the same time, the uniform distribution of Fe-bearing phases in the structure of the cold-rolled sheets contributes to the preservation of the dimple mechanism of the fracture toughness. This helps to maintain the same level of ductility for the cold-rolled sheet Fe-containing alloys as for Fe-free alloys. It has been shown, based on the data obtained, that adding Fe, Si, Mg, and Zn to the base Al–2%Cu–1.5%Mn alloy in a total amount of more than 3% makes it possible to retain the ductile fracture patterns of the base alloy and obtain a fairly higher level of mechanical properties. This suggests the fundamental possibility of using a variety of secondary raw materials (containing the main elements present in aluminum alloys of different alloying systems) to prepare a base alloy that does not require homogenization or thermal hardening.

1. Introduction

The production and consumption of aluminum alloys are continuously increasing because of the unique combination of their operational and technological properties with an abundance of raw materials [1,2,3,4,5,6,7,8]. From the viewpoint of environmental friendliness and economy, the production of aluminum alloys using recycled materials as the main charge is much more attractive than the use of primary aluminum. It should be noted that primary aluminum is more expensive than scrap aluminum, and its production is more harmful to the environment [9,10].
The secondary aluminum raw materials contained all the main constituents (Cu, Si, Mg, Zn, and Mn) of the branded alloys. The latter fact makes it possible to reduce the amount of primary additives and alloy preparation costs (including reduction in the melting time) [11,12,13]. On the other hand, aluminum scrap contains elevated concentrations of impurities (in particular, Fe), thus preventing the production of alloys with strict impurity limitations [14,15,16,17]. One possible solution is to design new alloys for which high–quality structures and operational properties can be achieved with high impurity contents [15,16,17].
Earlier works [18,19,20,21,22,23] dealt with the possibility of creating new–generation high-tech wrought Al-Cu-Mn system based aluminum alloys with elevated strength and heat resistance. These properties are achieved due to the presence of Al20Cu2Mn3 compound dispersoids in the structure of the alloys in amounts of 7–8 vol.%. Based on the calculated and experimental data obtained, the optimal copper (1.5–2%) and manganese (1.5–2%) concentrations were determined, which provided the best combination of manufacturability and physical and mechanical properties.
Taking into account the fact that scrap (especially mixed) contains, along with copper and manganese, other impurities, like iron, silicon, magnesium, and zinc, it is advisable to continue these studies, focusing on expanding the raw material base [24,25,26,27]. At the same time, it should be noted that iron, unlike other elements, is practically insoluble in an aluminum solid solution (hereinafter (Al)) and forms Fe-bearing phase inclusions, the morphology of which significantly affects the mechanical properties.
In aluminum alloys, iron forms various phases (including those with Mn, Si, Cu, and Mg), most of which have an unfavorable morphology (particularly needle-like or rod-like) [28,29,30]. As stress concentrators, they contribute to the premature initiation of microcracks during loading and, as a result, a decrease in mechanical properties, especially ductility, fracture toughness, and fatigue characteristics [31,32,33]. The most desirable phases are those whose eutectic inclusions have skeletal shapes, e.g., Al8Fe2Si [34,35] and Al15(Mn,Fe)3Si2 [36,37,38,39]. The former is typical of 6xxx series alloys because they are prone to spheroidization upon homogenization annealing [40,41,42,43]. Since the Al–2wt.%Cu–1.5wt.%Mn non-heat treatable sheet alloy was designed for a technological process that does not include homogenization annealing, this morphology improvement method is indispensable. Furthermore, with such a high manganese content, the formation of the Al15(Mn,Fe)3Si2 phase is more likely, as shown earlier [23]. However, it should be taken into account that the combined influence of the Al–Cu–Mn–Fe–Si–Mg–Zn alloying system elements, which are the basic elements for recycled alloys, has not previously been analyzed for the considered concentrations of copper and manganese.
Based on this, the aim of this work is to clarify the effect of Fe-bearing phases on the mechanical properties and fracture mechanism of the Al–2wt.%Cu–1.5wt.%Mn non-heat treatable sheet alloy containing various combinations of potential impurities, i.e., Fe, Si, Mg, and Zn.

2. Experimental Section

The test materials were 8 model alloys containing 2% Cu and 1.5% Mn: 4 alloys without Fe and Si additions (B, BM, BZ, and BMZ) and 4 alloys with 0.5% Fe and 0.4% Si (BF, BFM, BFZ, and BFMZ). The two alloys were without Mg and Zn additions (B and BF), and 1% Mg and 1% Zn were added separately and together to the other alloys. All these alloys (they are summarized in Table 1) constitute 2 series (Series B: small Fe and Si amounts and Series BF: high Fe and Si amounts), allowing the study of the effect of Fe-bearing phases on the mechanical properties and fracture mechanism of the Al–2wt.%Cu–1.5wt.%Mn (Mg,Zn) non-heat treatable sheet alloy.
The experimental alloys were prepared in a resistance furnace (GRAFICARBO) in a graphite-chased crucible. Pure raw metals (99.85% aluminum, 99.9% copper, 99.9% magnesium, and 99.9% zinc) and master alloys (Al–10%Mn, Al–10%Fe and Al–12%Si) were used to obtain the selected compositions. When the material batch melted, it was held for approximately 10 min for homogenization and then poured at 750 °C into a flat graphite mold 10 × 40 × 180 mm in size. The cooling rate during the solidification was approximately 20 K/s. The chemical composition of the experimental alloys, according to spectral analysis (ARL 3460 (Thermo Fisher Scientific, Waltham, Massachusetts, USA)), is presented in Table 1. It can be seen that the actual compositions were close to the target ones.
The ingots of experimental alloys were subjected to hot rolling at 400 °C. Before rolling, the ingots were annealed at 400 °C for 1 h. The ingots were then rolled to a thickness of 2 mm (80% reduction rate). Cold rolling was carried out to a thickness of 0.5 mm (75% reduction rate). The cold-rolled sheet products were prepared using a laboratory-scale rolling mill machine (Chinetti LM160 (Chinetti snc, Cavaria con Premezzo (Varese) Italy)). The final step involved annealing the samples at 400 °C for 3 h. The parameters of the experimental alloy processing methods are presented in Table 2.
The microstructure was examined by optical microscopy (OM, Axio Observer MAT), scanning electron microscopy (SEM, TESCAN VEGA 3, TESCAN GROUP, Brno-Kohoutovice, Czech Republic), electron microprobe analysis (EMPA, OXFORD Aztec, Milton Keynes, UK) and transmission electron microscopy (TEM, JEM-2100, JEOL Ltd., Tokyo, Japan). The specimens were prepared using mechanical and electrolytic polishing. Electrolytic polishing was performed at a voltage of 12 V in an electrolyte (6 C2H5OH, 1 HClO4, and 1 glycerine). The initial microstructural observations were carried out using OM, and detailed studies were performed using SEM and TEM. Thin foils for TEM were prepared via ion thinning using the Precision Ion Polishing System (PIPS, Gatan, Pleasanton, USA) and studied at 160 kV [44].
Room-temperature tensile tests were conducted on the cold–rolled sheets using an Instron 5966 machine. The loading rate was set to 10 mm/min. 4 specimens (120 × 10 × 0.5 mm) were used for each alloy.
The phase composition of the Al–Cu–Mn–Mg–Zn–Fe–Si system (isothermal sections, phase fractions, and element concentrations in (Al)) was calculated using Thermo-Calc software (TTAL5 database [45], Thermo-Calc Software AB Råsundavägen, Solna, Sweden).

3. Experimental Results

3.1. Microstructureof As-Cast Ingots

Initially, the joint effects of Fe, Si, Mg, and Zn on the phase composition of the base alloy were investigated. As observed from the isothermal sections, calculated at 400 °C (Figure 1), these additives significantly complicated the phase composition. When Fe was added separately, an Al7Cu2Fe compound was first formed, followed by Al6(Fe,Mn) (Figure 1a). The separate addition of silicon should lead to the appearance of the Al15Mn3Si2 compound, and for the combined addition of Fe and Si, the Al15(Mn,Fe)3Si2 phase (which is a solution of iron in the same ternary compound) is expected. In the presence of 1% Mg and 1% Zn, the Mg2Si and Al2CuMg (S) phases were added to the above-mentioned ones. All these phases can be present in graded aluminum alloys and have been studied quite well [46,47,48,49,50,51,52,53,54].
The analysis of the microstructure of the ingots (10F) focused on the morphology of the excess phases (Figure 2), their identification (Table 3), and the determination of the aluminum solid solution composition (Table 4). At the same time, it was taken into account that the formation of the ingot structures upon nonequilibrium solidification always occurs under production conditions, and the phase composition is often very different from the equilibrium composition.
A common feature of the structures of all the experimental alloys was the presence of small amounts of Cu-containing phase particles that formed as a result of non–equilibrium solidification. These particles are located along the boundaries of the dendritic cells of the primary (Al) crystals. The Mg-free alloys consisted of the Al2Cu phase (Figure 2a,c,e,g), whereas in the 1% Mg alloys, Al2CuMg phase particles were also detected (Figure 2b,d,f,h). In Series B alloys, the number of Fe-bearing phases is expected to be small (Figure 2a–d). The Al6(Mn,Fe) phase particles were predominantly detected, but the Si-bearing particles were identified as Al15(Mn, Fe)3Si2. The particles of both Fe- bearing phases were distributed along the boundaries of the dendrite cells, and their morphologies differed slightly from those of the Al2Cu and Al2CuMg phase particles (Figure 2a–d).
In the Series BF alloys, almost all the analyzed particles were of the Al15(Mn,Fe)3Si2 phase, with a skeletal morphology typical of this phase (Figure 2e–h). A small amount of the Mg2Si phase was also detected in the BFM and BFMZ alloys (Figure 2f,h). In general, as expected, the microstructures of the iron and silicon–containing alloys differ from those of alloys with low contents of these elements by significantly larger amounts of Fe-bearing particles (compare the left and right SEM images in Figure 2). The uniform distribution of these particles and the relatively favorable morphology (absence of needle-shaped particles) suggest a fairly high deformation manufacturability of the Series BF alloy ingots.
According to the EMPA data, the Series B alloys contained almost all of the manganese dissolved in (Al). In the Series BF alloys, the concentration of Mn in (Al) is somewhat lower because, according to Table 1, its concentration in these alloys is approximately 0.1% less than that in Series B alloys. Moreover, this can be attributed to the fact that part of this element is bound to the Al15(Mn,Fe)3Si2 phase. Zinc (in alloys where it is present) is completely dissolved in (Al), while magnesium and copper are dissolved only partially because they form Al2Cu and Al2CuMg eutectic phase particles.

3.2. Microstructure of Hot-Rolled and Cold Sheets

Defect-free sheet alloys were obtained via hot rolling. Analysis of the microstructure of the hot–rolled sheets showed that Fe-bearing phase particles, formed during casting (Figure 2), remained after rolling (Figure 3) due to the low iron solubility in (Al). The numbers of Al2Cu and Al2CuMg particles decreased slightly because magnesium and copper were partially dissolved in (Al) during heating at 400 °C (before and during rolling). It should be noted that there was a significant improvement in the morphology of crystallization–origin particles due to their fragmentation during deformation. This effect was most pronounced in the structures of the Series BF alloys (Figure 3e–h).
In addition to the changes in the microstructure described above, heating at 400 °C during the hot rolling, according to our previous work, should lead to the partial decomposition of (Al) and the subsequent formation of Al20Cu2Mn3 dispersoids [18,19,55]. Since their sizes did not exceed 100 nm, they were not detected (Figure 3).
The compact (nearly globular) morphology of the eutectic particles, formed as a result of hot rolling (Figure 3), provided for high-quality cold–rolled sheets with a thickness of 0.5 mm (75% reduction rate). Because the structure of the hot–rolled sheets still contained Cu-containing particles of eutectic origin, they were detected in the structure even after cold rolling, along with the Fe-containing phases (Figure 4). In this case, the morphology of the crystallization–origin particles improved due to their fragmentation during deformation. After annealing at 400 °C, when a close-to-equilibrium state was achieved, the number of Cu- and Mg-containing phase particles decreased, as shown by comparing the microstructures shown in Figure 2a–d and Figure 4a–d. Since the Fe solubility in (Al) is very low, the amount of Fe-bearing particles changes only slightly. However, the compact morphology of these particles and their uniform distribution allows one to rate the structure of the Series BF alloys as favorable, as can be seen from Figure 4e–h.
Cold rolling leads to an increase in the dislocation density and formation of a cellular structure, as illustrated in Figure 5a for the BFMZ alloy by way of an example. Al20Cu2Mn3 dispersions with sizes less than 100 nm were also detected (Figure 5b). Annealing at 400 °C preserved a completely non-recrystallized structure in all the alloys. As shown in Figure 5c, the size of the subgrains did not exceed 1 μm, and the dislocation density was still quite high locally. The size of most dispersoids was less than 100 nm, and only a few of them grew to 150–200 nm (Figure 5d). It is obvious that the Al20Cu2Mn3 dispersoids prevent the recrystallization and growth of subgrains, as shown in our previous work [18,19,55].
Table 5 shows the experimental data for the composition of the aluminum matrix in the as-annealed cold–rolled sheets. The data show that the copper concentration increased significantly compared to that in the as-cast state (compare the values in Table 4 and Table 5). However, one must keep in mind the significant differences in the nature of these values. In the as-cast state, they correspond to the composition of (Al); in the 0.5CR400 state, they refer to a mixture of (Al) and Al20Cu2Mn3 dispersoids (Figure 5d).

3.3. Mechanical Properties of Cold-Rolled Sheets

Table 6 presents the tensile test data for the cold–rolled sheets in the initial state (0.5CR) and after annealing at 400 °C (0.5CR400). The results suggest that the addition of magnesium in the 0.5CR state significantly increases the tensile strength and yield strength, while the effect of zinc is negligible. The relative elongation of all the alloys in this state was low (within 0.3–2.3%), which can be associated with a high dislocation density (Figure 5a). However, more indicative is the 0.5CR400 state, when a relatively stable structure is formed. Annealing increases the ductility of all the alloys to approximately the same level (5.8–9.6%) and reduces the strength properties. However, the degree of this reduction varies quite widely.

3.4. Fracture Mechanism of Cold-Rolled Sheets

The Fracture surface analysis of the experimental alloys showed a dimple structure b in the initial (Figure 6 and Figure 7) and as-annealed states (Figure 8 and Figure 9). However, the depth of the dimples in the as-annealed state was significantly greater, and their size was smaller, which was associated with greater deformation before failure. It should be noted that the nature of the fracture of the Series B and BF alloys is practically the same, which can be associated with the uniform distribution of the Al15(Mn,Fe)3Si2 phase inclusions and their compact morphology (Figure 4e,f).

4. Discussion

The data on the mechanical properties (Table 6) show that the combined addition of 0.5% Fe and 0.4% Si to Series B alloys had virtually no effect on their ductility (Figure 10a) but reduced the ultimate tensile strength (UTS) (Figure 10b) and yield strength (YS) (Figure 10c) values for the cold–rolled sheets annealed at 400 °C. At the same time, the presence of magnesium and zinc had virtually no effect on this trend. To explain this effect, differences in the structure and phase composition between the experimental alloys were analyzed.
The sufficiently high El of the Series BF alloys, while maintaining a non-recrystallized structure, can be accounted for by the uniform distribution of the globular Al15(Mn,Fe)3Si2 phase particles, the average size of which does not exceed 3 μm (Figure 4e–h). It is well known that this phase is preferable to other Fe-bearing phases present in aluminum alloys, particularly the needle-like phases (Al3Fe, Al7Cu2Fe, and Al5FeSi) [33,35,56,57]. However, the formation of this phase, as determined from the phase diagram, depends on the alloy composition. In 2xxx series alloys, particularly 2219, with a high Cu:Mn ratio, the presence of iron typically leads to the formation of the Al7Cu2Fe phase, which tangibly reduces the mechanical properties [58,59,60]. Therefore, the concentration of iron in graded alloys is strictly limited.
The effects of the particle morphology on the nucleation and growth of the micropores are shown in Figure 11. In the case of needle-like particles, elongated pores were formed (as a result of cracking or fracture at the interfacial boundary) (Figure 11a). Since they are stress concentrators, they lead to the formation of sharp microcracks in the early stages (including the process of rolling). In the case of globular particles, the micropores assumed a compact shape and grew to merge without the formation of microcracks (Figure 11b). As a result, the plastic deformation covers a much larger volume, which leads to the formation of deep dimples (Figure 9).
As follows from previous works on Al–Cu–Mn system alloys in the concentration range considered [18,19,20,21,22,23], annealing at 400 °C and higher temperatures makes it possible to obtain a close-to-equilibrium state. Therefore, appropriate calculations of the phase composition were carried out for the test alloys. It can be seen from Table 7 that the fraction of the Al15(Mn,Fe)3Si2 phase in the Series BF alloys is 4–5 times that in the Series B alloys. This was generally consistent with the microstructure observed in the 0.5CR400 state (Figure 4). However, it should be noted that equilibrium phase composition calculations do not allow the estimation of the distribution of manganese between particles of crystallization origin and dispersoids.
Based on the experimental data on the chemical composition of the alunimum solid solution (Table 5), it was calculated (by Thermo-Calc) the expected fraction of the dispersoids in the aluminum matrix (thus, aluminum matrix consists of aluminum solid solution and secondary precipitates (including Mn-containing dispersoids)). As can be seen from Table 8, the amount of the Al20Cu2Mn3 phase in all Series B alloys is approximately the same (~6.1–6.5 wt.%) and is almost insensitive to the presence of magnesium and zinc, which are mainly included in the (Al) composition. On the other hand, the calculated amount of Mn-containing dispersions was significantly lower in the Series BF alloys (~3.1–5.2 wt.%). At the same time, in the BF and BFZ alloys, in addition to Al20Cu2Mn3, the formation of diseproside Al15Mn3Si2 can also be expected, which is typical of some 3xxx series alloys [40,61,62]. It should be noted that for some alloys (BFM, BFMZ) with a high fraction of the Al15 phase (Table 7), the estimated (based on the experimental data) result of the Al20 dispersoids fraction (Table 8) is somewhat increased (even taking into account that the fraction of the (Al) is ~92 wt.%). The latter result is coupled with the fact that the calculated data in Table 7 are obtained for the equilibrium state; thus, all the phases have an equilibrium composition. However, the Al15(Mn,Fe)3Si2 compound is, in fact, a line phase, and its actual composition can differ significantly in comparison with the equilibrium. Due to the latter, the calculated fraction of the phases in Table 7 can differ from what actually exists. On the other hand, the data in Table 8 were obtained based on the experimental data and, therefore, deserve greater confidence.
This is likely due to the difference in the quantity (QM) of Mn-containing dispersoids, which is the main cause of the reduced strength of the Series BF alloys compared to the Series B alloys (Figure 10b, c). Earlier data suggest [55] that the size of the subgrains (grain boundary strengthening, which in the general case should depend on QM) should make the greatest contribution to the YS of the BMZ alloy. This is supported by a comparison of the degree of reduction in the YS, UTS, and QM (the difference between the respective values of Series B and BF alloys). As can be seen from Table 9, QM correlates well with YS and UTS (some differences may be related to other factors). This suggests that the decrease in strength can be compensated by increasing the concentration of Mn (or other dispersoid-forming metals) in (Al). It should also be noted that although the Mn-containing disperoid fraction is the main cause of the reduced strength of the Series BF alloys, there may be other less significant factors (e.g., disperoid size and distribution depending on the actual alloy composition), the total effect of which may have some influence on the results obtained, and the influence of which requires additional precision studies to be determined.
From the results obtained, it can be concluded that adding Fe, Si, Mg, and Zn to the base alloy in a total amount of more than 3% makes it possible to retain the ductile fracture patterns of the base alloy and obtain a fairly higher level of the mechanical properties. This suggests the principal possibility of using a variety of secondary raw materials (containing the main elements present in aluminum alloys of different alloying systems) to prepare a base alloy that does not require homogenization and hardening.

5. Summary

The effect of the Fe-bearing phases on the structure, mechanical properties, and fracture mechanism of an Al–2%Cu–1.5%Mn(Mg,Zn) non–heat–treatable model sheet alloy designed for Al20Cu2Mn3 dispersoids was studied. The conclusions are shown in the following points:
  • The addition of 0.5% iron and 0.4% silicon to the base alloy, regardless of the presence of Mg and Zn, led to the formation of Al15(Mn,Fe)3Si2 eutectic phase inclusions.
  • Despite the relatively high proportion of eutectic particles, non-homogenized ingots of alloys containing Fe-bearing phases demonstrated sufficiently high deformation processability during hot (400 °C) and cold rolling, making it possible to obtain suitable sheets (with reduction degrees of 80 and 75%, respectively).
  • During the rolling process, the Al15(Mn,Fe)3Si2 phase particles are fragmented into compact inclusions with sizes less than 2 μm, and their distribution becomes uniform.
  • The combined addition of 0.5% Fe and 0.4% Si to the “pure” alloys Series B has virtually no effect on the ductility of cold–rolled sheets but reduces the strength properties.
  • The uniform distribution of the Fe-bearing phases in the structure of the cold–rolled sheets contributes to the preservation of the dimple fracture toughness mechanism. This helps to maintain the same level of ductility in cold–rolled sheets as in “pure” alloys.
  • This work has demonstrated the fundamental possibility of using a variety of secondary raw materials (containing the main elements present in aluminum alloys of different alloying systems) for the preparation of a base alloy that does not require homogenization and hardening.

Author Contributions

Conceptualization, N.B.; investigation, K.T.; methodology, T.A., S.C. and N.A.; supervision, N.B.; writing—original draft, N.B. Writing—review, and editing: T.A. All authors have read and agreed to the published version of the manuscript.

Funding

The authors gratefully acknowledge the financial support provided by the Russian Science Foundation project no. 23-79-30015).

Data Availability Statement

The raw/processed data required to reproduce these findings cannot be shared at this time due to technical or time limitations.

Conflicts of Interest

The authors declare no conflict of interest.

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Figure 1. Isothermal sections of Al–Cu–Mn–Fe–Si (a) and Al–Cu–Mn–Mg–Zn–Fe–Si (b) phase diagrams at 2% Cu, 1.5% Mn and 400 °C: (b) at 1% Mg and 1% Zn, right drawings—fragments of left drawings at low concentrations of Fe and Si, all phase fields contain (Al), Al20-Al20Cu2Mn3, Al6-Al6(Mn,Fe), Al15-Al15(Mn,Fe)3Si2, S-Al2CuMg, Al7-Al7Cu2Fe.
Figure 1. Isothermal sections of Al–Cu–Mn–Fe–Si (a) and Al–Cu–Mn–Mg–Zn–Fe–Si (b) phase diagrams at 2% Cu, 1.5% Mn and 400 °C: (b) at 1% Mg and 1% Zn, right drawings—fragments of left drawings at low concentrations of Fe and Si, all phase fields contain (Al), Al20-Al20Cu2Mn3, Al6-Al6(Mn,Fe), Al15-Al15(Mn,Fe)3Si2, S-Al2CuMg, Al7-Al7Cu2Fe.
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Figure 2. Microstructure of experimental alloys in as-cast ingots (10F), SEM: (a) B, (b) BM, (c) BZ, (d) BMZ, (e) BF, (f) BFM, (g) BFZ, (h) BFMZ.
Figure 2. Microstructure of experimental alloys in as-cast ingots (10F), SEM: (a) B, (b) BM, (c) BZ, (d) BMZ, (e) BF, (f) BFM, (g) BFZ, (h) BFMZ.
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Figure 3. Microstructure of experimental alloys in hot rolled–sheets (2HR), SEM: (a) B, (b) BM, (c) BZ, (d) BMZ, (e) BF, (f) BFM, (g) BFZ, (h) BFMZ.
Figure 3. Microstructure of experimental alloys in hot rolled–sheets (2HR), SEM: (a) B, (b) BM, (c) BZ, (d) BMZ, (e) BF, (f) BFM, (g) BFZ, (h) BFMZ.
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Figure 4. Microstructure of experimental alloys in cold rolled–sheets after annealing at 400 °C (0.5CR400): (a) B, (b) BM, (c) BZ, (d) BMZ, (e) BF, (f) BFM, (g) BFZ, (h) BFMZ.
Figure 4. Microstructure of experimental alloys in cold rolled–sheets after annealing at 400 °C (0.5CR400): (a) B, (b) BM, (c) BZ, (d) BMZ, (e) BF, (f) BFM, (g) BFZ, (h) BFMZ.
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Figure 5. TEM structures of alloy BFMZ in cold–rolled–sheets: (a,b) 0.5CR, (c,d) 0.5CR400.
Figure 5. TEM structures of alloy BFMZ in cold–rolled–sheets: (a,b) 0.5CR, (c,d) 0.5CR400.
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Figure 6. Fracture surfaces of cold–rolled sheets of alloys B (a), BM (b), BZ (c), and BMZ (d) in the 0.5CR state after tensile testing via SEM (left: BSE image; right: SE image).
Figure 6. Fracture surfaces of cold–rolled sheets of alloys B (a), BM (b), BZ (c), and BMZ (d) in the 0.5CR state after tensile testing via SEM (left: BSE image; right: SE image).
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Figure 7. Fracture surfaces of cold–rolled sheets of alloys BF (a), BFM (b), BFZ (c), and BFMZ (d) in the 0.5CR state after tensile test, SEM (left: BSE image; right: SE image).
Figure 7. Fracture surfaces of cold–rolled sheets of alloys BF (a), BFM (b), BFZ (c), and BFMZ (d) in the 0.5CR state after tensile test, SEM (left: BSE image; right: SE image).
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Figure 8. Fracture surfaces of cold–rolled sheets of alloys B (a), BN (b), BZ (c), and BMZ (d) in the 0.5CR400 state after tensile testing by SEM (left: BSE image; right: SE image).
Figure 8. Fracture surfaces of cold–rolled sheets of alloys B (a), BN (b), BZ (c), and BMZ (d) in the 0.5CR400 state after tensile testing by SEM (left: BSE image; right: SE image).
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Figure 9. Fracture surfaces of cold–rolled sheets of alloys BF (a), BFN (b), BFZ (c), and BFMZ (d) in 0.5CR400 state after tensile test, SEM (left: BSE image, right: SE image.
Figure 9. Fracture surfaces of cold–rolled sheets of alloys BF (a), BFN (b), BFZ (c), and BFMZ (d) in 0.5CR400 state after tensile test, SEM (left: BSE image, right: SE image.
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Figure 10. Comparison of the mechanical properties of alloys with low Fe and Si (B, BM, BZ, BMZ) and alloys with high Fe and Si (BF, BFM, BFZ, BFMZ): (a) El, (b) UTS, and (c) YS. (1) B– BF, (2) BM– BFM, (3) BZ– BFZ, (4) BMZ– BFMZ.
Figure 10. Comparison of the mechanical properties of alloys with low Fe and Si (B, BM, BZ, BMZ) and alloys with high Fe and Si (BF, BFM, BFZ, BFMZ): (a) El, (b) UTS, and (c) YS. (1) B– BF, (2) BM– BFM, (3) BZ– BFZ, (4) BMZ– BFMZ.
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Figure 11. Influence of morphology of Fe-bearing particles on nucleation and growth of micropores: (a) needle-like particles and (b) globular particles.
Figure 11. Influence of morphology of Fe-bearing particles on nucleation and growth of micropores: (a) needle-like particles and (b) globular particles.
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Table 1. Chemical compositions of experimental alloys.
Table 1. Chemical compositions of experimental alloys.
Alloy DesignationConcentration, wt.%
CuMnMgZnFeSiAl
Series B
B2.061.660.030.030.110.08balance
BM2.11.651.090.0020.120.08balance
BZ21.720.021.100.120.07balance
BMZ2.071.61.021.110.130.08balance
Series BF
BF2.241.670.020.030.490.43balance
BFM2.211.541.250.020.480.37balance
BFZ2.11.580.030.980.50.4balance
BFMZ2.251.471.221.060.520.39balance
Table 2. Processing routes for the experimental alloys.
Table 2. Processing routes for the experimental alloys.
ProcessObtained ProductDesignation
CastingFlat ingot with sizes 10 × 40 × 180 mm10F
Hot rolling (at 400 °C) of foundry ingotHot rolled sheet 2 mm in thickness2HR
Cold rolling of hot rolled sheet Cold rolled sheet 0.5 mm in thickness0.5CR
Annealing of cold rolled sheet at 400 °C (3 h)Annealed cold rolled sheet0.5CR400
Table 3. Detected excess phases in the experimental alloys.
Table 3. Detected excess phases in the experimental alloys.
Alloy 1Excess Phases 2 of Eutectic Origin
As-cast ingot (10F)Annealed cold rolled sheet (0.5CR400)
BAl2Cu, Al6, Al15Al2Cu, Al6, Al15
BMS, Al6, Al15Al2Cu, Al6, Al15
BZAl2Cu, Al6, Al15Al2Cu, Al6, Al15
BMZS, Al6, Al15Al2Cu, Al6, Al15
BFAl2Cu, Al15Al2Cu, Al15
BFMAl2Cu, S, Al15, Mg2SiS, Al15, Mg2Si
BFZAl2Cu Al15Al2Cu, Al15
BFMZAl2Cu, S, Al15, Mg2SiS, Al15, Mg2Si
1 see in Table 1; 2 Al6—Al6(Mn,Fe), Al15—Al15(Mn,Fe)3Si2, S—Al2CuMg.
Table 4. Chemical composition of aluminum solid solution in as-cast ingots (10F).
Table 4. Chemical composition of aluminum solid solution in as-cast ingots (10F).
Alloy 1Concentration in (Al), wt.%
CuMnMgZnSiFe
B0.881.29<0.01<0.010.04<0.01
BM0.851.270.77<0.01<0.01<0.01
BZ1.121.29<0.010.85<0.01<0.01
BMZ0.541.220.440.71<0.01<0.01
BF0.681.16<0.01<0.010.08<0.01
BFM0.681.060.63<0.010.02<0.01
BFZ0.781.09<0.010.850.06<0.01
BFMZ0.811.180.710.840.06<0.01
1 see in Table 1.
Table 5. The chemical composition of aluminum matrix in annealed cold–rolled sheets (0.5CR400) was experimentally defined (by EMPA).
Table 5. The chemical composition of aluminum matrix in annealed cold–rolled sheets (0.5CR400) was experimentally defined (by EMPA).
Alloy 1Concentration in (Al), wt.%
CuMnMgZnSiFe
B2.001.30<0.01<0.010.02<0.01
BM1.681.280.90<0.010.02<0.01
BZ1.851.27<0.011.030.03<0.01
BMZ1.621.370.891.06<0.01<0.01
BF1.521.1<0.01<0.010.14<0.01
BFM1.441.110.89<0.010.02<0.01
BFZ1.521.09<0.011.110.15<0.01
BFMZ1.421.020.81.040.05<0.01
1 see in Table 1.
Table 6. Mechanical properties of cold–rolled sheets (0.5 mm).
Table 6. Mechanical properties of cold–rolled sheets (0.5 mm).
Alloy 1State 2UTS, MPaYS, MPaEl, %
B0.5CR3453380.6
0.5CR4003042267.1
BM0.5CR4114080.3
0.5CR4003522776.4
BZ0.5CR3423281.3
0.5CR4003032267.7
BMZ0.5CR4224171.6
0.5CR4003582806.0
BF0.5CR3333092.3
0.5CR4002761895.8
BFM0.5CR4033971.3
0.5CR4003112097.8
BFZ0.5CR3583352.0
0.5CR4002731809.6
BFMZ0.5CR4424380.5
0.5CR4003222027.3
Average deviations±15±10±1.2
1 see in Table 1 2; see in Table 2.
Table 7. Calculated fraction of phases in experimental alloys at 400 °C.
Table 7. Calculated fraction of phases in experimental alloys at 400 °C.
Alloy 1Fractions of Precipitates. wt.%
Al20Al15Al2CuMg2SiS(Al)
B7.230.96000balance
BM7.230.9600.020.02balance
BZ7.580.93000balance
BMZ6.91.0500.0020.08balance
BF3.065.180.7100balance
BFM4.053.9300.221.24balance
BFZ3.014.920.3900balance
BFMZ3.364.2900.21.38balance
1 see in Table 1.
Table 8. Calculated composition of the aluminum solid solution and fraction of dispersoids for the experimental alloys at 400 °C.
Table 8. Calculated composition of the aluminum solid solution and fraction of dispersoids for the experimental alloys at 400 °C.
Alloy 1Concentration in (Al), wt.%Fraction of Dispersoids 2, wt.%
CuMnMgZnSiFeAl20Al15
B1.110.06<0.01<0.010.02<0.016.29-
BM0.790.060.94<0.01<0.01<0.016.15-
BZ0.980.06<0.011.10.03<0.016.07-
BMZ0.660.070.951.13<0.01<0.016.56-
BF1.040.07<0.01<0.010.02<0.013.391.29
BFM0.670.070.92<0.01<0.01<0.015.26-
BFZ1.070.06<0.011.170.03<0.013.161.42
BFMZ0.720.070.771.090.01<0.014.8
1 see in Table 1; 2 calculated from the composition of the aluminum matrix (see in Table 5).
Table 9. Comparison of the decrease in the fraction (QM, mass.%) of dispersoids 1 and the decrease in strength properties 2 between alloys of series B and BF.
Table 9. Comparison of the decrease in the fraction (QM, mass.%) of dispersoids 1 and the decrease in strength properties 2 between alloys of series B and BF.
Pair of Alloys 3Decrease of Values
QM, mass.% YS, MPa 2UTS, MPa 2
B—BF 25.616.49.2
BM—BFM 14.524.511.6
BZ—BFZ 24.520.49.9
BMZ—BFMZ 26.727.910.1
1 see in Table 8; 2 see in Table 6; 3 see in Table 1.
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Belov, N.; Akopyan, T.; Tsydenov, K.; Cherkasov, S.; Avxentieva, N. Effect of Fe-Bearing Phases on the Mechanical Properties and Fracture Mechanism of Al–2wt.%Cu–1.5wt.%Mn (Mg,Zn) Non-Heat Treatable Sheet Alloy. Metals 2023, 13, 1911. https://doi.org/10.3390/met13111911

AMA Style

Belov N, Akopyan T, Tsydenov K, Cherkasov S, Avxentieva N. Effect of Fe-Bearing Phases on the Mechanical Properties and Fracture Mechanism of Al–2wt.%Cu–1.5wt.%Mn (Mg,Zn) Non-Heat Treatable Sheet Alloy. Metals. 2023; 13(11):1911. https://doi.org/10.3390/met13111911

Chicago/Turabian Style

Belov, Nikolay, Torgom Akopyan, Kirill Tsydenov, Stanislav Cherkasov, and Natalia Avxentieva. 2023. "Effect of Fe-Bearing Phases on the Mechanical Properties and Fracture Mechanism of Al–2wt.%Cu–1.5wt.%Mn (Mg,Zn) Non-Heat Treatable Sheet Alloy" Metals 13, no. 11: 1911. https://doi.org/10.3390/met13111911

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