<figure><img src="image_53.png" /><figcaption>**Fig. 51.** The effect of adding trace elements on the microstructure of recycled casting alloys. The micrographs in the middle are selected from [866,873,875] and reproduced with permission [865–869].</figcaption></figure>

but at higher strains [884]. The authors also found that coarser particles (>3 ~ 4 µm) tend to break more than once. This becomes more detrimental because the local density of the cracks will increase and their interaction will generate faster propagation of damage.

The damage mechanism, from the perspective of particle failure, is either fracture of the particle or its decohesion from the matrix [879,883,884]. This depends on many factors including morphology, size, alignment around the applied load, mismatch stress between the inclusion and the matrix. Babout et al. [885] have shown that the damage mechanism is also related to matrix hardness. For a soft matrix, particle-matrix debonding is the main damage mechanism, while if matrix strength is greater the same particles tend to break rather than debond. The number of damage events caused by both particle decohesion and cracks increase with plastic strain. However, in the case of a hard matrix the particles break at lower strains, and the number of cracked particles increases sharply with plastic strain. The morphology of the particles plays an important role in the localization of the strain and consequently in the damage mechanism. Two of the most common IMPs in Al 6xxx alloys are plate-like β-Al~5~FeSi and round α-Al~1~2(Fe,Mn)~3~Si particles [881,886,887]. The first is more detrimental due to its elongated morphology. The effect of loading direction associated with the morphology shows that β-Al~5~FeSi particles aligned parallel to the loading direction break into fragments, while those particles (either α-Al~1~2(Fe,Mn)~3~Si or β-Al~5~FeSi) perpendicular to the loading direction generate void nucleation through particle debonding from the matrix [881]. In general, when the angle between the long axis of the particle and the loading direction is >45°, particles tend to undergo interface decohesion. Otherwise, they fail by cracking. The cracks are normal to the main loading direction and the larger plate-like particles (β-Al~5~FeSi) tend to break into several fragments [881,883,884].

In the literature, ductile fracture is reported to be the main mechanism in the failure of 6xxx alloys [877,883]. This is characterized by the nucleation of voids, followed by their growth and coalescence. Petit et al. [877] showed that the coarse particles do not localize at the grain boundaries, but rather are distributed inside the grains. However, the particles also tend to form along the subgrain boundaries. As the damage initiation is dominated by failure of coarse Mg~2~Si or Fe-rich IMPs and these are distributed mostly within the grains, the main mechanism for the fracture is shown to be transgranular (see Fig. 52). In contrast, Shen et al. [879] have shown that the coarse particles are distributed along the grain boundaries, and since the damage nucleates and propagates through these precipitates along grain boundaries the fracture is assumed to be dominated by intergranular crack propagation. These different observations may be attributed to the use of different alloys and heat treatments. It is shown in [877,879] that large Mg~2~Si particles break at lower strains compared to other IMPs, even though their mean size is lower than that of IMPs. This can be seen in Fig. 53, where the damage originates from Mg~2~Si particles. Upon higher plastic strains, the cracks propagate through large IMPs. As a result, the fracture surfaces of 6061 samples are found to include many dimples caused by the breakage of Mg~2~Si particles. Only a few other IMPs are shown to be fractured, generating what are thought as of secondary cracks.

The Fe content in these alloys is a decisive parameter for the damage mechanism because, on the one hand, it contributes to the formation of Fe-containing IMPs, and on the other, it determines which Fe-containing phase will form in the microstructure. The increase in Fe significantly reduces ductility, while slightly enhancing yield strength [100]. One of the earliest studies by Blind et al. [888] showed that toughness increases due to the higher volume fraction of dispersoids in low-Fe alloys (<0.06 wt%), but decreases as the Fe content becomes higher (~0.30 wt%). This is attributed to formation of coarse IMPs [100,888]. Sweet et al. [886] showed that at lower Fe content (~0.1 wt%) β-Al~5~FeSi particles are dominant in the microstructure, while α-Al~1~2(Fe,Mn)~3~Si particles become the major type of intermetallic at increased Fe content (~0.5 wt%) (see Fig. 54). Substitution of plate-like particles with a more compact α-Al~1~2(Fe,Mn)~3~Si phase in the form of round or Chinese-script-shaped particles improves ductility. Adding Mn has been found to
suppress the development of β-Al~5~FeSi particles and to promote the formation of the more spherical-shaped α-Al~12~(Fe,Mn)~3~Si precipitates. This transformation can be triggered by homogenization at relatively high temperatures [887,889,890]. A reduced Si level has also been reported to enhance the spheroidization of intermetallic particles containing Al-Fe-Si [889]. It has also been observed that in order to enforce transformation of the β-Al~5~FeSi phase into the α-Al~8~Fe~2~Si phase, a Mn:Fe ratio of 0.5 is particularly desirable, producing an α-Al~8~(Fe~2~Mn)Si phase or an α-Al~12~(Fe,Mn)~3~Si phase in the material [100]. When considering IMPs as damage initiation spots, the concentration of Fe and Mn becomes important. The magnesium content of 6xxx alloys has no significant effect on the formation of Fe-rich IMPs, but rather on the formation of Mg~2~Si particles. There is also evidence that Mg addition can promote the formation of π-Al~8~FeMg~3~Si~6~ particles with a hexagonal structure that forms at lower temperatures than β-Al~5~FeSi [886]. These particles also have Chinese-script morphology. Ji et al. [100] found that the size and morphology of the α-Al~12~(Fe,Mn)~3~Si particles can be altered by the cooling rate. At higher cooling rates (over 10~3~ K/s), Chinese-script-like particles form and assume a fine compact morphology, and their particle size is almost insensitive to variations in Fe content, with average particle sizes around 0.75 µm. However, their volume fraction increases by approximately 30% if the Fe content is altered from 0.5 to 1.5 wt% [100]. In contrast, at lower cooling rates (below 10~2~ K/s) particle size and the volume fraction of α-Al~12~(Fe,Mn)~3~Si particles both increase notably with Fe content, with a maximum average size of ~60 µm observed for 1.8 wt% Fe.

Formation of precipitate-free zones (PFZ) adjacent to grain boundaries can also decrease toughness and facilitate damage propagation in this class of aluminum alloys [891–893]. Moregeneyer et al. [891] studied the quench sensitivity of toughness in 6156 alloys, and found that the energy required for crack propagation drops substantially in air-cooled samples with relatively large PFZs compared to those which are water-cooled. In air-cooled samples the cracks are also observed to be thoroughly intergranular, and most cracking occurs parallel to the loading direction. By conducting compression tests with different strains, Christiansen et al. [894] showed that PFZs and grain interiors in 6060 alloys can exhibit different orientations. The misorientation between the grain interior and PFZs becomes more pronounced at larger strains. The grain-like PFZs are expected to inhibit void nucleation at grain boundaries, but their boundary with interior grain will be susceptible to damage initiation [891].

<figure><img src="image_54.png" /><figcaption>**Fig. 52.** Fracture surfaces of two 6061-T6 alloys subjected to different aging times of 4 and 16 h. The fracture mechanism in both cases is transgranular. Dimples correspond to broken Mg~2~Si precipitates and circled areas correspond to very large IMPs. Reproduced with permission from [877].</figcaption></figure>
5.4. Influence of scrap-related contaminant elements on the quench sensitivity of aluminum alloys

Thermal treatment of precipitate-strengthened alloys usually involves three steps: solution treatment, quenching and aging [895]. The purpose of solution heat treatment is to fully dissolve the soluble alloying elements into the aluminum matrix. After rapid quenching, a supersaturated solid solution (SSSS) begins to form and finally decomposes during artificial aging. The quenching rate strongly affects the supersaturation degree and thus subsequent age hardening. Quench sensitivity refers to the reduction in final properties (e.g. strength, ductility/toughness) due to the reduction of the quenching cooling rate [538,896]. The effect of quench sensitivity on the final properties of aluminum alloys and the characterization methods of quench sensitivity are summarized in Fig. 55.

The quenching rate changes the width of the PFZ by affecting the concentration of as-quenched-in vacancies. A low quenching rate allows more vacancies to become annihilated at sinks like the nearest grain boundaries (or dislocation jogs), favoring a greater PFZ width. Thereby the quenching rate substantially affects many critical damage properties, including stress corrosion cracking (SCC) susceptibility, fatigue and fracture. A faster quenching rate often corresponds to higher peak-aged hardness and greater SCC resistance. Quench sensitivity can also be measured by two indicators: (i) the upper critical cooling rate (UCCR) (higher than which a fully supersaturated solid solution can be reached); and (ii) the corresponding relationship between the strength and the cooling rate [538]. More attention has been paid to the quench sensitivity of thick plates [897–900]. The quench rate is relatively low in the center but high at the surface, resulting in different solute precipitation sizes [896]. The cooling rate near the core is lower than the UCCR. After solution treatment, the solute atoms precipitate during quenching to form coarse particles, which greatly reduces solute supersaturation and precipitation strengthening during subsequent aging.

The effect of quench sensitivity should also be considered in aluminum alloy sheets used for hot stamping. Because of the high forming temperature, the equipment must be specially designed to achieve rapid cooling of quench-sensitive alloys to avoid the formation of coarse secondary phases at grain boundaries and dislocations [901,902].

Quench sensitivity depends on alloy composition [896,905]. Therefore, it is important to study the effect of contaminant elements on the quench sensitivity of recycled aluminum alloys. There are three types of alloying element in the different aluminum alloy systems [904]: (i) the major alloying elements; (ii) the impurity elements; (iii) deliberately-added minor elements. Taking 7xxx alloys

<figure><img src="image_55.png" /><figcaption>**Fig. 53.** Nucleation and propagation of the damage in a 6061 aluminum alloy in T6 condition. The damage initiates from the decohesion of Mg₂Si particles (e.g. number 1, circled in blue) and then continues by propagation and coalescence with other voids until it travels through the captured region. The circled regions correspond to Mg₂Si precipitates (blue), Fe-containing IMPs (red), and pre-existing voids (purple). Reproduced with permission from [879]. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)</figcaption></figure>
<figure><img src="image_56.png" /><figcaption>**Fig. 54.** Morphology (deep etched) of two different intermetallic phases by variation of Fe content in a 6060 aluminum alloy: (a) the plate-like β-Al~5~FeSi phase with 0.1 wt% Fe; and (b) the Chinese-script α-Al~12~(Fe,Mn)~3~Si phase with 0.5 wt% Fe. Reprinted with permission from [886].</figcaption></figure>

as an example, Zn, Mg and Cu are the major alloying elements, Mn, Cr, Zr, and Sc are the minor elements, and Fe, Si, etc. are the impurity elements. The Al-Zn-Mg-(Cu) aluminum alloy is considered to be the most sensitive to quenching rate because it usually contains a high concentration of alloying elements [538]. For this reason the quench sensitivity of 7xxx aluminum alloys has been studied extensively [413,896,897,900,906]. Quench sensitivity increases with the total content of alloying elements (Zn + Mg + Cu) due to the enhanced tendency of the supersaturated solid solution to decompose during the cooling process [899,905]. The increase in Mg content in 7085 aluminum alloys may reduce the solubility of the other main alloying elements, leading to an increase in the driving force for η phase precipitation during quenching [897]. It is believed that the 7xxx series aluminum alloy with a higher Zn/Mg ratio has relatively low quench sensitivity [905]. The possible explanations for this are (i) the decelerated decomposition rate of the solid solution to a stable phase [907] and (ii) the significant suppression of the precipitation of the S phase (Al~2~CuMg) during continuous cooling [904]. In consequence more Cu and Mg can be retained in solid solution, resulting in reduced quench sensitivity. Adding Cu results in higher quench sensitivity [749]. However, Livak and Papazian [908] demonstrated that a small addition of Cu to an Al-Zn-Mg alloy could decrease the quench sensitivity by promoting the formation of solute-vacancy complexes. In 6xxx alloys, changes in the content and ratio of the main alloying elements also seriously affect quench sensitivity. Milkereit et al. [411] pointed out that increasing the total content while keeping the Mg/Si ratio unchanged increases quench sensitivity. A comparative study also reveals that higher Mg and Si content induces a faster critical cooling rate and thus greater quench sensitivity [909]. In 2xxx aluminum alloys increasing the Cu content will also increase quench sensitivity [910,911]. For example, compared with 2024, 2219 with a higher Cu content is more sensitive to the cooling rate and is more likely to generate an θ phase during cooling [538]. Therefore both the content and the proportion of alloying elements must be controlled to avoid unacceptable quench sensitivity.

Adding minor elements such as Mn, Cr and Zr to the aluminum alloy can increase the latter's recrystallization temperature and inhibit the occurrence of recrystallization [896,905]. The reduction in grain size will increase the alloy's quench sensitivity by increasing the number of both nucleation sites for particles during cooling and sites where excess vacancies can annihilate. In 7xxx series alloys, the minor elements added usually accumulate in the matrix in the form of second phases, e.g. Al~18~Cr~2~Mg~3~ (E phase) and an AlMgMn phase [905], which can act as nucleation sites for the equilibrium η phase. Zr has been reported to inhibit recrystallization through the presence of Al~3~Zr dispersoids and to reduce the alloy's quench sensitivity [497,896]. The crystallographic orientation relationship plus the interfacial energy between the dispersoid phase and the precipitate play a substantial role in quench sensitivity. In Al-Zn-Mg alloys Ge can significantly inhibit quench-induced coarse precipitates [896,912]. The Al-Zn-Mg-Cu alloy containing both Ag and Ge has a lower coarse equilibrium phase density after slow quenching. It is believed that Ag/Ge trap free vacancies and reduce the amount of coarse precipitation on particles containing either Cr or Zr during quenching, leading to the lower quench sensitivity [896]. Fe and Si are common impurities in 7xxx series aluminum alloys, and their content should be controlled to remain at very low levels [913,914]. The mechanism underlying the effect of Fe and Si on quench sensitivity is still a subject of debate. Si may combine with Mg to form the Mg~2~Si phase, while Fe can promote the formation of Al~15~(FeMn)~3~Si~2~ with Mn. These primary intermetallic phases are insoluble coarse particles which may cause premature cracks and failure [915]. The traditional belief is that adding Fe and Si will increase quench sensitivity due to the presence of dispersoids and such intermetallic phases [905]. However, the impurity elements Fe and Si mainly exist in Al-Zn-Mg-Cu alloys such as the A1~7~Cu~2~Fe phase and the Mg~2~Si phase [916]. The formation of these phases consumes large amounts of Cu and Mg, reducing the solutes supersaturation and the equilibrium phase precipitation during slow quenching.

Adding Mn and Cr will increase the quench sensitivity of 6xxx aluminum alloys [917–919]. In a way similar to 7xxx alloys, the dispersoids containing Mn or Cr formed during the cooling process promote β/β' precipitation. Taking the 6063 alloy as an example, a cooling temperature of 0.8 °C/s followed by artificial aging can obtain 90% strength in T6 state. When the alloy contains excessive Mn/ Cr, the cooling rate needs to be increased to 25 °C/s to reach the same level [917–919]. Saito et al. [920] found that a small amount of Cu (~0.1%) will not change the quench sensitivity in an Al-Mg-Si alloy. The average precipitate size reduces significantly, while the precipitate number density increases after subsequent artificial aging in the copper-containing alloy after either air cooling or water quenching [920]. However, Kassner et al. pointed out that when too much Cu is introduced, it increases quench sensitivity [921]. Tzeng et al. [922] reported that adding a small amount of Sc (0.04 wt%) can reduce the quench sensitivity in Al-Mg-Si alloys. They
<figure><img src="image_57.png" /><figcaption>**Fig. 55.** The effect of quench sensitivity on the final properties of aluminum alloys and characterization methods for the material’s quench sensitivity: (a) the effect of quenching rate on the peak aging microstructure; (b) the effect of quenching rate on the vacancy concentration adjacent to a grain boundary and the width of PFZ [903]; (c) continuous cooling precipitation diagram for 7150 [538]; (d) tensile strengths of eight aluminum alloys as a function of average cooling rate during cooling [904]. The diagrams in (c) and (d) are reproduced with permission.</figcaption></figure>

believe that Sc atoms have the ability to preferentially trap vacancies and form Al₃Sc during the cooling process, effectively hindering the movement of Mg atoms and thereby delaying the formation of Mg₂Si. In 2xxx aluminum alloys, adding trace elements such as Mg and Mn also affects quench sensitivity [538,910,923,924]. During cooling, S phase Al₂CuMg and T phase Al₂₀Cu₂Mg₃ tend to precipitate at the grain boundaries along with a PFZ, all of this having a negative impact on the strength of the alloy. Fig. 56 summarizes the effects of various elements on the quench sensitivity of three heat-treatable aluminum alloys.

## 5.5. Influence of scrap-related contaminant elements on the fatigue behavior of aluminum alloys

Scrap-related elements such as Si, Fe, Mn and Cu are commonly present in commercial aluminum alloys as precipitates or secondary phases, and they accumulate during recycling. The influence of Si on the fatigue behavior of hypo-eutectic Al-Si cast alloys is twofold. On the one hand, the eutectic Si particles can cause crack deflections, and slow down crack propagation [925,926]. On the other, large eutectic Si particles can increase the probability of crack nucleation and promote formation of crack networks [925,927]. Which mechanism dominates depends strongly on the stress or strain levels. Under low stress conditions with only minor plastic deformation (i.e. high cycle fatigue (HCF)), the number of cracks is very small and normally only a single main fatigue crack dominates. In such scenarios a higher Si content would be very harmful, because the eutectic Si particles would then be the main reason for crack nucleation [927]. At higher fatigue deformations (i.e. low cycle fatigue (LCF)), cracks can nucleate nearly everywhere in the eutectic structure, and the number of crack nucleation events is usually sufficiently high to enable the further growth of cracks. At the early stage of fatigue with relatively high local plastic strain, when many cracks form and propagate, they can be deflected by hard eutectic Si particles. Here only a few cracks grow further and cause fatigue damage.

In cast aluminum alloys Fe usually causes the formation of Fe-Al based intermetallic compounds, e.g. β-Al₅FeSi in needle- or plate-like shapes, which can have a detrimental effect on mechanical properties [339,928]. The effect of Fe content on the fatigue life of a cast Al-7 %Si-0.7 %Mg (A357) alloy with secondary dendrite arm spacing (SDAS) of 20–80 µm was studied by Wang et al. [929] with a sinus load frequency of 55 Hz. They found that an increase in Fe content from 0.08 to 0.14 wt% generates a decrease in the fatigue life of the cast alloy for larger SDAS (>60 µm), while its influence at small and intermediate SDAS values (<60 µm) is insignificant [929]. Horng et al. [930] studied Al-7 %Si-0.3 %Mg (A356) alloys with an Fe content varying from 0.14 to 0.97 wt%. They found that the primary Fe-rich IMCs are π-Al₈Mg₃FeSi₆ and α-Al₁₅Fe₃Si₂ at an Fe content of below 0.57 wt%, while β-Al₅FeSi is formed where there is higher Fe content, e.g. 0.97 wt% [930]. Compared to π-Al₈Mg₃FeSi₆ and α-Al₁₅Fe₃Si₂, the β-Al₅FeSi IMC is more detrimental to fracture resistance, and has an increasing effect on the mechanical properties of the materials with increasing Fe content [930]. It has also been observed that fatigue cracks propagate along Si particles and Fe-rich IMCs [930]. Gall et al. [931] investigated the fatigue crack paths in a cast Al-7 %Si-0.3 %Mg (A356) alloy containing modified IMCs, and found that during the later stage of fatigue the cracks propagate preferentially along Fe-rich (Al₁₅(Fe,Mn)₃Si₂) IMCs and Si particles.
<figure><img src="image_58.png" /><figcaption>**Fig. 56.** The influence of various alloying elements on the quench sensitivity of 2xxx, 6xxx and 7xxx aluminum alloys.</figcaption></figure>

The effect of Fe on fatigue behavior also depends strongly on stress and cycling time. Yi et al. [932] found that in the short lifetime regime (<10⁵ cycles) of the conventional cast Al-7 %Si-0.3 %Mg (A356) alloy the influence of Fe is insignificant, but that Fe-rich IMCs generate a remarkable decrease in fatigue life in the long lifetime regime (>10⁶ cycles) (see Fig. 57). With high levels of applied stress in the short lifetime regime, crack propagation dominates fatigue life. The growth of small cracks can be hampered by the large plate-like Fe-rich particles via crack branching and meandering in the high Fe content cast A356 alloy [932]. However, in the long lifetime regime with relatively low levels of applied stress, the large Fe-rich particles stimulate crack initiation by increasing the stress-strain concentration [932].

When Fe and Mn accumulate together during recycling, their joint effect on fatigue behavior is particularly interesting. Ceschini et al. [933] studied the effects of Fe and Mn content on the fatigue properties of an Al-10 %Si-2 %Cu casting alloy. The needle-like β-Al₅FeSi was the main intermetallic phase in the high-Fe content (0.5 wt%) alloy samples, while α-Al₁₅(Fe,Mn)₃Si₂ and π-Al₈Mg₃FeSi₆ intermetallic phases with a Chinese-script morphology were also present in the alloy samples containing Mn [933]. It was found that the α-Al₁₅(Fe,Mn)₃Si₂ and π-Al₈Mg₃FeSi₆ intermetallic compounds increased the fatigue resistance at high applied stress, while they decreased the cycles to failure at low applied stress [933]. These findings relate to the following facts. Dislocation movement is more effectively blocked by intermetallic compounds than by eutectic silicon particles under low applied stress, and thus the accumulation of dislocations at the interface between the alloy matrix and the intermetallic compounds leads to stress concentrations and consequently promotes crack nucleation with the matrix-particle decohesion or the cracking of particles. However, under high applied stress microcrack propagation determines the fatigue life [929,930,933]. Because Mn could modify the morphology of the needle-like β-Al₅FeSi and promote the formation of α-Al₁₅(Fe,Mn)₃Si₂ with a Chinese-script-like morphology (less harmful shape) [933], the overall fatigue and other mechanical properties could be optimized.

The influence of Sr (~0.2 wt%) on the fatigue behavior of an AlSi₆Cu₄ alloy was investigated by Huter et al. [927]. They observed a significantly longer fatigue lifetime for the Sr-modified alloy in both the high- and low-cycle fatigue regimes. Sr enables the eutectic silicon particles to assume smoother shapes and smaller sizes by promoting phase nucleation, thus acting as a refinement or modifying agent in the alloys [927,934]. It was found that the elastic strain energy necessary for interface cracking in the AlSi₆Cu₄ alloy without Sr modification is lower than that required in the Sr-modified alloy; thus crack nucleation in the former starts earlier during fatigue [927]. This suggests that Sr has a beneficial effect on preventing fatigue crack initiation. Moreover, because by adding Sr it was possible to induce more internal eutectic Si particles as barriers to fatigue crack propagation, average crack propagation speed could be reduced compared to the AlSi₆Cu₄ alloy without Sr modification.

The influence of Cu on the fatigue endurance limit of an AlSi₇Mg(Sr) alloy was also reported by Huter et al. [927]. In the initially solidified state, some primary Al₂Cu precipitates (the θ phase) can be present, and they can evolve into the coherent θ'' phase and the
partially incoherent θ' phase upon solution annealing treatment and artificial aging. Accordingly, Cu generates enhanced strength in the alloy due to precipitation. The Cu-induced precipitation-hardening effect can also increase the fatigue limit under HCF [927]. One possible explanation was that the fatigue load can be internally shed onto the precipitates, and the fatigue endurance can be enhanced via the Orowan mechanism [927]. However, when fatigue proceeds in the plastic regime, alloys with lower Cu content have higher ductility and hence show higher fatigue life [927].

Fig. 58 summarizes the above discussions about the effects of some scrap-related elements (Si, Fe, Mn, Sr and Cu) on the fatigue behavior of aluminum alloys. The scrap-related elements influence the fatigue properties mainly by precipitation and/or modification of intermetallic compounds or secondary phases. The related effects are strongly dependent on the stress and plasticity levels involved in fatigue. The shape of the precipitates is also important for the fatigue properties, e.g., the Chinese-script-like α-Al₁₅(Fe,Mn)₃Si₂ is less harmful than the needle-like β-Al₅FeSi, as discussed above.

Zhang et al. [935] recently suggested an alternative, or additional, method of making high-strength Al-alloys more resistant to fatigue. This involves a specific microstructure design that exploits the mechanical energy imparted during the initial cycles of fatigue to dynamically heal the weak points inherent in the microstructure. This concept substantially improved the fatigue life of ultrahigh-strength Al-alloys and may also work in impurity-contaminated alloy variants.

## 5.6. Influence of scrap-related contaminant elements on the corrosion of aluminum alloys

Aluminum and its alloys feature high chemical activity, but they usually exhibit excellent corrosion resistance owing to the formation of a thin, uniform and stable oxide film [936]. Modifying the alloying elements can alter the microstructure and affect the anti-corrosion performance of recycled aluminum alloys, as summarized in Fig. 59. The impurity elements accumulated in the recycled aluminum alloys usually promote the formation of intermetallic compounds, which can induce micro-galvanic cells due to the electrochemical potential difference between the compounds and the matrix. If the electrochemical potential of intermetallic compounds is higher than that of the aluminum matrix, these compounds (e.g., Al₂Cu) will become cathodic with respect to the anodic aluminum matrix, causing localized corrosion [937,938]. For instance, Fe contamination in aluminum alloys promotes electrochemical corrosion due to the formation of local electrochemical cells [939]. As the Fe content increases, a large number of anodically acting needles (i.e., Al₆Fe, Al₃Fe and Al₁₂Fe₃Si₂ intermetallics) form and produce more potential corrosion sites on the surfaces. Therefore, enhanced Fe content lowers the corrosion performance of recycled aluminum alloys [67,940]. Adding Mn or Cr can help to neutralize the adverse effect of Fe on corrosion resistance due to the formation of Al₆(Mn,Cr,Fe) and Al₁₂(Mn,Cr,Fe)₃Si IMCs, whose electrochemical potentials are similar to those of aluminum [941,942].

The size of the intermetallic compounds is also important, affecting the alloys' formability and their corrosion resistance in particular due the formation of micro-galvanic elements. More specific, large intermetallic compounds with micrometre dimensions usually cause corrosion performance to degrade substantially, while sometimes no corrosion attack at all can be found on or around nanometre-sized intermetallics [943,944]. For instance, if Zr is added to an 1070 Al-alloy as a model impurity this can trigger refinement and improved dispersion of Fe-containing intermetallics, an effect which mitigates the negative influence of the Fe-containing intermetallics in the alloy [732]. In general, the following aspects related to these beneficial elements contribute to neutralizing intermetallics' negative effects on corrosion [732,941,942]: (i) enhanced passivation performance, which slows the formation of corrosion pits; (ii) reduced potential difference between intermetallic compounds and the matrix; (iii) the refinement and improved dispersion of particles to prevent localized corrosion due to propagation. Note that non-metallic impurities (e.g., Al₄C₃ particles and entrapped oxides) generally cause corrosion resistance to deteriorate, as they may react with water at the surface

<figure><img src="image_59.png" /><figcaption>Fig. 57. Stress cycle number curves for a cast Al-7 %Si-0.3 %Mg (A356) alloy with low Fe content (<0.1 wt% Fe, denoted as "Low-Fe") and high Fe content (>0.55 wt% Fe, denoted as "High-Fe"). S<sub>f</sub> and S<sub>p</sub> specimens refer to specimens taken from the top and bottom positions of the wedge A356 ingot, respectively. Diagrams reproduced with permission from [932].</figcaption></figure>
[941,942].

Impurity elements also exert an influence on the corrosion behavior of aluminum alloys through their interaction with precipitates, either inside grains or at grain boundaries [945-947]. Examples of such interaction are seen in 6xxx alloys. In these materials the Zn content has to be strictly controlled to values below 0.03 wt% in recycled Al-Mg-Si ternary alloys, because larger amounts can cause corrosion attack on grains and interfaces in alkaline solutions which would be harmful to the anodized surface appearance.

Tolerance to Zn levels in 6xxx alloys made from scrap can be increased if a similar level of Cu (e.g., 0.1 wt%) is added. Preferential grain boundary corrosion attack rather than grain surface attack then occurs due to the segregation of Cu or the formation of precipitates (e.g. the Q phase) at grain boundaries. The accumulation of Zn was observed to increase the Zn/Cu ratio in the Q phase. This effect seems to change the potential difference of the Q phase relative to the neighbouring matrix [149].

In Al-Mg-Si-Cu-Zn alloys (i.e. modified 6xxx and most 7xxx alloys) with relatively high Zn/Cu content, intergranular corrosion (IGC) resistance first degrades with an increase of the Zn level to 1 wt%, and then improves with its increase to values above 2.0 wt% [946]. This effect is related to the compositional states of the Zn-containing grain boundary precipitates and the precipitation free zones (PFZs).

Ag was found to reduce the severe IGC susceptibility in the Si-excess Al-Mg-Si alloys. However, an accumulation of Ag above 0.3 wt % reverses this trend due to the more significant formation of intergranular anode β' phases [947].

In addition to scrap-related impurity elements (Fe, Zn, etc.) and microstructure-modifying elements (Mn, Zr, Cr, etc.), the main alloying elements Cu, Si, Mg can also affect the corrosion performance of recycled aluminum alloys [948] if they deviate from standard specifications. For example, adding Cu can cause the formation of Cu-containing phases or the segregation of Cu at grain boundaries, thereby reducing the corrosion resistance of these alloys [945,949,950]. Even trace levels of Cu (as low as 0.06 wt%) can induce substantial IGC sensitivity in Al-Mg-Si alloys, due to the formation of Cu-containing precipitates at high-angle grain-boundaries [950]. Enhanced scrap-related Si content in recycled alloys can be beneficial, as it generates (e.g.) higher strength, but it also increases corrosion tendency [951]. Because the concentration of Si atoms in PFZs is lower than that in grain interiors and in Si-containing grain boundary precipitates, the corrosion potential of the Si depletion layer (PFZ) is not only lower than that of the adjacent precipitates, but also lower than that of the adjacent grain interior. This can have severe anodic dissolution effects and lead to corrosion of the PFZs around the Si particles, thus increasing the material's overall IGC sensitivity [952]. This Si effect is important, because future scrap will probably have a higher overall Si content due to the massive number of cast vehicle parts from the automotive sector which will be returned in the next two decades.

In many cases adding Mg serves to enhance the corrosion resistance of aluminum alloys. However, too much Mg inherited from scrap can lead to the precipitation of anodically active Al₃Mg₂ IMCs at grain boundaries, creating the sensitization problem characteristic of 5xxx alloys. Therefore care must be taken to control the compositions of alloys made from 5xxx scrap or from mixed light metal scrap, to prevent them from triggering the unwanted grain boundary precipitation of Mg-rich IMCs.

Mg has a strong affinity with almost all of the other major alloying elements, e.g., Si, Cu and Zn, and can promote precipitation of unwanted phases that contain these elements in the grain interior [952,953]. The grain interiors then become less noble, leading to a smaller potential difference between the grain boundaries and the adjacent grain interiors. This effect reduces the material's IGC sensitivity [949].

The scrap-related tramp elements discussed above primarily act on specific microstructural and galvanic features. However, there are also impurity elements that are surface-sensitive and influence corrosion through a mechanism referred to as anodic activation [954,955]. Components pertaining to this group are, for example, the low-melting elements Pb, Sn, Bi and In. Most of the detailed

<figure><img src="image_60.png" /><figcaption>**Fig. 58.** Diagram summarizing the influences of some scrap-related elements on the fatigue behavior of aluminum alloys.</figcaption></figure>
<figure><img src="image_61.png" /><figcaption>Fig. 59. Effect of various scrap-related tramp elements on the corrosion of recycled alloys.</figcaption></figure>

studies have been conducted on the role of Pb [956,957].

When these elements accumulate above their solubility limits they segregate to the alloys' surfaces and can generate impurity-enriched metal-oxide films, for instance when the alloys are exposed to processing steps where these elements are sufficiently mobile. This changes the band-gap of the oxides, which leads to anodic activation, meaning that the oxides are destabilised, particularly under the effect of chloride ions. The phenomenon can appear in the form of enhanced pitting and even activate the entire oxide surface. Countermeasures include additional doping with noble elements such as Cu to change the surface potential [958].

Reviewing these different types of mechanism shows that a better understanding of the effects on corrosion of scrap-related composition deviation is crucial for the design of recycling-friendly aluminum alloys [959]. It is important to mention that advanced processing technologies, e.g. high pressure die-casting [960], solid-state recycling [940], severe plastic deformation [279] and casting methods featuring higher solidification rates can also lessen the difference between the behavior of secondary-synthesis alloys and primary-synthesis alloys in terms of corrosion resistance [936]. The modification of microstructure in recycled aluminum alloys with scrap-related contaminant elements through processing was discussed in more detail in section 3.6.

## 5.7. Trapping of contaminant elements at defects and in precipitates

Aluminum alloys generally show high sensitivity to embrittlement induced by impurities. For example, alkaline impurities such as Na and Ca which intrude into the material during the aluminum recovery process, can cause high-temperature embrittlement [961]. They segregate to grain boundaries and reduce their cohesive strength.

Another examples is hydrogen and the associated hydrogen embrittlement (HE) [962] which can have catastrophic impact on the mechanical properties of materials made from contaminated scrap. Organic matter in particular, which is inevitably introduced by post-consumer scrap into recycled aluminum alloys, acts as a hydrogen source. Therefore the hydrogen content in some recycled aluminum alloys may in certain cases be higher than that in aluminum alloys made entirely from primary materials. The potential of low-temperature atom probe tomography to observe the distribution of hydrogen (and deuterium) at various structural defects in metal alloys was demonstrated recently by Chen et al. [963], Breen et al. [964–966] and Zhao et al. [133]. Hydrogen trapped at defects (vacancies, dislocations, grain boundaries, sub-grain boundaries) and precipitates [963,967,968] can reduce stress corrosion resistance [133] and enhance the HE propensity of recycled aluminum alloys [131,967]. At room temperature, the solubility of hydrogen in aluminum is below $0.1 \times 10^{-6}$ wt%. Excessive hydrogen intrusion affects the mechanical properties of aluminum alloys in two ways: (i) segregated hydrogen induces interfacial decohesion at grain boundaries (HE) and (ii) micropores composed of hydrogen molecules cause ductile fracture [962,969]. Oger et al. [131] studied the segregation of hydrogen in 7xxx aluminum alloys in different artificial aging stages. They proposed that the interface between the precipitates and the Al matrix is an effective trap site for hydrogen, and that the presence of precipitates can effectively reduce hydrogen trapping by grain boundaries. Their findings thus emphasize that strictly controlling precipitates may aid resistance to HE [131]. Tsuru et al. [967] confirmed by first-principles calculation that the interface between MgZn₂ and the Al matrix is a more favorable site for trapping hydrogen compared to the grain boundaries in an Al-Mg-Zn alloy. Chao et. al. [968] found that the θ phase in 2xxx aluminum alloys is also capable of trapping hydrogen. Therefore, refinement of grains, introduction of dislocations/vacancies and the control of H-trapping precipitates are strategies that can be employed to effectively reduce HE in recycled aluminum alloys.

Trapping of scrap-related tramp elements also occurs in precipitates. It was observed that traces of Zn can introduce precipitates into an Al-Mg-Si alloy enhanced with Zn, resulting in a change in the precipitate's structure but without changing its sequence [831]. If the Zn content exceeds 1 wt% Zn will segregate to grain boundaries under peak aging conditions, leading to high IGC sensitivity. These observations suggest that trapping contaminant Zn in precipitates is an effective way to avoid the segregation of Zn at grain boundaries. Contaminants can also get trapped by their preferential enrichment around precipitates. For example, when Ag is added to an Al-Cu alloy some portion of the Ag will partition to the coherent interface at the θ' precipitates to form an Ag-rich double-layer structure [690], which is believed to reduce the interfacial energy between θ' and the matrix. Some Ag also segregates to the semi-
coherent θ'-matrix interface and acts as an obstacle to Cu diffusion. This interfacial trapping may hinder or prevent the lateral growth of θ' precipitates. Fig. 60 summarizes the distribution of various minor contaminant elements at defects/precipitates and their associated influence on the macroscopic properties of recycled aluminum alloys.

# 6. Effects of scrap-related impurities in cast alloys: Constitution, microstructure, processing and properties

## 6.1. Brief introduction to cast alloys

Given their high fluidity, low tendency to hot cracking and good as-cast surface quality, aluminum casting alloys offer good castability. Like wrought alloys, they are divided into alloy series according to their main alloying elements: Si, Cu, Mg, Zn and Sn [970,971] (see Table 15). Unfortunately, however, no globally recognized designation system exists. Many countries use specific alloy names. In the following reference is therefore made to the system of the Aluminum Association (AA), but where necessary a designation according to the chemical composition is also provided. The AA uses a four-digit numerical system. The first digit indicates the alloy group; the following two digits identify the different aluminum alloys in the group; and the digit that follows the decimal indicates the product form. Variations in specified compositions are denoted by the prefix letters A, B, C or D to define differences in impurity limits, in particular relating to Fe content. For example, the upper limit for Fe in A356 is 0.15 wt%, but in C356 it is only 0.07 wt%.

Castability is best with the 3xx and 4xx series and decreases on the order of 5xx, 2xx and 7xx.

Besides the major alloying elements, (i) minor elements, (ii) microstructure-modifying elements; and (iii) impurities must also be considered, as follows:

(i) *Minor elements*: Ni for enhancement of the elevated temperature service strength

(ii) *Microstructure-modifying elements*: Ti and B for refining primary aluminum grains; Sr, Na, Ca, Sb for modifying the morphology of the eutectic silicon phase in Al-Si alloys; Mn, Cr, V, Mo for changing the structure and morphology of the monoclinic β-Al<sub>5</sub>FeSi phase to the cubic α-Al<sub>15</sub>(MnFe)<sub>3</sub>Si<sub>2</sub> phase, which is much less harmful to ductility

(iii) *Impurities*: e.g. Fe, Cr and Zn

Various commercial alloys are subject to tight restrictions on a number of different elements that may negatively affect material properties. However, with regard to the influence of scrap-related impurities on property profiles, it is basically true that aluminum casting alloys are much more tolerant and less sensitive than wrought alloys [72]. This is mainly due to two things: aluminum casting alloys' generally higher alloying element content, and the significantly wider tolerance range of these alloying elements [972].

The alloys of the 3xx-series are the most popular in the casting industry because of their superior casting characteristics and good strength properties [973]. Secondary Al-Si-Cu-Mg alloys dominate the market for powertrain components and die-cast parts, where strength at elevated temperature and wear resistance are more important than ductility. Al-Si-Mg alloys – mostly primary alloys – are used for wheels, structural castings and suspension parts, which require medium to high strength and good ductility [972]. The 4xx-series alloys are used when good castability in combination with particularly good ductility (Mg-free) or particularly good corrosion resistance (Cu-free) is required [970].

Cast alloys of the 5xx-series are characterized by high resistance to corrosion, weathering and seawater. They are well suited for decorative anodic oxidized components. Conventional Al-Mg alloys have medium strength values, but exhibit good ductility properties. However, their castability is comparatively limited, i.e. they lack fluidity and are sensitive to hot cracking. Therefore they are not often used for die-cast components [971]. Contaminants and impurities should be avoided, as many applications involve surface finishing (polishing). Alloys suitable for die-casting typically contain about 1 wt% Si to form an Al-Mg<sub>2</sub>Si eutectic. Their mechanical properties are rather modest, with yield strength values around 90–120 MPa and elongation at fracture of roughly 5% [974]. More recent developments in the 5xx family concern alloy type AlMg5Si2Mn. They enable the die-cast production of large thin-walled structural parts with improved ductility. Depending on the wall thickness, good mechanical properties are achieved without heat treatment, i.e. in the “naturally hard” state (yield strength 140–170 MPa, elongation to fracture 9 – 14% for wall thickness 5 mm). However, high ductility requires the use of low-iron primary alloys with Fe < 0.2 wt% [975].

2xx-series cast alloys have been used in various industrial sectors such as the automotive and aerospace industries due to their high strength and good elevated temperature properties. The age-hardenable Al-Cu-Mg variants exhibit the highest strength levels of all aluminum casting alloys. With carefully controlled levels of impurities, excellent ductility is also achieved. The alloy A206 (AlCu4.5 Mg0.25), for example, reaches yield strength values of 350 MPa at 12% elongation to fracture, but the Fe content is limited to 0.15 wt% and the Si content to 0.05 wt% [976]. The corrosion resistance of 2xx alloys is not very high, and certain compositions may be susceptible to stress corrosion cracking in the high-strength state. Generally the alloys of the 2xx series are not preferred candidates for die-casting due to their poor fluidity and relatively high tendency to hot cracking, but they are increasingly being used in squeeze casting [971,972].

The alloys of the 7xx-series usually contain Mg and minor additions of one or more of the elements Cu, Fe, Cr and Mn. The alloys in the as-cast state exhibit a pronounced tendency for natural aging, but artificial aging is only used in special cases. Yield strength values of over 250 MPa can be achieved, but ductility is modest. The casting properties are relatively poor and they show a strong tendency to hot cracking [977]. On the whole, the alloys can be regarded as special and are not considered further in the next section.
<figure><img src="image_62.png" /><figcaption>**Fig. 60.** Schematic diagram summarizing potential trapping sites for contaminant elements at defects and precipitates. The influences of contaminant elements on macroscopic performance are also included.</figcaption></figure>

## 6.2. Main contaminants and contaminant effects in cast alloys

Cast alloys are sometimes more suitable for absorbing material from mixed alloy scrap streams than wrought alloys, which mostly have little tolerance for impurity elements and scrap-related compositional variations.

While a number of contaminating elements cause property changes which are comparable in all alloy groups, specific elements have particularly detrimental effects on some alloys. In the following, the most important alloy series are considered individually; where there are generally applicable impacts, this is explicitly pointed out. It is important to note that cast components frequently undergo post-processing heat-treatment operations which bring the microstructure from a non-equilibrium state to a state closer to thermodynamic equilibrium. In consequence, the effects of the contaminants may change (see Table 16).

By far, the most important contaminant introduced by scrap is Fe. However, Cu and Si can also have considerable contamination effects. It must therefore always be kept in mind that major elements in some alloys may be impurity elements in others. An important example is Cu, as Cu-free 3xx-series cast alloys for structural applications are often regarded as more corrosion-resistant than Cu-containing counterparts. However, Fragner et al. [978,979] pointed out that even a small increase in tolerable Cu content could lead to a considerable increase in potential scrap input – without any noticeable influence on corrosion properties. Si is Janus-faced in the 5xxx series. While in binary die-cast alloys it is considered a harmful element even in relatively low amounts if it forms Mg₂Si precipitates at grain boundaries, it can contribute to better castability if it is present in sufficient quantity to form an Al-Mg₂Si eutectic. Zinc is generally considered an acceptable impurity element in many scrap-based alloys. It is quite neutral in small concentrations, and neither improves nor diminishes the properties of an alloy.

Possible measures for reducing or eliminating property deterioration due to contamination and optional alloy design possibilities outside the standard compositions are discussed below. They address only the most important alloy series: 3xx, 5xx and 2xx. The assessment of Fe contamination is accorded the greatest importance.

The 3xx-series alloys 319 (AlSi6Cu3.5Zn) and A380 (AlSi9Cu3.5FeZn) are the most important scrap-based secondary alloys, and are primarily used in engine blocks and cylinder heads. Their usage volume equals over 80% of all die-cast alloys [971]. Because their impurity tolerance is significantly greater than that of other alloys, they are currently the most important sinks for scrap. Despite their wide compositional tolerance, these alloys are easy to process and exhibit good strength properties, even at higher temperatures up to 180 °C. However, their wide tolerance range is accompanied by great variations in strength and elongation values. In addition, a significant susceptibility to corrosion makes these alloys unsuitable for use in safety-critical components. Pucher et al. [99,980,981] systematically investigated the options for microstructure design and the associated property changes in AlSi9Cu3FeZn. The tolerance
<caption>**Table 15**</caption>

Aluminum Association four-digit numerical system for identification of aluminum alloys. The examples highlighted in color were considered in the preceding sections.

<table><thead><tr><td>Number series</td><td>Major alloying elements</td><td>Example alloy</td></tr></thead><tbody><tr><td>1xx.x</td><td>99% minimum Al content</td><td>Al99.7</td></tr><tr><td>2xx.x</td><td>Al + Cu</td><td>AlCu4Mg0.25 (206)</td></tr><tr><td>3xx.x</td><td>Al + Si + Mg, Al + Si + Cu, Al + Si + Mg + Cu</td><td>AlSi10.5Mg0.4<br>AlSi9Cu3MgFeZn (A380)<br>AlSi7Mg (A356)</td></tr><tr><td>4xx.x</td><td>Al + Si</td><td>AlSi10FeCuNi (A413)</td></tr><tr><td>5xx.x</td><td>Al + Mg</td><td>AlMg5Si2MnTi<br>AlMg5Zn3.5Cu0.5</td></tr><tr><td>7xx.x</td><td>Al + Zn</td><td>AlZn5.5Mg0.5</td></tr><tr><td>8xx.x</td><td>Al + Sn</td><td>AlSn6</td></tr></tbody></table>

range defined by the standards for all alloying elements presents wide scope for specific adaptations to the alloy composition to meet increased requirements, both for as-cast and age-hardened conditions. Depending on the chemical composition within the tolerance band chosen (8 &lt; Si &lt; 11; 2 &lt; Cu &lt; 3.5; 0.1 &lt; Mg &lt; 0.5, 0.15 &lt; Mn &lt; 0.55, 0.4 &lt; Fe &lt; 1.2; in wt.%), the yield strength and elongation to fracture of permanent mold and high pressure die-cast components in the as-cast state range from 100 MPa to 200 MPa and 0.35% to 3.9% respectively, and in the T6 state from 135 MPa to 400 MPa and 0.15% to 7% respectively (see Fig. 61).

As far as ductility is concerned, the morphology of the intermetallic phases which contain Fe is the decisive aspect, just as it is in wrought alloys (see section 8). The occurrence of the monoclinic β -Al₅FeSi phase must be prevented in favor of the cubic α-Al₁₅(MnFe)₃Si₂ phase. This can be achieved by adding so-called Fe-corrector elements, with Mn and Cr being the most important and effective ones [971].

Fig. 62 shows for the quaternary Al-Si-Fe-Mn system the effect of Mn on the Fe-containing IMCs for the case of a non-equilibrium solidification according to Scheil (as-cast condition) and for the equilibrium state at 495 °C (after solution heat treatment).

The tolerance bands for Mn and Fe of alloys 319 (Mn &lt; 0.5 wt%, Fe &lt; 1 wt%) and A380 (Mn &lt; 0.5 wt%, Fe &lt; 2% wt.) mean that greater Fe content will lead to the formation of the unfavorable primary β -Al₅FeSi phase there. It may be expected that at greater Mn content outside the tolerance band the preference for the α-Al₁₅(MnFe)₃Si₂ phase over the β -Al₅FeSi phase will result in improvements in ductility, although the total volume of IMCs which contain Fe will increase.

Hwang et al. [983] investigated the influence of Mn on the mechanical properties of alloy 319 containing 0.5 wt% Fe. They varied the Mn content in steps from 0.02 to 0.85 wt% and showed that the content of the β -Al₅FeSi phase was 0.5 vol% at 0.3% Mn but that no β -Al₅FeSi phase was present at 0.65% Mn (note that this is outside the tolerance range for Mn). Despite an overall higher volume fraction of IMCs of 0.29 vol% (α-Al₈Fe₂Si phase) versus 0.2 vol% (α-Al₈Fe₂Si + β-Al₅FeSi phase), the alloy with 0.65% Mn was significantly more ductile. Jin et al. [984] investigated the influence of the combined addition of Mn and Mo on the formation of β -Al₅FeSi and α-Al₈Fe₂Si phases in alloy 319, with 0.3 and 0.7 wt% Fe. They showed that adding Mo strongly promotes the formation of α-Al₈Fe₂Si and suppresses the precipitation of the β -Al₅FeSi phase in alloys containing both small and large amounts of Fe. However, the effect of Mo on the suppression of the β -Al₅FeSi phase depends on the Fe content. In alloys containing low Fe content, as little as 0.37% Mo completely suppresses the formation of the β -Al₅FeSi phase, whereas in alloys containing high Fe content only the pre-eutectic β -Al₅FeSi phase can be reduced as the Mo content increases, while the eutectic β -Al₅FeSi phase remains unchanged. Only the combined addition of Mn (0.24 wt%) and Mo (0.41 wt%) completely suppresses the β -Al₅FeSi phase in this case.

The influence of Cr and Co on the formation of α-Al₈Fe₂Si and β-Al₅FeSi phases in alloy A413 (AlSi10FeCuNi) with Fe content ranging from 0.5 to 2.5 wt% was investigated in [985]. The authors showed that an Fe/Cr ratio of 3 should be aimed for to eliminate the β -Al₅FeSi phase. For Co, an effect similar to that of Cr was observed, but a higher alloy content was identified in the formation of an α-Al₈Fe₂Si Chinese script instead of a β -Al₅FeSi phase. An optimal ratio of Fe/Co amounts to ≈ 1.0.

A team at Brunel University [986] investigated the effect of solution and aging on the microstructure and mechanical properties of AlSi9Cu3.5 die-cast alloys with ≈ 0.5 wt% Mn and a modified Fe content of 0.1 and 1.0 wt%. In the as-cast state the microstructures of the low-Fe alloy contained the intermetallic α-Al₈Fe₂Si phase, θ-Al₂Cu and Q-Al₂Cu₂Mg₈Si₆. The Fe-rich alloy also contained a β-Al₅FeSi phase and a higher α-Al₈Fe₂Si phase content. In the as-cast state the elongation to fracture of the Fe-rich alloy was reduced. The solution heat treatment caused the Cu-rich phases to dissolve and the eutectic Si phase to become spheroidized, resulting in significant improvement in mechanical properties after an additional aging treatment. Assuming that the Mn content of the alloys is higher, such that the Fe-rich variant would be in the sector of the β -Al₅FeSi transformation in Fig. 62, then an even more significant
<table><caption>**Table 16**

Influence of specific tramp elements on cast aluminum alloys.</caption><thead><tr><td>Impurity element</td><td>Undesired influence on alloy and processing</td></tr></thead><tbody><tr><td>Cu</td><td>Lowers corrosion resistance,<br>Lowers weldability</td></tr><tr><td>Fe</td><td>Forms coarse intermetallic phases,<br>Embrittling effect</td></tr><tr><td>Zn</td><td>Embrittling effect<br>Promotes oxidation of the liquid</td></tr><tr><td>Mg, Ca, Na</td><td>Promote oxidation of the liquid,<br>Lead to mold sticking</td></tr></tbody></table>

improvement in properties might be expected from high temperature annealing.

Pucher et al. [99] evaluated A380 alloys outside the tolerance band for Cu. As mentioned above, at high Si and Cu content not only α-Al<sub>8</sub>Fe<sub>2</sub>Si and β-Al<sub>5</sub>FeSi phases determine the ductility level, but also the intermetallic phases Q-Al<sub>2</sub>Cu<sub>2</sub>Mg<sub>8</sub>Si<sub>6</sub> and θ-Al<sub>2</sub>Cu. Systematic tests with Cu content reduced to slightly below 2.0 wt% generate a significant improvement in ductility. Permanent mold castings in the as-cast state exhibit an elongation to fracture of ≈10% at yield strength of 120 MPa.

These examples show that it is possible to increase the property profile of alloys 319 and A380 by partially bypassing the tolerance limitation. This potential for property improvement is particularly important for the following reason. Alloys 319 and A380 are the most important secondary alloys, for the synthesis of which almost all wrought alloys are ideally suited. This is best illustrated by Modaresi et al. [987] in Fig. 63. It shows the clear potential for recycling of many wrought alloy series into alloys 319 and A380 (this is indicated by the lighter shades in the red frame). However, the demand for alloys for internal combustion engines will decrease in the near future due to the electrification of vehicles (this topic is addressed in more detail in section 7.3). If variants are designed based on alloys 319 and A380, but with a balanced content of the main alloying elements and thus improved properties, they may however find applications outside of the field of combustion engines.

Of particular importance in the automotive industry are not only the casting alloys for aggregates described, but also those used in car bodies. Here the high ductility of the castings is absolutely essential for structural applications. The reference alloy for such applications is the primary alloy AlSi10MnMg, which is subject to challenging quality requirements (die-cast, T7 condition: yield strength > 120 MPa, elongation to fracture > 12%) and restrictions regarding the main impurity elements from scrap-iron, copper and zinc permitted. Consequently, the casting alloys used today are mostly based on primary aluminum. In all of them, Fe content is limited to values below 0.15 wt% to prevent the formation of the ductility-reducing β-Al<sub>5</sub>FeSi phase. Because increased use of secondary aluminum inevitably leads to increased Fe content, alloy design focuses primarily eliminating this IMP.

Boesch et al. [988] evaluated the optimum Mn content in age-hardenable AlSi10Mg0.4 alloys with ≈ 0.55 wt% Fe which would

<figure><img src="image_63.png" /><figcaption>**Fig. 61.** Mechanical properties and flow length of AlSi9Cu3(Fe,Zn) in as-cast and T6 states as a result of systematic variation of chemical composition within the alloy limitations. The range of properties is greatly extended by applying a heat treatment (solid solution treatment at 495 °C/8h + 160 °C/10 h). Blue: 0.4% Fe; red: 0.8% Fe; green: 1.2% Fe. Adapted from [981]. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)</figcaption></figure>
allow a high fraction of secondary aluminum. According to their systematic study, the Mn concentration most promising for neutralizing 0.55 wt% Fe in AlSi10Mg HPDC alloys is in the range of 0.4–0.6 wt%. A Fe/Mn ratio of slightly above 1 is particularly recommended, since an Mn content of 0.6 wt% leads to the formation of a finer morphology of the α phase compared to 0.4 wt% Mn. The associated mechanical properties are very satisfactory and meet requirements, with $R_{p0.2} \approx 150$ MPa and $\epsilon_f = 13\%$.

In [989] the influence of Mn and combined Mn plus Cr on the tensile properties of AlSi7Mg (A356) with 0.20 wt% Fe were investigated. With the addition of 0.20 wt% Mn, the acicular β-Al~5~FeSi phase is replaced by the granular α-Al~8~Fe~2~Si phase with the effect of a slight increase in strength and doubling of the elongation to fracture. Simultaneous addition of 0.13 wt% Mn and 0.13 wt% Cr also generates the transformation of the β-Al~5~FeSi phase but causes an almost threefold improvement in ductility. The mechanical properties of this Mn + Cr-modified Fe-enriched alloy variant exceed those of the commercial reference alloy A356 containing < 0.15 wt% Fe.

Basaz and Babu [871] investigated the effect of Cu on the microstructure and properties of a recycled AlSi6 cast alloy with 2 wt% Fe, i.e. the upper limit for the content of Fe impurities in recycled Al-Si alloys. They reported that the harmful β-Al~5~FeSi phase was destabilized by the addition of Cu, thereby apparently providing an alternative to traditional Mn addition. Solution annealing of variants with a Cu content > 4 wt% at 500 °C leads to the formation of the ω-Al~7~Cu~2~Fe phase instead of the β-Al~5~FeSi phase, which in turn decomposes into the θ-Al~2~Cu phase at lower temperatures. Although the mechanical properties obtained are rather modest – yield strength ≈ 100 MPa at elongation to fracture < 3% – the study points towards interesting options for the modification of secondary alloys with high Fe content.

Lin et al. [990] investigated the formation of IMCs and the mechanical properties of vanadium-modified AlSi7Mg cast alloys (A356) with 1.5 wt% Fe. Experiments which systematically varied V-content up to 1 wt% showed that an increasing V-content promotes the formation of an α-Al~15~(Fe,V)~3~Si~2~ phase and reduces the size of the β-Al~5~FeSi phase. With regard to mechanical properties, a V content of 0.8 wt% proves particularly favorable. A yield strength of almost 160 MPa is achieved at an elongation at break of just above 5%.

The above-mentioned scrap-related modifications of 3xx-series cast alloys have one aspect in common: they aim to prevent the formation of the plate/needle-shaped monoclinic β-Al~5~FeSi phase by adding transition metals in well balanced amounts. However, such alloy modifications have two further significant effects, firstly the reduction or elimination of “die soldering”, and secondly the formation of “sludge” particles.

Die soldering is a casting defect which is of considerable importance in high-pressure die-casting. It is caused by an interfacial reaction between the liquid aluminum alloy and the ferrous die-casting material [991]. A boundary layer is formed with various intermetallic compounds that adhere to the tool surface. This leads on the one hand to permanent damage of the tool surface, and on the other to the sticking of the castings to the mold, with negative consequences for process time and product quality. The tendency towards die soldering is clearly dependent on the chemical composition of the casting alloy. Increased content of Fe and Mn, but also of Cr and other Fe correctors, reduces the risk of die soldering [992–994]. In this respect, all measures for increasing ductility by transforming β-Al~5~FeSi into the α-Al~8~Fe~2~Si phase are also beneficial to the prevention of die soldering. This in turn implies that “chemically corrected” 3xx-series casting alloys may be particularly suitable for HPDC.

Fe correction (i.e. the addition of Mn and Cr to secondary cast alloys and their incorporation into the α-Al~8~Fe~2~Si phase), however,

<figure><img src="image_64.png" /><figcaption>**Fig. 62.** The isothermal section at 495 °C in the Al-Si7-Fe-Mn system illustrates the border between microstructures free of the β-Al~5~FeSi phase and β-containing structures depending on Fe and Mn content (solid line). In addition, the demarcation of areas with the primary β phase after non-equilibrium solidification according to Scheil is shown (calculated with the Pandat software using the database PanAl2019 [982]).</figcaption></figure>
increases the total fraction of IMCs and their thermal stability. Their liquidus temperature may exceed the casting temperature [995,996], with the result that large, primary α-Al₁₅(FeMnCr)₃Si₂ phase particles with polyhedral and star-like morphologies are already formed in the holding furnace [997,998]. Due to their high density (ρ ≈ 3.6 gcm⁻³) they fall to the bottom of the furnace as sediment, which in the long term creates process-related problems, or they are transported with the melt into the castings and create ductility-reducing inclusions [988,999]. To evaluate the tendency of alloys to form sludge and to avoid undesired phase formation, a segregation factor sludge factor was derived by Jorstad [999] (SF = Fe + 2 Mn + 3 Cr [wt.%]) which should not exceed a value of 1.8. Thus, the content of Fe, Mn and Cr in an alloy cannot be increased arbitrarily, because with sludge factor > 1.8 the occurrence of a coarse α-Al₈Fe₂Si phase becomes increasingly more likely [999]. The result is a significant reduction in ductility [988]. This effect should be taken into account when designing scrap-tolerant 3xx-series casting alloys.

The contaminant effects in 5xx-series alloys must be considered separately in the context of conventional binary alloys, conventional complex die-cast alloys, and recently developed die-cast variants for use in thin-walled structural components. As mentioned above, impurities should be avoided in corrosion-resistant binary alloys intended for decorative applications. These alloys usually have a polished surface whose optical quality is unacceptably reduced by second-phase particles. Thus Fe should be kept to a minimum, and Si content also requires particular attention. Thermodynamic calculations for solidification in non-equilibrium show that AlMg4.5 alloys with an Si content of 0.25 wt% already precipitate approx. 0.5 mol.% Mg₂Si from the residual melt [982], reducing ductility and corrosion resistance.

However, if increased strength with acceptable ductility is more important than high corrosion resistance, which is the case in many applications, increased scrap content may be considered. For this reason, Zhu et al. [1000] investigated the influence of Fe on the structure and properties of AlMg5 and AlMg5Mn0.6 alloys. The Fe content was varied systematically from 0 to 3.5 wt%. The Fe-rich intermetallics – monoclinic Al₁₃Fe₄ in AlMg5 and orthorhombic Al₆(Fe, Mn) in AlMg5Mn0.6 – influence the mechanical properties of the casting alloys significantly. Adding Fe increases the yield strength from ≈130 to 150 MPa, but significantly reduces elongation from 20 to 5%, independent of the Mn content. Nevertheless, elongations of >15% are achieved at an Fe content of 1 wt%.

The available industrial secondary 5xx die-cast alloys feature a very wide range of compositions, which are defined in country-specific specifications. At an average Mg and Mn content of 4.5 wt% and 0.5 wt% respectively, the concentrations of Si, Fe and Cu vary considerably in the range of 0.3 < Si < 1.5, 0.35 < Fe < 1.2, 0.03 < Cu < 1.0. The mechanical properties are accordingly rather modest. For example, for the alloy GK-AlMg5Si with an Si content between 0.9 and 1.5 wt% and an Fe content of 0.4 wt%, yield strength of 110–150 MPa and elongation to fracture of 2–5% were reported [1001].

Significantly more ductile variants of alloy type AlMg5Si2Mn have been successfully developed over the last few years [974,1002],

<figure><img src="image_65.png" /><figcaption>**Fig. 63.** Options for and restrictions on recycling paths of aluminum alloy series due to differing alloy content. The shades of gray indicate the percentage of a source alloy (scrap) that could be used in the production of the sink alloy. Reprinted with permission from [987].</figcaption></figure>
with the aim of enabling the production of large thin-walled structural parts. The optimum composition of the alloy is specified as 5.0–5.5 wt% Mg, 1.5–2.0 wt% Si, 0.5–0.7 wt% Mn, 0.15–0.2 wt% Ti. Rods of 6.4 mm diameter produced by HPDC achieve yield strength values of 140 MPa at 18% elongation. Cu increases the tendency towards hot cracking and affect ductility. A content of 0.4 wt % already reduces the elongation to fracture by a factor of 2. An Fe content of up to 0.25 wt%, however, can be tolerated, because the increased Mn content enables the formation of a compact α-Al<sub>8</sub>Fe<sub>2</sub>Si phase [974]. At higher content the elongation to fracture decreases unacceptably. Adding 0.1–0.3 wt% Cr hinders the tendency towards die soldering [1003]. Higher contents should be avoided, as otherwise an Al<sub>7</sub>Cr phase is formed, which significantly reduces ductility [1004]. In general, the effect of contaminants on the properties of alloy AlMg5Si2Mn can be regarded as rather moderate. However, during smelting it is important to reduce the content of alkali and alkaline earth elements as much as possible. Calcium and sodium are limited to a maximum of 15 ppm, as these increase the tendency of castings towards hot cracking and lead to reduced fluidity [975]. When using recycled material, monitoring the calcium and sodium content is strongly recommended.

Alloy design based on the crossover principle is proving an interesting method (see section 8.4). One example is new high-pressure die-casting alloys based on Al-Mg-Zn-(Cu) for structural cast components [1005]. Compared to the binary AlMg4.5 variants, the scrap-relevant elements Zn and Cu are also taken into account here as main alloying elements. A decisive advantage of these alloys is their potential for artificial aging. The experimental alloy with 4.8 %Mg, 3.4 %Zn, 0.8 %Mn and 0.5 %Cu is a good example. After pre-aging at 100 °C/3h, an aging treatment at temperatures in the range of 150 to 175 °C leads to a considerable increase in hardness within an industrially relevant time range of 2 to 20 h. The formation of T-phase (Mg₃₂(Al, Zn)₄₉) precursors is responsible for the age hardening [1005,1006].

The influence of Fe and Si contamination on the structure and properties of 2xx-series alloys can be best illustrated using the example of alloy 206 (AlCu4.5Mg0.25). Again, Fe is a property-determining impurity and is normally limited to 0.15 wt%. In addition, Si content is restricted to < 0.1 wt%. A study by Liu et al. [1007] showed how the Fe-containing IMCs are modified when alloys containing 0.15 wt% Fe are treated with Mn and Si. Adding either Mn or Si or both promotes the formation of the α-Al<sub>8</sub>Fe<sub>2</sub>Si phase and inhibits the occurrence of the β-Al<sub>5</sub>FeSi phase. The cooling rate also has a positive effect on β - to α-Al<sub>8</sub>Fe<sub>2</sub>Si phase conversion. At 0.3 wt % Mn and 0.3 wt% Si the formation of the β-Al<sub>5</sub>FeSi phase is completely prevented. In a further work [1008], the same group is examining Al-Cu 206 cast alloys with different Fe content in the range of 0.15 to 0.5 wt% at Si and Mn content of 0.1–0.3 wt% and 0.1–0.5 wt%, respectively. At comparable Si and Mn content the elongation to fracture decreases with increased Fe content due to the increasing volume and size of the Fe-rich intermetallic compounds. With a simultaneous Si and Mn content of 0.3 wt% each, quite good mechanical properties are achieved in condition T4, with a yield strength of 260 MPa and an elongation to fracture of ≈8%. Zhang et al. investigated the influence of Fe in the range 0.1 – 1.5 wt% for AlCu5 alloys with 0.6 wt% Mn and 0.1 wt% Si. In the squeeze-casting process, very good properties are achieved with an Fe content of 0.5 wt% (R<sub>p0.2</sub> ≈ 220 MPa, A≈18%). The latter two studies may serve as incentives in the development of contaminant-compliant 2xx-series casting alloys.

6.3. *End-of use recycling of engine block alloys and gearbox housings*

It can be assumed that the demand for cast components as aggregates of passenger cars will decrease as the electrification of private transport progresses [62]. Engine blocks and gearbox housings, for example, are nowadays made from composition-tolerant recycling alloys, mainly from the aforementioned alloys 319 and A380, for which there is no other relevant use requiring large quantities. If such cast components are returned to the scrap cycle at the end of the vehicles' lives, it is predictable that there will no longer be any direct recycling use for them due to their high Si, Cu and Fe content. This brings up the question of possibilities for metallurgical modification with the aim of alternative use in structural components subject to increased demands on mechanical and electrochemical properties. Three options can be mentioned here: (i) alloy and process design of scrap-tolerant alloys; (ii) alternative manufacturing processes which allow rapid solidification; and (iii) removal of contaminants by physical methods with reuse of the purified fraction.

(i) Alloy and process design were already mentioned above, and demonstrated with two examples. The first suggestion was to reduce the Cu content of A380 while balancing the Fe/Mn ratio. Such measures, however, require a willingness of industry and customers to break away from traditional standardization constraints. The second measure was the use of thermal treatment with the aim of dissolving a portion of the IMCs and spheroidizing the eutectic Si [98,986]. A prerequisite for this, however, is low gas content in the components to avoid blistering, which would require a vacuum-supported high-pressure die-casting process.

(ii) Alternative manufacturing processes which allow rapid solidification offer the advantage of refining the IMCs, which with their smaller size and reduced aspect ratio have fewer ductility-reducing effects. Examples are processes for additive manufacturing (AM) and the spray forming technique. The possibilities of AM are described in detail in section 9.1. In particular, a study by Suchy et al. [1009] is referred to there in which A380 powder was compacted by laser powder bed fusion, resulting in good mechanical properties. Here a brief reference is made to the experience with spray forming (Osprey). The recycling of contaminated casting alloys using the spray-forming technique is not a new approach. A series of publications show that the spray-forming production of Al-alloys with acceptable mechanical properties can be achieved even where there is high iron content [269]. Ferrarini et al. [268] investigated the structural and property changes of spray-compacted A380 and show that the spray-forming process can strongly suppress the formation of the β-Al<sub>5</sub>FeSi phase. However, increased porosity levels are an inherent problem in spray forming, necessitating subsequent treatment by extrusion or forging to achieve ductility properties for structural applications [1010]. This significantly reduces the attractiveness of the process.

(iii) On the subject of removal of contaminants by physical methods, two techniques are addressed below which enable significant reduction in harmful contaminants. The first covers the separation of the solid phase and melt of alloy A380 in the semi-solid state, and the second the separation of sludge-forming IMCs by means of gravity or supergravity segregation.
In hypoeutectic alloys 319 and A380 in the semi-solid state, the chemical composition of the solid fcc-Al phase differs significantly from the basic composition, since the majority of the IMCs are precipitated in the melt. For example, for an AlSi8Cu3.5Mn0.55Mg0.5 alloy with 0.8 wt% Fe, the fcc-Al phase at 560 °C would have a mass fraction of ≈ 0.5 and a chemical composition of 1.38 wt% Si, 1.12 wt% Cu, 0.17 wt% Mg, 0.02 wt% Mn and only 0.003 wt% Fe (calculated with Pandat software using database PanAl2019 [982]). If the solid fcc phase is mechanically separated from the melt and the IMCs, considerable cleaning performance is achieved. Such a method is described in the Austrian patent specification AT412348B [1011]. It is based on the principle of mechanical squeezing of a liquid-filled sponge.

Sludge removal by gravity separation is based on the difference in density between sludge and melt. The literature reports on Fe reduction in Al-Si alloys by the intentional addition of Mn and Cr to form sludge (see section 6.2), which is then separated from the melt by physical sedimentation or centrifugation, but also by electromagnetic sedimentation or filtration [1012-1016]. At this point, instead of detailing the various processes, an estimation of the purification potential of sludge removal in alloy A380 is useful. It is based on thermodynamic calculations using Pandat with the database PanAl2019 [982]. For alloy A380 (AlSi8Cu3.5Mn0.55 Mg0.5Fe0.8) at 620 °C, which is about 30 °C above the liquidus temperature of the fcc-Al phase, Table 17 shows the calculated values for the mass fraction of the α-Al<sub>8</sub>Fe<sub>2</sub>Si phase (V<sub>α</sub>, sludge to be removed) and for the composition of the melt (without the α-Al<sub>8</sub>Fe<sub>2</sub>Si phase). The data are given for the original alloy with 0.8 wt% Fe and 0.55 wt% Mn, and after addition of 1.0 wt% Mn and 1.5 wt% Mn. The calculated data indicate that optimal process control of the sludge removal may enable the Fe level of the secondary alloy to approach that of a primary alloy.

# 7. Effects of scrap-related impurities in wrought alloys: Constitution, microstructure, processing and properties

## 7.1. Brief introduction to wrought alloys

Unlike cast aluminum alloys, wrought alloys have narrow and sometimes low compositional limits. This means that wrought alloys do not easily absorb mixed alloy scrap streams and that primary aluminum remains an important backbone for them [72,171,172].

The alloying elements used in the wrought aluminum alloy series 1xxx to 8xxx are listed in Fig. 64. For typical alloys within the different classes, major, common minor and less common minor alloying elements according to [1017,1018] are indicated. Elements are considered to be intentionally added if there are a lower and an upper limit in the composition range of the standard. Based on the allowable compositional ranges, some critical elements can be specified (Fig. 64) and these can give a first indication of whether or not a wrought alloy series is easy to recycle and tolerates impurities. The 1xxx, 7xxx and 8xxx series are considered “tough” (to produce from scrap) due to the compositional restrictions set by current standards. The 2xxx and 6xxx series can be judged as “challenging”. Only the 3xxx and 4xxx wrought alloys exhibit a “moderate” challenge for production from scrap. The most problematic elements according to this evaluation are Fe, Si, and Cu, and these are well known in the aluminum recycling community [51].

Ashkenazi et al. [1019] provide a rather long list of problematic impurities from recycling which includes Cr, Cu, Fe, Mg, Mn, Ni, Pb, Si, V and Zn. Their list is consistent with Fig. 64, but also includes minor common alloying elements as critical (here Pb, Ni, V). From this it is important to acknowledge that not only the main elements present in the different alloy classes can cause problems upon recycling; even trace elements present in specific alloys can have drastic effects on other wrought alloys. For example, adding a few parts per million of Sn can dramatically suppress or accelerate aging behavior in 6xxx [553,555] and 2xxx series alloys [383] respectively. Nowadays Sn is used as a substitute for poisonous Pb in some alloys for machining operations, to enhance chip breaking [341,1020]. Other exotic elements are also often used in 8xxx-series alloys, which are sometimes referred to as miscellaneous alloys [970]. However, the introduction of new or exotic alloying elements in alloy development needs to consider their impact on other alloy classes and the implications for recycling. Most recent studies have considered such impacts from a very general perspective without examining the detailed metallurgical effects such elements may have [71,1021]. In the recycling of scrap from the automotive sector a comprehensive analysis was in fact performed to examine whether Ce, which may be used in some new casting alloys, will impact the future properties of alloys produced from recycling [1022]. However, up to now this type of study seems to be the exception rather than the rule.

Fig. 65 displays the tensile properties covered by wrought aluminum alloys. A broad range of properties are covered by classes 1xxx to 8xxx. Alloys where the properties are most sensitive to contamination by tramp elements are indicated by colored group contours.

There is no obvious direct correlation between the properties presented in Fig. 65 and sensitivity to tramp elements. However, it is clear that the alloys providing the extreme values for yield strength or elongation are subject to “tough” compositional restrictions in the context of recycling. With decreasing yield strength and elongation values, the compositional restrictions change from “tough” to “challenging” and finally to “moderate”. In other words, the most-developed alloys from the material point of view are also the most difficult to recycle. This may indicate an inherent connection, or simply that recyclability was not explicitly considered by materials scientists in alloy design over the last few decades.

The alloy series where the impurity content is critical comprises the high-strength alloys often used in aircraft [78]. The content of Fe and Si as tramp elements is especially crucial in the 7xxx series, but also in the 2xxx series of alloys frequently applied in the aircraft industry [1017]. Besides the high-strength alloys, very soft alloys like the 1xxx series are also highly sensitive to elements such as Mg (see Fig. 64). Those alloys are typically used in foils (e.g. for food equipment), decorative and gloss applications, heat shielding plates or electric applications [1017,1018]. Some alloy series exhibiting medium strength and elongation values are less sensitive to tramp elements (Fig. 64). These mostly involve 3xxx-series alloys, which are used heavily in beverage cans [1017,1018]. 3xxx-series alloys are also utilized in architecture and as the core material in heat exchangers, where 4xxx series alloys are often applied as clad material for soldering applications [1018]. In contrast, other lower-to-medium-strength alloys such as the 5xxx series alloys feature tight Si
<caption>**Table 17**</caption>

Thermodynamically calculated values for the mass fraction of the α-Al<sub>8</sub>Fe<sub>2</sub>Si phase, V<sub>a</sub>, and the chemical composition of the purified melt [in wt.%]; alloy A380 with 8 %Si, 3.5 %Cu, 0.5 %Mg, 0.55 %Mn and 0.8 %Fe at 620 °C.

<table><thead><tr><th></th><th>V<sub>a</sub></th><th>Si</th><th>Cu</th><th>Mg</th><th>Mn</th><th>Fe</th></tr></thead><tbody><tr><td>A360 (0.8%Fe, 0.55%Mn)</td><td>0.0153</td><td>7.97</td><td>3.55</td><td>0.51</td><td>0.30</td><td>0.60</td></tr><tr><td>+ 1.0 at.% Mn</td><td>0.054</td><td>7.88</td><td>3.70</td><td>0.53</td><td>0.50</td><td>0.28</td></tr><tr><td>+ 1.5 at.% Mn</td><td>0.071</td><td>7.83</td><td>3.77</td><td>0.54</td><td>0.57</td><td>0.20</td></tr></tbody></table>

limits and, in several cases, tight Cu limits. These alloys are typically deployed for inner panels in automotive applications or for other structural parts in transportation, pressure vessels, boat hulls and ballistic armour plates [1024–1026]. The 6xxx-alloy class is judged to be moderate according to Fig. 64, but actually shows diverse tolerance limits if examined in more detail. Fortunately Si is a major alloying element in 6xxx-series alloys and thus tight compositional control is not critical. However, some alloys exhibit strict Fe tolerance limits, and others strict Mn tolerance limits. This strongly depends on the actual alloy application, which ranges widely for the 6xxx series and encompasses transportation, automotive industry and mechanical engineering [1017,1027,1028].

Combinations of different wrought alloys series are deployed in several fields of application. In the following we discuss how wrought-alloy diversity impacts potential recycling in three specific cases: beverage cans, automobiles and aircraft.

Recycling of old scrap for aircraft alloys seems very difficult, because the alloys involved are typically in use for a long period and tolerance limits are very strict (especially in the 7xxx series, but also in some 2xxx wrought alloys), excluding nearly all impurity elements (Fig. 64). One reason for the low tolerance of Fe and Si impurities, which can form brittle IMCs, are aircraft applications’ stringent fracture toughness requirements [1029,1030]. Manufacturing for the aircraft industry usually involves complex machining operations, and typically a high proportion of the material is lost during production (this is known as the “buy to fly ratio”) [3,1031]. Here closed-loop recycling may be an option [1032]. There the challenge is more logistical than metallurgical. Nevertheless, the issue of old-scrap recycling will also arise in this area, and solutions must be found to the problem of narrow tolerance limits for tramp elements.

Aluminum beverage cans present an extreme contrast to the situation described above. Today packaging is the second-largest source of aluminum scrap, and here cans are a well suited example to the transition towards a circular economy [30]. Can usage time is relatively short, and the recycling of used beverage cans (UBC) is well established [30,299]. In some countries with dedicated collection systems (i.e. Finland [30]) the recycling rate is as high as 97%. The composition of UBCs is very suitable for recycling because few alloys are involved (i.e. 3004 for the body and 5182 for the lid [30]). Therefore recycling is simple as to composition; the fact that the main alloy system (3xxx) is moderately tolerant of foreign elements (Fig. 64) is also helpful. Other types of packaging scrap may even be deployed for cans. However, according to recent work can-to-can recycling seems the better option with regard to climate change [30].

<figure><img src="image_66.png" /><figcaption>**Fig. 64.** Alloying elements in wrought aluminum alloys and their critical limits (sorted according to the most common alloys given in [1017,1018]).</figcaption></figure>
The third complex, challenging and most important example of aluminum in terms of recycling is automotive. Today's automotive industry is already the largest consumer of secondary aluminum due to the huge number of used castings that are relatively easily recycled into other castings [72]. All car manufacturers have increased their use of aluminum over the last few years because it is lightweight, and this is forecast to increase further (see Fig. 66) [72].

At the same time, however, vehicles' wrought alloy content (sheets, forging and extrusion) is strongly increasing [44,71,72,987,1033]. This is illustrated in Fig. 67, which shows the tremendous growth in scrap potentially available from automotive body sheet (ABS) material [36]. This colossal wave of ABS scrap requires us to think carefully about appropriate recycling strategies. This is even truer if we consider that Zhu et al. [36,113] calculated that in the best-case scenario in 2017, only 70% of aluminum scrap could be recycled domestically in the US due to compositional constraints. Currently the issue is partially masked by scrap exports (in particular to China, and concerning mainly shredded auto scrap). A further increase in demand for wrought alloys, to produce power-supply boxes in electric vehicles, is predicted [168].

Without a technological leap, the huge increase in wrought products in the automotive sector may become incompatible with current casting-focused recycling of end-of-life vehicles (ELVs), or even generate excess scrap by reducing automotive casting demand (at present the automotive industry is the biggest consumer of secondary castings) [35,72]. The complexity of the problem is amplified because data is lacking on the alloy quantities and types needed in future ELVs [72]. This is a general problem, because most data sources list only the total amount of aluminum produced in wrought, cast or extruded form, and data on specific alloys are seldom available [72]. Table 18 shows some typical alloys used in various different aluminum components in cars [72].

In Table 18, while the problem of mixed cast and wrought alloys is not addressed, the large number of different wrought alloys used in cars is immediately apparent. 1xxx, 2xxx, 3xxx, 4xxx, 5xxx, 6xxx and 7xxx alloys have all been deployed there. However, the most frequent automotive wrought alloys are the 6xxx, 5xxx and 3xxx series [72]. Besides the fact that contamination from cast alloys is problematic for most wrought alloys, recycling the wrought material from cars into new wrought alloys is a great challenge [36]. This can be deduced by comparing alloying limits (Fig. 64). A good example considers the mix of 6xxx and 5xxx series alloys, which are typically tightly joined as inner (e.g. 5182) and outer (e.g. 6016) car door parts [1034]. Although Mg is a major alloying element in 6xxx alloys, the Mg content limit can be more than one order of magnitude lower in 6016 series alloys than in 5182 series alloys [1017,1018]. Even more problematic is the very low Si limit in 5xxx series alloys: Si content can be >7 times higher in 6016 than in 5182 [1017,1018]. The metallurgical origins of these tight controls are discussed in subsequent sections.

Drastic improvements in sorting technology are clearly mandatory if wrought alloys are to be recycled efficiently to new wrought alloys in ELVs. Full separation of alloys is not currently possible, because shredder-based recycling practices show natural limits when faced with complex product designs and the various associated joining means [1035]. Hence impurity levels will rise along with the increase in the number of different alloying elements, because of the various types of wrought alloys used [72]. The growing number of alloys used in cars will generate a concurrent increase in joining means such as screws, rivets and adhesive bonding, while welding will decrease [116]. The trend towards multi-material design will reduce the potential of higher-purity wrought alloy recycling streams and increase the proportion of material with a higher Fe content [116,1036,1037].

IMCs which contain Fe can be present in most wrought alloys in the form of brittle needles or plates, and are known to cause detrimental effects such as crack initiation [78,252], reduced ductility and bendability [322], reduced fracture toughness [1017] and

<figure><img src="image_67.png" /><figcaption>**Fig. 65.** Comparison of tensile properties for 1xxx to 8xxx wrought aluminum alloys. Alloys where properties are most sensitive to contamination by tramp elements according to the evaluation in Fig. 64 are indicated [1023].</figcaption></figure>
deteriorated surface finish [1038]. However, in wrought aluminum alloys Fe is not solely detrimental. In 1xxx and (especially) 8xxx series alloys, the situation is reversed: high fractions of Fe-containing IMCs are required to control the grain size and properties of foil [1039].

Despite their importance, the effects of Fe-containing IMCs have not been studied as intensively in wrought alloys as in cast alloys, because most contaminated old scrap usually goes for the production of cast alloys. This section briefly addresses only wrought alloys or the basic principles of Fe-containing IMCs. Details on the effects of specific Fe-containing IMCs can be found in the corresponding sections on specific wrought alloy classes.

Fundamental research on Fe-containing IMCs, such as the recent first-principles study by Fang et al. [1040] on the solubility of Si in some Fe-containing IMCs, is gaining increasing importance in light of the implications for recycling. Feng et al. [1041] also recently studied the formation of primary Fe-containing IMCs using in-situ X-ray radiography. They report a correlation between inoculation by TiC and TiB₂, the solidification conditions and an increased number density of Fe-containing IMCs [1041]. Similar to the basic studies of the Al-Fe system in [1041–1043], there have been attempts to better control the size, distribution and morphology of Fe-containing IMCs [1044–1047] by understanding the fundamentals of nucleation, growth and phase selection, in order, potentially, to facilitate better toleration of higher Fe content in wrought alloys [1044–1050]. Efforts have also been made to search for potent nuclei for primary Fe-containing IMCs, at least for cast alloys [175]. However, while some authors reported a weak inoculation effect of α-Al₂O₃, TiC and TiB₂ [1049] in experimental alloys near the 1xxx series, Lui et al. [1046] suggested that TiB₂ might be associated with nucleation of Fe-containing IMCs in 6xxx alloys. Many questions remain in this important field of phase selection.

Another element whose limit is often considered critical in wrought alloys is Si (Fig. 64). Si is present as a major element in cast alloys to improve casting properties. It is also seen in the 6xxx series of wrought alloys, but can have severe negative effects in other classes. Highly sensitive alloys include the 7xxx and 5xxx series, because there Si can generate large, brittle IMCs. For example, if Mg content is high, the primary Mg₂Si phase will form even at low Si content [78]. Primary Mg₂Si particles have been shown to initiate fracture even at low strains [877].

A third element that is critical in some wrought alloys is Cu (Fig. 64). Cu is the main element for age hardening in 2xxx series wrought alloys, and also promotes strength in other alloy classes such as the 6xxx and 7xxx series. However, Cu can negatively affect corrosion properties [1017]. This problem is often discussed in the context of the automotive applications of 5xxx and 6xxx series alloys. However, many 6xxx series alloys used in the automotive sector contain Cu as an alloying element. Even incorporation of Cu in automotive 5xxx series alloys to mitigate strength loss upon paint baking has recently been examined [1051]. It is true that some corrosion properties can deteriorate here, but they can also be kept above a level satisfactory for the application; and Cu can even improve performance in the area of specific corrosion types [1052]. A similar effect is known in 7xxx series alloys, where adding Cu mitigates stress corrosion cracking [1053]. Nevertheless, it is a fact that many wrought alloy standards place tight restrictions on Cu.

In conclusion, it must be emphasized that in general there is no such thing as a “bad element” in wrought aluminum alloys. All

<figure><img src="image_68.png" /><figcaption>**Fig. 66.** Average vehicle curbside weight and the aluminum content in cars in North America. Reprinted with permission from [72].</figcaption></figure>
<figure><img src="image_69.png" /><figcaption>**Fig. 67.** Estimation of the ABS scrap generated by vehicles with the highest US sales: Ford F-150, Super Duty, Expedition, and Lincoln Navigator. Reprinted with permission from [36].</figcaption></figure>

<table><caption>**Table 18**
Typical alloys utilised by the automotive industry. Reprinted with permission from [72].</caption><thead><tr><th>Body & Inner Panel</th><th>Body Closures</th><th>Heat Exchangers</th><th>Heat Shields</th><th>Misc. Engine</th></tr></thead><tbody><tr><td>2008, 5030, 5052, 5182,<br>5454, 6009, 6016, 6111</td><td>2008, 2036, 6009, 6016,<br>6010, 6383, 6061, 6111</td><td>6060, 6061, 6063, 6106, 5049, 7072, 1145, 4047, 4004,<br>4045, 4343, 3003, 8079, 6006, 1200, 1050, 1100</td><td>1056, 3003,<br>5052, 5182</td><td>226, AlSn20Cu,<br>AlZn5Bi4</td></tr><tr><td><strong>Cradles & Frames</strong></td><td><strong>Wheels</strong></td><td><strong>Steering System</strong></td><td><strong>Fuel System</strong></td><td><strong>Engine/<br>Cylinders</strong></td></tr><tr><td>5182</td><td>356, 6081, 6061</td><td>6082, 7108, 7021</td><td>6063, 3103,<br>5049, 5754</td><td>380, 319, Al-Si</td></tr><tr><td><strong>Collision</strong></td><td><strong>Brake System</strong></td><td><strong>Suspension Parts</strong></td><td><strong>Trans</strong></td><td><strong>Pistons</strong></td></tr><tr><td>6013, 7021, 7029</td><td>359 or 360 + SiC</td><td>AlSi7Mg, 6013, 6082</td><td>380.2</td><td>4032</td></tr></tbody></table>

elements introduced by recycling can have positive effects. Even the impurity most often discussed in the literature, Fe, is beneficial in 8xxx series alloys. It is, rather, the combination of a certain wrought alloy with a certain tramp element that causes problems. A holistic view of the effects of elements in alloys can help to allay the specific problems caused by some currently unavoidable tramp elements in wrought alloys.

## 7.2. Use of scrap in 1xxx and 8xxx alloys for electric, decorative, chemical and packing products

As Fig. 65 indicated, the 1xxx and 8xxx series alloys for electrical, decorative, chemical and packaging products are very sensitive to many accompanying elements. Table 19 shows their typical compositions and respective uses. 1xxx alloys are subject to very tight limits on most of the important alloying elements used in other wrought alloy classes. These alloys include Mg, Cu, Mn and Zn, which are often kept below 0.05 wt% in alloys used in food equipment, decorative and shiny applications and heat protection plates [1018,1054]. It is thus immediately clear that given such strong restrictions on impurities, no other wrought alloys recycled from old scrap can be used in these types of product. Hence production is highly dependent on primary aluminum, and only alloy-to-alloy recycling (e.g. from production scrap) is an option. This also applies to some electrical and electronic products which use 1xxx alloys with high-purity aluminum (e.g. 1199 with 99.99 wt% Al) [1054].

In the past, alloys of the 1xxx series (e.g. 1145, 1100, and 1200) were widely used for foil stock [1054]. However, due to increased demands on strength and ductility higher levels of Fe (up to 2 wt%), Si (up to 1.1 wt%) and sometimes Mn (up to 1.0 wt%) are now often used in foils to control the microstructure during recovery and recrystallization via IMCs, which turned those alloys from 1xxx to 8xxx series alloys [1039,1054]. High levels of Fe increase the strength of foils (unlike that of all other wrought products), and generates ductile material behavior via IMCs. Fe acts as effective grain refiner (highly important in thin foils) [1039]. Fig. 68 shows one of the rare cases where Fe is a welcome element. Here Fe acts in exactly the opposite way than it does in aluminum alloys used for structural applications. Instead of limiting ductility, it increases elongation when its amount is increased. However, this is an indirect effect and results mainly from the increase in the volume fraction of IMCs acting as grain refiners upon recrystallization, as the maximum Fe content in the solid solution is only 0.05 wt% [1039].

In the recycling of 8xxx alloys used for foils, the situation is essentially the same as for 1xxx alloys. The limits on elements other than Si and Fe are only in rare cases slightly less stringent than in 1xxx alloys, and production from old scrap via recycling is difficult.

A more differentiated picture emerges when considering the use of these 1xxx and 8xxx alloys in recycling for the production of other alloys. The high-purity and 1xxx alloys are the ideal starting material in terms of composition and are comparable to primary aluminum in terms of composition. However, recycling then takes place as downcycling in a cascade recycling chain, with such low-
alloy aluminum at the top of the cascade [169]. However, with increased amounts of Fe and Si, using 1xxx and especially 8xxx alloys becomes more challenging (compare with Fig. 65). Although mixed packaging waste might be deployed in (e.g.) can bodies, Niero et al. saw can-to-can recycling as more sustainable [30]. In this context, it is important to know that the EU has a target of recycling > 75% of packaging waste in the next few years [30].

Packaging scrap from households is typically sorted into ferrous and non-ferrous alloys, resulting in a mixture of cast and wrought alloys which is mostly absorbed in cast alloy production. This results in continuous downcycling of the wrought scrap fraction [31]. According to Paraskevas et al. [169], container and packaging scrap might theoretically be managed using a separate closed-loop recycling strategy for the same applications, which may solve the problem.

Different alloys are often used in rather similar products, increasing recycling complexity and making scrap separation in particular more challenging. A good example are wrought alloys used for heat exchangers, where (among others) both 1xxx and 8xxx alloys are deployed. The variety of alloy types in heat exchanger components shown as an example in Fig. 69 necessarily produces a large number of tramp elements [72].

The 1xxx and (even more) the 8xxx alloys intentionally contain Fe, an element which is typically detrimental to most other wrought alloy series. Allen et al. [1048] summarized the various stable and metastable phases which can be present in 1xxx alloys. Although only the phase Fe₄Al₁₃ (often denoted as FeAl₃, but structurally complex with 100 atoms / unit cell) exists at dilute conditions in binary Al-Fe in the thermodynamically stable state, [1048] lists at least 6 structurally distinct phases in that form under industrially standard non-equilibrium solidification conditions. The presence of Si further increases the phase complexity of the compositionally simple 1xxx alloys. Eight structurally different phases are listed for the common α- and β-Al₅FeSi phases [1048]. Fig. 70 shows the composition ranges of some of the most common phases in dilute Al-Fe-Si alloy systems.

It is significant that the phase selection can be influenced by varying the casting conditions (local solidification velocity, cooling rate, grain refinement). In compositionally simple systems the underlying mechanisms of nucleation and growth in Fe-containing IMCs have been much investigated. Because it is important to control the size and change the morphology and distribution of Fe-containing IMCs [1048,1049,1055,1056], existing basic information on 1xxx alloys may be an interesting source of knowledge in the effort to better control Fe-containing IMCs in other wrought alloys with high recycling content. Further properties such as corrosion have also been investigated in these alloys. For example, Dorin et al. [1057] showed that an increased cooling rate upon solidification has a positive influence on corrosion properties. Therefore direct strip casting was concluded to be an interesting candidate technology for improving the recyclability of Fe-containing Al-alloys, as it enables high solidification rates.

Several more exotic manufacturing and industrially less-common methods were studied in the context of Al-Fe alloys with high Fe content. This includes work on rapid solidification [1058–1060], mechanical alloying [1061–1063] and severe plastic deformation [1064,1065]. Such work may also provide useful knowledge for the future recycling of high-Fe aluminum scrap.

## 7.3. Use of scrap in 2xxx alloys intended for high-strength components

### 7.3.1. Introduction

With the development of the duralumin alloy, the Al-Cu-based 2xxx alloy family gained the first age-hardenable aluminum alloys used in high-strength aerospace applications [1066]. Today this alloy family comprises three main sub-types, based on their combination of solute elements: (i) Al-Cu based alloys, strengthened by the θ-Al₂Cu precipitation sequence including GP zones, θ and θ' phases (typical alloys: 2219, 2017) [1067]; (ii) Al-Cu-Mg alloys, strengthened by the S-Al₂CuMg precipitation sequence including clusters, GPB zones, and the S phase (typical alloys 2024 and its variations 2124, 2524, 2624) [1068,1069]; and (iii) Al-Cu-Li-(Mg) alloys, strengthened by combinations of phases, including T₁-Al₂CuLi, δ'-Al₃Li, and the phases of the two aforementioned sequences (typical alloys 2195, 2198, 2099, 2070) [1070–1072].

### 7.3.2. The role of impurities in 2xxx alloys: Chemistry effects

2xxx series alloys are generally used in high-strength components. The Li-containing alloys reach strengths comparable to those of 7xxx alloys. Al-Cu-Mg-based alloys are frequently used in tension, so their damage tolerance is of paramount importance. The fatigue resistance of these alloys is also one of their key properties in such conditions. Al-Cu-based alloys are used primarily at elevated temperatures, with high Cu levels (up to 6 wt%).

In all these alloys, evidence has shown that coarse intermetallic particles are the primary cause of fracture [891,1073–1075] both during initiation (short crack nucleation) and final fracture (toughness-controlled). The damage tolerance has been shown to be degraded when the amount of Fe in particular is increased [1076–1078]. This effect may be linked to the phase diagram of the Al-Cu-Fe

<table><caption>**Table 19**
Typical compositions and uses of 1xxx and 8xxx wrought alloys [1054].</caption><thead><tr><th>Designation</th><th>Si</th><th>Fe</th><th>Cu</th><th>Mn</th><th>Mg</th><th>Zn</th><th>Cr</th><th>Ti</th><th>Comments/Typical application</th></tr></thead><tbody><tr><td>1100</td><td>1 (Fe + Si)</td><td></td><td>0.05–0.20</td><td>0.05</td><td></td><td>0.10</td><td></td><td></td><td>Al > 99/Foils, sheets, plates, tubes, wires spun hollow-ware, food equipment</td></tr><tr><td>1200</td><td>1 (Fe + Si)</td><td></td><td>0.05</td><td>0.05</td><td></td><td>0.10</td><td>0.05</td><td></td><td>Al > 99/Foils,</td></tr><tr><td>1145</td><td>0.55 (Fe + Si)</td><td></td><td>0.05</td><td>0.05</td><td>0.05</td><td>0.05</td><td>0.03</td><td></td><td>Al > 99.45/Foils, sheet, semi-rigid containers</td></tr><tr><td>1199</td><td>0.006</td><td>0.006</td><td>0.006</td><td>0.002</td><td>0.006</td><td>0.006</td><td>0.002</td><td></td><td>Al > 99.99/Electrical and electronic foil</td></tr><tr><td>8011</td><td>0.50–0.9</td><td>0.6–1.0</td><td>0.10</td><td>0.10</td><td>0.05</td><td>0.10</td><td>0.05</td><td>0.08</td><td>Household foils, bottle caps,</td></tr><tr><td>8006</td><td>0.40</td><td>1.2–2.0</td><td>0.30</td><td>0.30–1.0</td><td>0.10</td><td>0.10</td><td></td><td></td><td>Foils, finstock for heat exchangers</td></tr></tbody></table>
<figure><img src="image_70.png" /><figcaption>**Fig. 68.** Effect of Fe content on strength and elongation of 8021 foil in the O-temper. Reprinted with permission from [1039].</figcaption></figure>

system [1079,1080]. Adding Fe in the presence of Cu induces the formation of the Al~7~Cu~2~Fe phase (denoted as β or ω) in the form of elongated, brittle platelets [1081] during solidification. In addition, the presence of Fe-containing intermetallics may reduce the al-loy's corrosion resistance [943].

The effect of adding Si has mostly been investigated in tandem with the addition of Fe, because the two are usually added or removed together when the alloy's purity is modified. However, Si has a more complex effect: it can be associated with some of the hardening precipitates, such as the S phase in the 2024 alloy [1082], or change the precipitation sequence with the appearance of the Q phase and its precursors, as well as induce the formation of complex multi-component intermetallic phases [1083].

Other impurities may enter the alloys if mixed with other families by recycling. Adding Mn is a useful way to modify the nature of intermetallic particles, favoring phases such as Al~6~(Fe,Mn) or Al~3~(Fe,Mn) and changing the β-Al~5~FeSi phase into the less detrimental α phase (in association with Si) [1084]. When coupled with Zr, Mn can be a very useful addition to an Al-Cu binary alloy to stabilize θ' precipitates through interfacial segregation, providing exceptional thermal stability and high-temperature strength [719,1085]. Adding Ni may also change the nature and morphology of Fe-rich intermetallic particles [249].

If mixed with 7xxx scrap, 2xxx alloys may end up with some Zn content. This addition was shown to be favorable, strength-wise, for

<figure><img src="image_71.png" /><figcaption>**Fig. 69.** Range of elements in different heat exchanger alloys. Reprinted with permission from [72].</figcaption></figure>
<figure><img src="image_72.png" /><figcaption>Fig. 70. Compositions of common IMCs in dilute Al-Fe-Si systems. Reprinted with permission from [1048].</figcaption></figure>

the Al-Cu-Mg-based 2024 alloy [1086] and for Al-Cu-Li-Mg-based alloys [1087,1088], due in both cases to the incorporation of Zn into the strengthening precipitates. The affinity between Li-containing alloys and 7xxx alloys was confirmed by the discovery in 7xxx alloys of the Y phase [1089], which shows a striking similarity to the highly beneficial T1 phase in Al-Cu-Li [801,1090].

Other minor additions have been used to improve the properties of 2xxx alloys, in particular Sc, which in conjunction with Mg was shown to improve the elevated temperature strength of the 2219 alloy [1091,1092]. However, these are unlikely to be added through recycling.

### 7.3.3. Mitigating the role of impurities by processing

The detrimental effect of Fe-rich intermetallics on Al-Cu alloys is (besides their brittle nature) largely due to their size, which often exceeds 10 µm in some directions. Some of the detrimental effects of these phases can be mitigated by specific processing techniques.

When processing in the liquid phase, squeeze casting under pressure was investigated with Fe amounts of up to 1.5 wt% [1084,1093]. It was shown that the applied pressure promotes the formation of less detrimental phases, improving ductility in comparison with conventional casting routes. However, the degradation of fracture properties is still clear, at above 0.5 wt%.

Another processing approach is to cast under in conjunction with ultrasonic treatment. This method has been shown to promote higher dispersion of intermetallic particles and therefore lead to higher ductility in the final product [1094].

Fine-scale microstructures can be obtained by additive manufacturing, although Al-Cu alloys are prone to hot cracking and thus difficult to process in this way [1095]. However, Al-Cu and Al-Cu-Mg alloys have been processed successfully by laser power bed fusion, with a microstructure comprising fine-scale intermetallics [1096–1098].

The other possibility for reducing the size and improving the dispersion of intermetallics is to deploy plastic deformation. Very high-strain processes such as high-pressure torsion or equal angle channel processing are the most efficient, but are not very practical. Among more scalable processes, multi-directional forging was shown to be an efficient way to induce the spheroidization of intermetallics in alloy 2219 [1099], and friction stir processing was shown to cause a similar effect [1100].

## 7.4. Use of scrap in 3xxx and 4xxx alloys intended for packaging, buildings and infrastructure

Wrought aluminum alloys of the 3xxx series use Mn as a major alloying element, usually in the range of between 0.5 and 2 wt% (see Table 20). They can also contain Mg, Cu and Si as minor alloying ingredients. Interestingly, Mn is – among all the conventional alloying elements used in Al – the most sustainable in terms of emission and energy footprint. Additional Mg doping (0.8–1.3 wt%) increases solid solution strength and strain hardening. 3xxx alloys are usually not heat-treatable and are strengthened by cold working. Al-Mn alloys have moderate mechanical strength, weldability, very high formability and excellent corrosion resistance. They are used predominantly in products such as thin sheets for beverage cans and similar rigid containers, foils, roof sheets, cooking utensils and radiators. The fine dispersoids in 3xxx alloys tend to decorate grain boundaries and stabilize the alloys against grain coarsening during annealing, providing Hall-Petch strengthening and maintaining high formability. Mn also improves castability and reduces shrinkage during solidification [1101].

Recycling of 3xxx alloy materials is an important topic due to their sometimes very short lifetimes, particularly when serving in packaging. However, 3xxx alloys used in buildings (such as for roofs) have very long lifetimes and often enter the recycling stream only after several decades.

In the context of packaging it must be considered that cans consist of 2 different alloys: the lid is made of a slightly different alloy than the aluminum used for the base and sides. The body of the container is generally made of 3004 or 3104 alloys, which have similar chemical compositions (about 1 wt% Mn, 1 wt% Mg, up to 0.25 wt.%Cu). Some variants can even tolerate quite high Fe content (up to
0.7–0.8%). The lids consist of the 5182 alloy, with about 4-5 wt% Mg and only 0.20–0.50% Mn. The third alloy, usually 5042 with less Mg but similar Mn content, is used for the opening bank, and makes the smallest contribution to the total bank weight. Because these 3 alloy groups are chemically quite compatible, the whole product can essentially be remelted. This also applies to the minor alloy ingredients, i.e. the Cu, Si, Fe and Mn values in the 5182 alloy are in principle all below those of the 3004 alloy. This means that the joint remelting of the two introduces no impurity incompatibilities. The only exception here is Mg, which is considerably higher in the 5182 alloy than in the 3004 alloy. However, owing to its higher oxidation tendency compared to the base metal, Mg is substantially reduced in the remelting process and also during fluxing and purification [1102].

Recent developments in machine-learning-enhanced LIBS-based scrap sorting are currently improving alloy separation even among these rather closely related materials (i.e. 3xxx and 5xxx alloys) [59,769,781,1103,1104].

Another aspect of high relevance for recycling of packing alloys is the use of paint. It was observed that Ti in particular can intrude as a paint debris contaminant into melts produced from secondary sources [119,123,127]. This suggests that such contamination should be prevented by removing any paint by adequately pre-treating the scrap and properly adjusting the melting temperature.

Wrought Al-Si silicon alloys of the 4xxx alloys series form only a small group in the overall alloy market. They include both heat-treatable and non-heat-treatable materials, although most of them (except alloy 4032 with 1 % Mg and alloy 4145 with 4% of copper) are non-heat-treatable. Besides their major alloying element, Si, they may contain Cu, Mg, Ni and Be as minor alloying ingredients. Si is blended in these materials at up to 13.5% (alloy 4032) to reduce the liquidus temperature and increase fluidity, because these alloys are mainly used in filler materials such as welding wires or brazing rods. Some are also used in buildings for architectural reasons owing to their characteristic gray color in the anodized state.

## 7.5. Use of scrap in 5xxx alloys intended for automotive, buildings and infrastructure

The classic alloys of the 5xxx series generally have narrow Si tolerance limits and are therefore rated as “tough” with regard to recycling. Compositions and uses of typical 5xxx wrought alloys are shown in Table 21. Even at low concentrations, Si can react with the Mg present in the melt in 5xxx alloys to form the primary Mg₂Si phase, which is usually undesirable [78]. This is illustrated in Fig. 71, where the Mg content is comparable to that of the common alloy 5182. It can be seen that at elevated Si content, formation of Mg₂Si occurs, and there is no option to dissolve Si into the matrix by homogenization or solution annealing. Only at below 0.1% Si does a gap begin to appear between the Mg₂Si solvus and the solidus temperature of the alloy, providing an opportunity to dissolve Mg₂Si. Even in the 5xxx series of wrought alloys made mainly from primary aluminum (e.g. 5182), primary Mg₂Si can be found which has survived homogenization treatments [1105]. Mg₂Si was seen to form fractured stringer-like particle structures in rolled sheets of 5754 [1106]. Moreover, it is known from observations in other wrought alloys that primary Mg₂Si particles fracture at even lower strains than Fe-containing IMCs [877], highlighting the detrimental effect of primary Mg₂Si.

Although there are some 5xxx alloys where Cu is used for precipitation hardening [1051,1052,1109], there are usually also strict limits on Cu in the important 5xxx alloys (see Table 21). Besides the effect of hardening due to low Cu addition, Cu together with added Zn can also reduce the volume fraction of β-Al₃Mg₂. This introduces a less anodic phase and thereby mitigates intergranular cracking, which can be problematic in 5xxx alloys at elevated temperature [1110]. This illustrates the fact that sometimes elements that are subject to tight restrictions can in some cases be used to generate beneficial properties. Not only Cu exhibits such effects on certain types of corrosion in 5xxx series alloys; many other minor elements have also been studied in this context [1111,1112]. Interestingly, Ce has recently been shown to be beneficial [1113], and may appear in some automotive scrap in the future [1022].

In addition to Si and Cu, Fe is set to relatively low values in specific alloys. This has to do with the common issue of reduced formability at increased Fe content in aluminum alloys. For example, a twin-roll-casted (TRC) Al-Mg alloy with a high impurity content (0.72 wt% Si, 0.58 wt% Fe) was investigated and exhibited a fine distribution of Fe-containing intermetallics and Mg₂Si at higher cooling rates (≈ 10³ K/s), and hence improved tensile and fracture properties. Although the secondary phase particles in the as-cast TRC strip were shown to be fine and discrete by 2D analysis, 3D analysis showed that these particles were arranged in intact networks [221]. Hence, 2D evaluation of particles can lead to misleading assumptions and should be carefully scrutinized. However, after homogenization treatment, these networks were broken up into clusters and Fe-bearing intermetallics and Mg₂Si spheroidized [221]. These aspects will be discussed in detail later in this section in the context of 5xxx series alloys.

5xxx series alloys are typically used in structural plates or sheets and in extrusions for transportation and engineering, where their high ductility (even at cryogenic temperatures), good weldability and high corrosion resistance is important. Typical products are marine applications, tanks and pressure vessels, and ballistic armour plates [970,1018]. In transport applications where high formability is required, especially in the automotive sector, 5xxx series alloys are the exclusive choice among non-heat-treatable wrought alloys because of their superior strain hardening behavior [1018,1114]. However, a drawback of the 5xxx series alloys is the

<table><caption>**Table 20**

Chemical compositions of some wrought 3xxx aluminum alloys (wt.%).</caption><thead><tr><th>Alloy designation</th><th>Cu</th><th>Mn</th><th>Mg</th></tr></thead><tbody><tr><td>3003</td><td>0.05–0.20</td><td>1.0–1.5</td><td>—</td></tr><tr><td>3004</td><td>0.25 max</td><td>1.0–1.5</td><td>0.8–1.3</td></tr><tr><td>3005</td><td>0.30 max</td><td>1.0–1.5</td><td>0.2–0.6</td></tr><tr><td>3105</td><td>0.30 max</td><td>0.3–0.8</td><td>0.2–0.8</td></tr></tbody></table>
appearance of Lüders elongation, dynamic strain aging, and deteriorating surface quality via formation of strain localization upon forming [86,87,1115]. In addition to processing-related measures, it appears that less-pure alloys may improve this behavior. It has been shown that adding Cu and Zn can reduce Lüders elongation and dynamic strain aging [87]. Interestingly, large IMCs can also help to suppress Lüders elongation [86]. Ebenberger et al. [86] showed that the differences in thermal expansion behavior between IMCs and the aluminum matrix can be utilized to introduce mobile dislocations via quenching that are not blocked by dissolved Mg atoms (illustrated in Fig. 72), thus eliminating detrimental Lüders elongation. In [86] IMCs of type Al<sub>13</sub>(Mn,Fe)<sub>6</sub> were studied in alloy 5182, but this basic mechanism is applicable to other IMCs with different thermal expansion coefficients than Al. This can be seen as a potentially beneficial side-effect of introducing tramp elements through recycling, which often generate coarse IMCs.

In most 5xxx series alloys used today for forming and automotive applications, impurities are still typically kept to a minimum [78]. Nevertheless, an automotive alloy of the 5xxx series with a high content of recycled material already exists. Novelis UK and Jaguar Land Rover have developed a 5754 alloy with up to 75% recycled scrap, called RC5754. In the context of such alloys, twin-roll casting has also been investigated as a method of achieving high cooling rates, which favor the formation of less-harmful Fe-containing IMCs [224]. Recently Al-Helal et al. [202] showed that even 100% scrap can be used. However, in this study Taint Tabor scrap (a special clean wrought alloy scrap from the UK which includes a large fraction of production scrap [202]) was used. This scrap was shredded further, thoroughly re-sorted, and cast via melt conditioned direct chill casting (MC-DC) [1116]. 5754 billets produced from 100% Taint Tabor scrap were then extruded, cold rolled to sheets and annealed [202]. In the best-case scenario, the industrial limits for 5754 were reached directly. Here it can be noted that the MC technique can also be used in twin-roll casting [224]. Recently, Das et al. [1117] concluded that to make a 5754 alloy recycling-friendly the processing conditions need to be adjusted, because additional IMCs and dispersoids affect microstructure formation. Modified heat treatments should be applied to improve properties such as self-piercing riveting performance, where a high level of formability through a combination of optimised grain size and balanced texture is required.

Que et al. [1118] made the important point that “understanding the formation of Fe-rich IMCs during solidification of Al-alloys is essential for developing effective approaches to mitigate their harmful effect on mechanical properties.” Here it is important to recognise that appropriate grain refinement was also shown to be significant in obtaining a good distribution of primary IMCs when casting recycled 5xxx series wrought alloys [175]. Although the literature mainly addresses Fe-containing IMCs, it must be considered that in 5xxx series alloys with increased impurity content, one is dealing with an array of particles, not just Fe-containing IMCs (e.g. Mg<sub>2</sub>Si, the S phase (Al<sub>2</sub>CuMg), and the Q phase (AlMgSiCu)) [175,1044].

<table><caption>**Table 21**

Composition [970,1018] (if no range is given, the values are upper limits) and typical application of 5xxx alloys.</caption><thead><tr><th>Designation</th><th>Si</th><th>Fe</th><th>Cu</th><th>Mn</th><th>Mg</th><th>Zn</th><th>Cr</th><th>Ti</th><th>Typical application</th></tr></thead><tbody><tr><td>5182</td><td>0.20</td><td>0.35</td><td>0.15</td><td>0.2-0.5</td><td>4.0-5.0</td><td>0.25</td><td>0.15</td><td>0.1</td><td>Packaging products, automotive body panels and reinforcements</td></tr><tr><td>5083</td><td>0.40</td><td>0.40</td><td>0.10</td><td>0.4-1.0</td><td>4.0-4.9</td><td>0.25</td><td>0.05-0.25</td><td>0.15</td><td>Transportation equipment, marine parts, pressure vessels</td></tr><tr><td>5754</td><td>0.40</td><td>0.40</td><td>0.10</td><td>0.50</td><td>2.6-3.6</td><td>0.30</td><td>0.20</td><td>0.15</td><td>Vehicle bodies, tread plates, shipbuilding, rivets</td></tr><tr><td>5251</td><td>0.40</td><td>0.50</td><td>0.15</td><td>0.1-0.5</td><td>1.7-2.4</td><td>0.15</td><td>0.15</td><td>0.15</td><td>Panelling and pressings, vehicle panels, furniture</td></tr><tr><td>5005</td><td>0.30</td><td>0.70</td><td>0.20</td><td>0.20</td><td>0.5-1.1</td><td>0.25</td><td>0.1</td><td>-</td><td>Architectural facades, furniture, high-strength foils, HVAC equipment</td></tr><tr><td>5052</td><td>0.25</td><td>0.40</td><td>0.10</td><td>0.10</td><td>2.2-2.8</td><td>0.10</td><td>0.15-0.35</td><td>-</td><td>Architectural facades, tread plates, fence systems, furniture</td></tr><tr><td>5086</td><td>0.40</td><td>0.50</td><td>0.10</td><td>0.2-0.7</td><td>3.5-4.5</td><td>0.25</td><td>0.05-0.25</td><td>0.15</td><td>Pressure vessels, transportation equipment</td></tr></tbody></table>

<figure><img src="image_73.png" /><figcaption>**Fig. 71.** Evolution of Mg<sub>2</sub>Si and the liquid phase in an Al-4.5 %Mg-x%Si system as a function of temperature when increasing Si content from 0.1% to 0.5% Si in thermodynamic equilibrium (calculated with FactSage 8.0 [1107] software using the FACT light-metal alloy database FTlite 2020 [1108]).</figcaption></figure>
For all these IMCs, solidification rate plays an important role. High-speed twin roll casting was shown to be beneficial for an Fe-
containing 5xxx series alloy [222]. Liu et al. [1119] also showed that cooling rate plays a critical role and that the Al₆(Fe,Mn) IMCs can
be given fine Chinese-script morphology by near-rapid cooling (e.g. continuous strip casting). The same group also investigated the
effect of increasing Fe and Si content in 5083 as a function of the cooling rate during solidification [1120,1121]. An increase in Fe
content leads to an increase in the amount and size of the Al₆(Fe,Mn) phase. An increase in Si content generates an increase in the
Mg₂Si phase, which forms a network structure. Increasing the cooling rate at high Fe and Si content not only significantly refines the
IMCs, but also changes their type to a quaternary Al₁₅(Fe,Mn)₃Si phase with a fish-bone or Chinese-script structure, which is claimed to
be beneficial for the strip casting of recycling-based 5xxx series wrought alloys.

The currently high usage of Si in casting alloys, together with the usage of other impurities, will make the production of con-
ventional 5xxx series alloys via ELV recycling difficult. In particular, a mix with 6xxx wrought alloys from ABS is problematic here
[1122]. Hence, at present 5xxx series alloys are only easily converted into specific 5xxx series alloys, or are fed into some 6xxx series
alloys in ABS recycling by removing Mg through oxidation [1122].

Que et al. demonstrated one way to address this problem. They intensively studied an Al-5Mg-2Si-0.7Mn-1.1Fe alloy [1118,1123-1127], which is an interesting model system for highly contaminated 5xxx-series alloys from ELVs. Originally they studied the alloys from the perspective of cast alloys. However, [1118] investigated the formation of Fe-containing IMCs during solidification in twin-roll casting of strips for automotive applications made of this alloy. In their model alloys the authors obviously considered high Si content and also addressed high Fe contamination. Unfortunately, no reports on the mechanical properties of the sheets made from this alloy seem to be available so far.

Although high-pressure die casting was used for processing in the work of Zhu et al. [1000] and Liu et al. [1128], the base alloy
composition resembles a typical wrought 5xxx series alloy. From Fig. 73 it can be estimated that there is potential for significantly
increasing the Fe content tolerated in 5xxx series alloys that do not require superior ductility under rapid solidification conditions
[1000].

Algendy et al. [1129] showed that the formation of Fe-containing IMCs depends not only on the commonly studied elements Fe, Mn and Si. The range of Mg content in the 5xxx series alloys must also be taken into account, and different IMCs can form. Que et al. [1130] also demonstrated that the morphology of the formed eutectic Al₆(Fe,Mn) changed from needle- and plate-like to Chinese script in Al-1.4Fe-0.7Mn-xMg if the Mg content is increased from 0 to 3%, even though Mg is not found in the Al₆(Fe,Mn) IMCs. They also reported that an increase in Mg content also leads to grain refinement via heterogeneous nucleation of α-Al on naturally formed MgAl₂O₄ particles. Interestingly, they observed the accumulation of Mg on the surface of Al₆(Fe,Mn) particles, with the likely formation of an Al₁₂Mg₁₇ phase a few nanometres thick. The authors concluded that growth rates along different growth directions in Al₆(Fe,Mn) particles are changed by Mg segregation in such way that growth becomes two-dimesional at low Mg and three-dimensional at high Mg content, generating the morphologies observed. These results may be of interest when using new alloying elements that indirectly influence the morphology of Fe-containing IMCs.

Of course, the formation of Fe-containing IMCs in 5xxx wrought alloys also depends on the solidification conditions, and hence, in
conventional DC-casting, on the local position in the cast ingot. Li and Arnberg [1131] reported in the context of an Fe-containing 5182
alloy that as a function of the distance from the skin of the ingot, the type of IMC changed from Al<sub>m</sub>(Fe,Mn) with skeletal morphology to
Al<sub>3</sub>(Fe, Mn) with plate-like or rod-like morphology. They explained the tendency to form Al<sub>m</sub>(Fe,Mn) in the eutectic reaction at higher
cooling rates by the higher growth temperature at higher growth velocity compared to Al<sub>3</sub>(Fe,Mn). In a further work [1132] they found
that the primary Al<sub>m</sub>(Fe,Mn) can be transformed, by homogenization via a eutectoid transformation, to a lamellar mixture of Al<sub>3</sub>(Fe,
Mn) and Al. The authors linked this with an easy break-up of primary particles during further hot rolling. Interestingly, they found no
Al<sub>6</sub>(Fe, Mn) in 5182, and explain this by the high Mg and low Mn content in the alloys. It should be noted that other elements, such as
Cr, are also important for IMC formation in some industrial 5xxx alloys [1133].

All these findings indicate that 5xxx series wrought alloys feature complex relationships between casting conditions, chemical composition, phase selection and the phase morphology of Fe-containing and other IMCs introduced by recycling. Only systematically analyzed and understood relationships can be used to develop 5xxx series alloys that are more suitable for recycling.

<figure><img src="image_74.png" /><figcaption>**Fig. 72.** Electron Channelling Contrast Image reevaling large IMCs (a) and mobile dislocations (b) in a 5182 alloy after quenching. Electron Channelling Contrast Imaging is a technique which reveals lattice defects such as dislocations directly in a scanning electron microscope. Reprinted with permission from [86].</figcaption></figure>
## 7.6. Use of scrap in 6xxx alloys intended for automotive and mobile communication applications

The 6xxx series wrought alloys are not extremely sensitive to impurities from recycling, but production from scrap can still be a challenge (see Fig. 84, section 7.1). Bearing this in mind, it is important to appreciate that 6xxx wrought alloys are frequently used [1134] and are expected to comprise the major part of recycling scrap in the future – even though today only 20% of 6xxx series scrap is recycled into 6xxx wrought alloys [3,252]. Table 22 illustrates the compositions and common uses of typical 6xxx alloys. The main restrictions in some alloys concern Fe and Cu, or Mn and Zn, which are all potential cross-contaminants from the recycling of aluminum scrap. Other elements (e.g. Cr) are also subject to tight limits in certain alloys, but are only minor alloying elements in specific alloys which usually represent minimal contamination in typical scrap streams. The fields of application for 6xxx alloys are wide and include the construction, automotive, marine and aircraft industries in rolled, extruded, drawn and forged form [1018]. This is due to their interesting combination of properties such as good formability (especially for extrusion), high corrosion resistance, excellent surface properties and weldability, combined with relatively low price [1054,1135]. The wide range of the 6xxx series alloys is also shown in Fig. 65 in section 7.1, where it is seen that only the extreme strength and elongation limits cannot be reached in the field of wrought aluminum alloys.

A 6xxx alloy that is particularly well suited for recycling is 6061 [51,171,269]. This is mainly due to the fact that this alloy can absorb a considerable amount of Mg, Si, Fe and Cu (see Table 22). It is one of the most commonly used 6xxx alloys. Some other alloys of the 6xxx series are much more limited in the concentration of impurities they tolerate. This is especially true of 6xxx alloys used as automotive body sheets (ABS), such as 6022.

However, from a plot of dilution-quality-losses for 6061 (Fig. 74) derived from a Monte Carlo-based decision support tool in [171] it becomes clear that even the production of the impurity-tolerant alloy 6061 can require significant dilution with primary aluminum if it is produced by recycling the material from scrap. This is even the case if scrap of superior quality from the wrought aluminum products of vehicles (SC2) is used. If scrap of other qualities is deployed, e.g. that from engines and vehicle transmissions (SC1) or used beverage cans (UBC) (SC3), the demand for primary aluminum is enormous. Mixing these scrap classes helps, but does not solve the problem. This example shows clearly that the recycling of used wrought alloys into new 6xxx series wrought alloys is currently a challenge.

Dilution losses can be only reduced at the moment by using compositionally similar grades (i.e. the same alloys) [171,1122]. This type of recycling is called “tolling” and is applicable in closed-loop systems directly during production, such as stamping scrap. The

<figure><img src="image_75.png" /><figcaption>Fig. 73. Yield strength and elongation of Al-5Mg-xFe and Al-5Mg-0.6MnxFe alloys. Reprinted with permission from [1000].</figcaption></figure>
Ford Motor Company used it for the F-series pickup truck, where up to 33% of primary aluminum can be saved with the strategy [78]. However, for end-of-life products and complex systems the approach does not seem very accessible at the moment.

The complex compositional challenge arises not only from the mixture of different alloy series from wrought and cast alloys and other metals typical of ELV recycling. Arowosola et al. [72] noted that over the last few decades several different 6xxx series alloys (6005, 6009, 6010, 6111, 6014, 6016, 6022 and 6451) have been registered for automotive skin applications alone. This makes for more complexity even within a single series of wrought alloys used for a single purpose, without taking into account that in reality the compositional space of a single standard cannot be fully exploited for the industrial production of high-performance alloys for many products. This is particularly true of automotive applications, where alloy composition requirements are generally very tough.

The following section examines the main metallurgical effects of critical elements introduced into 6xxx series alloys by recycling. Although many other elements can exceed standard limits and such considerations depend on the specific case (e.g. Mg from 5xxx series alloys can also be a problem in recycling), in this section we concentrate primarily on Fe and to a lesser extent on Cu, as these elements are the most widely discussed in the literature.

The Fe content in 6xxx alloys is limited due to the formation of primary Fe-containing IMCs [78], which can cause crack initiation, especially when present as brittle needles or plates [78,252]. The negative effects of these Fe-containing IMCs can be reduced during wrought alloy processing, i.e. via an appropriate homogenization practice and/or plastic deformation to break large Fe-containing IMCs into smaller ones [78]. However, it is not just Fe-containing IMCs which need to be considered in the context of fracture initiation in 6xxx series alloys. In studying a 6061 alloy, Petit et al. [877] showed that large Mg₂Si particles can fracture at low plastic strains. In addition, reduced ductility and bendability have been associated with increased Fe content in 6xxx series alloys [322] (see also Section 3.7). However, they are not always linearly connected with this Fe content. Lu et al. [1136] indicated that hemming and bending properties in 6xxx series alloys may result from a complex interdependence between Fe content and processing conditions such as the degree of hot rolling and quenching parameters, and that there can be sweet spots which do not require the lowest Fe content. The surface finish of 6xxx alloys can also be affected by Fe-containing IMCs [1137,1038].

Besides these Fe effects, the formation of primary Fe-containing IMCs also removes Si from solid solution [78]. This can significantly affect artificial aging kinetics, as shown in Fig. 75 [735], and needs to be considered if high Fe content is introduced via recycling. Interestingly, the effect of reduced Si super-saturation on aging kinetics is much stronger when novel aging strategies based on direct artificial aging are deployed [735].

Modification via Fe in the dissolved Si concentration in the aluminum matrix also depends on which IMC phase is actually formed [78]. In 6xxx series alloys, as-cast IMCs often exhibit an Si-rich composition around β-Al₅FeSi, which can be transformed into α-Al₁₂(Fe,Mn)₃Si by homogenization [887,1038,1138,1139]. This transition can release Si into the aluminum matrix [78]. Mn plays an important role in this transition [78]. Fig. 76 illustrates the effects that can be achieved through heat treatment and compositional measures. The metastable β phase in as-cast condition shown can be transformed into α by suitable homogenization.

Besides its effect on the Si supersaturation, the morphology of the Fe-containing IMCs formed is crucial for the properties of these alloys. For the often-discussed impurity-tolerant 6061 alloy, the most important IMCs are α-Al₈Fe₂Si, α-Al₁₅(FeMn)₃Si₂ and β-Al₉Fe₂Si₂ [269]. The α phases can show Chinese-script or polyhedral morphology, while the β phase comprises unfavorable plates or needles [269]. As discussed above, the detrimental β phase can be transformed by long-term, high-temperature treatments which generate spheroidization [78]. According to Wang et al. [1140], this was shown to significantly improve elongation and fatigue strength in an extruded 6060 alloy. Moreover, after exposing the material to a slightly corrosive environment its surface appearance can be improved,

<table><caption>**Table 22**

Typical composition and uses of 6xxx wrought alloys [1018,1054].</caption><thead><tr><th>Designation</th><th>Si</th><th>Fe</th><th>Cu</th><th>Mn</th><th>Mg</th><th>Zn</th><th>Cr</th><th>Ti</th><th>Typical application</th></tr></thead><tbody><tr><td>6016</td><td>1.0-1.5</td><td>0.50</td><td>0.20</td><td>0.20</td><td>0.25-0.6</td><td>0.20</td><td>0.10</td><td>0.15</td><td>Automobile body sheets widely used in Europe</td></tr><tr><td>6022</td><td>0.8-1.5</td><td>0.05-0.2</td><td>0.01-0.11</td><td>0.02-0.1</td><td>0.45-0.7</td><td>0.25</td><td>0.1</td><td>0.15</td><td>Automobile body sheets</td></tr><tr><td>6014</td><td>0.30-0.6</td><td>0.35</td><td>0.25</td><td>0.05-0.20</td><td>0.40-0.8</td><td>0.10</td><td>0.20</td><td>0.10</td><td>+0.05-0.20 Vanadium</td></tr><tr><td>6005</td><td>0.60-0.9</td><td>0.35</td><td>0.10</td><td>0.10</td><td>0.40-0.6</td><td>0.10</td><td></td><td>0.10</td><td>automobile body sheets, profiles</td></tr><tr><td>6111</td><td>0.60-1.1</td><td>0.40</td><td>0.50-0.9</td><td>0.20</td><td>0.10-0.45</td><td>0.15</td><td>0.10</td><td>0.10</td><td>General-purpose extrusions, rail vehicles, construction</td></tr><tr><td>6060</td><td>0.30-0.60</td><td>0.10-0.3</td><td>0.10</td><td>0.10</td><td>0.35-0.6</td><td>0.15</td><td>0.05</td><td>0.10</td><td>Automobile body sheets</td></tr><tr><td>6063</td><td>0.20-0.6</td><td>0.35</td><td>0.10</td><td>0.10</td><td>0.45-0.9</td><td>0.10</td><td>0.10</td><td>0.10</td><td>General-purpose extrusions, rail vehicles, construction, forgings</td></tr><tr><td>6061</td><td>0.40-0.80</td><td>0.70</td><td>0.15-0.40</td><td>0.15</td><td>0.8-1.2</td><td>0.25</td><td>0.04-0.35</td><td>0.15</td><td>Casing for electronic devices, architectural extrusions, pipes, door and window frames, heat exchanges, rail vehicles, bumpers</td></tr><tr><td>6082</td><td>0.70-1.3</td><td>0.50</td><td>0.10</td><td>0.40-1.0</td><td>0.6-1.0</td><td>0.20</td><td>0.25</td><td>0.10</td><td>Higher-strength applications, welded structures, rail vehicles, construction, bumpers</td></tr><tr><td>6013</td><td>0.40</td><td>0.50</td><td>0.10</td><td>0.2-0.7</td><td>3.5-4.5</td><td>0.25</td><td>0.05-0.25</td><td>0.15</td><td>Higher-strength applications widely used in Europe, extrusions</td></tr><tr><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td>Plates, chassis components, forgings, rail vehicles, construction, bumpers</td></tr><tr><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td>Aircraft sheets, fuselages</td></tr></tbody></table>
<figure><img src="image_76.png" /><figcaption>**Fig. 74.** Dilution-quality-losses plot for the production of 6061 from different types of scrap: Al engines and transmissions (SC1), wrought products from vehicles (SC2), UBC (SC3) and a mix of these. All model solutions are shown as scatter diagram for the addition of primary Al (dilution losses) and required alloying elements. The best case for recycling is at the lowest dilution losses and where there are small amounts of additional alloying elements. Reprinted with permission from [171].</figcaption></figure>

because micro surface defects due to the debonding of β-particles from the Al matrix can be averted by a suitable heat treatment [1140]. This is illustrated in Fig. 77.

The adverse impact of the α phase can be mitigated in 6xxx alloys with increased Mn content (a ratio of Mn/Fe > 0.5 is often suggested) [78]. This promotes α-type IMCs (see Fig. 76), which do not cause ductility to deteriorate as much. Hence, promotion of the α phase by adding Mn makes it possible to accept higher Fe content, while maintaining good tensile elongation [1141]. Adding Cr has a comparable effect [78]. However, these compositional measures also have their limitations: too much added Mn increases porosity and sludge formation, decreases machinability [252,339,1142,1143] and reduces strength and ductility [1144].

The type of IMC formed during casting also depends on Fe content. Sweet et al. [886] showed that in 6060 an increase from 0.1 wt% to 0.5 wt% Fe shifted the cast microstructure from β-Al~5~FeSi to a Chinese-script cubic α~c~-Al(Fe,Mn)Si phase (see Fig. 54). This observation may support the above-mentioned non-linear effect of Fe content on bending and hemming properties. In general, there are several Fe-containing IMCs (only in the lean 6060 alloys described in [886] have four different phases been found in as-cast conditions) which may form in 6xxx series alloys. The phase depends strongly on composition (involving many transition metals or even Mg for the π-Al~8~FeMg~3~Si~6~ phase) and the highly complex formation progression upon solidification, which includes a series of peritectic reactions [886,1145].

It seems very significant that the array of phases in 6xxx alloys depends on the solidification conditions. As already shown, the β phase is of a metastable nature, and its formation depends on the growth rate. Sha et al. [1050] demonstrated in 6xxx model alloys made by Bridgman furnace solidification that high growth rates above 120 mm/min strongly favor formation of the beneficial α phase. Consequently, various novel casting techniques are being utilized to improve the morphology and distribution of Fe-containing IMCs. Al-Helal [1146] et al. used high shear melt conditioning to create less-harmful Fe-containing IMCs in twin roll cast recycled 6111 strips. In this context, Wang et al. [1147] studied an alloy with 0.8 wt% Mg, 0.8 wt% Si and 0.7 wt% Fe which they produced by direct chill casting, and reported a refinement of Fe-containing IMCs, Al-grains and an increased fraction of small α-dispersoids in billets cast at high speeds. There are complex links between the microstructure, morphology, type and interconnectivity of the Fe-containing IMCs in the cast state which also influences the formation of other IMCs during processing. This includes Mg~2~Si, as this phase tends to nucleate at β-Al~5~FeSi [1044]. Inoculation and its interdependence with Fe-containing IMCs upon solidification has also been investigated in the literature [1046,1047].

It should be noted that high levels of Fe content also influence the conventional casting process of 6xxx series alloys. In the context of direct chill casting of a 6xxx series alloy, Nagaumi et al. [1148] reported that increased Fe content can significantly degrade hot cracking behavior.

Sweet et al. [1149] studied hot tear susceptibility in DC cast 6060 aluminum alloys with various amounts of Fe. They observed that materials with an Fe content ranging from 0.02 to 0.5 wt% had a substantial influence on load response during solidification and on the
<figure><img src="image_77.png" /><figcaption>**Fig. 75.** Dependence of age-hardening kinetics on the super-saturation of Mg plus Si (MSSS<sub>570-175</sub> of MgSi) expressed by a rate constant k according to the model used in [735] within the standard compositional limits of 6061. The results for direct artificial aging and artificial aging with storage at RT are shown. The effect of Fe (and in this case Ti) which form IMCs and trap Si is clearly shown ("mean" denotes a similar nominal Si and Mg content in the 6061 alloy used). Reprinted with permission from [735].</figcaption></figure>

<figure><img src="image_78.png" /><figcaption>**Fig. 76.** Isothermal section at 540 °C in an alloy with 1 wt% Si and 0.5 wt% Mg indicating the fields of existing IMCs in dependence on the Mn and Fe content. The effect of Scheil-cooling and the resulting extended metastable β phase field in as-cast condition are shown. This can be brought into equilibrium via homogenization (calculated with Pandat software using database PanAl2019 [982]).</figcaption></figure>

resulting hot tear susceptibility. They interpreted their observations in terms of the influence of Fe content on the frequency, size and morphology of β-Al<sub>5</sub>FeSi and α-AlFeSi IMCs and their influence on the coherency and coalescence of the microstructure.

In the same context, Sree et al. [1150] recently introduced an in-situ microstructural control method for changing the dispersion and distribution of second-phase particles via a mechanically melt conditioned direct chill (MC-DC) casting process, operated only 2–7 °C above liquidus. The approach makes use of a rotor–stator high-shear device. These authors successfully applied the method to an 6082 alloy containing 0.27 wt% Fe. The results show that the approach improves particle dispersion and thus helps to reduce or even suppress the need for subsequent expensive, time-consuming homogenization operations. The method thus clearly increases flexibility in the use of higher scrap fractions in such alloys, where second-phase dispersion is otherwise a challenge.

The second element to be considered is Cu, which shows two faces in the 6xxx series. It is strongly restricted in many alloys (often to below 0.1 wt%), but there are also 6xxx wrought alloys that contain >1 wt% Cu. This becomes important when we consider the 6xxx alloys used together in similar products, e.g. cars. A good example of misfit affecting recycling is seen in the frequently used ABS alloys 6022 and 6111, where compositional differences in Cu are extreme and complicate recycling [1122]. Cu acts in these alloys as follows. High Cu content (e.g. in 6111) is used to generate additional strengthening because it introduces the Q phase and its precursors [808]. The rationale for tight limits on Cu in 6022 (and many other primary alloys), however, is a bit less clear. One reason may be the deteriorative influence that Cu may have on various types of corrosion in 6xxx series alloys, even though this also depends in very complex ways on the composition and processing conditions affecting the different types of corrosion [1151,1152].

The Cu content of Al-alloys used in enclosures for electronic devices (such as 6063) is controlled carefully (<0.1 wt%), because it can have an impact on cosmetic appearance by affecting the color of the enclosure after anodizing. Fe can influence the appearance of Al-alloys and generate grayish colors [1153], possibly due to the disappearance of Fe-rich IMCs after anodizing, which leaves etching pits which reduce surface gloss [1154]. The formation of "foggy" oxides in Fe-containing alloys may also contribute to their unattractive appearance. Gray coloring can also arise from the presence of Si in the alloy. Anodization of pure aluminum generates the brightest anodic oxides compared to Al-alloys [1155]. The challenge of reproducing a defined compositional range to achieve
consistent cosmetic appearance increases with the degree of recycling and is related to the source and amount of recycling material. Cosmetic aspects are often linked to corrosion behavior. The effects of scrap-related tramp elements on corrosion are discussed in more detail in section 5.6.

In conclusion, we will mention some new recycling-based concepts for 6xxx alloys addressed in recent literature. In [252] Kotadia et al. explored extending the current composition limits of 6xxx alloys, particularly with regard to Si, Fe and Mg content, but they also considered further tramp elements. As a first step, they showed that ultrasonication upon casting can significantly refine grain structure and effectively modify Fe-containing IMCs in a beneficial way. Other new technologies, such as spray forming, are also being investigated to produce starting material for forming operations that can tolerate higher Fe content. Pereira et al. [269] showed in the context of wrought alloy 6061 that IMCs with detrimental morphologies can be avoided even at 1.4 wt% of Fe. The formation of the primary α-Al₁₅(Fe,Mn)₃Si₂ phase with polyhedral morphology and homogeneous distribution can generate significant ductility and enable the hot extrusion of such 6xxx material enriched with Fe [269].

To summarize, a major feature of current literature on the development of recyclable 6xxx series alloys is acceptance of higher levels of Fe contamination. This applies to research on novel casting technology, solidification and inoculation, heat treatments and alloy solidification techniques other than conventional casting.

## 7.7. Use of scrap in 7xxx alloys intended for high-strength applications

Al-alloys in the 7xxx series are heat-treatable, and those in the quaternary system Al-Zn-Mg-Cu in particular have the potential to reach very high strength levels (tensile strength up to 850 MPa [1156], and up to 1 GPa with ultrafine grain structures [629]) in comparison with the other Al-alloy classes. These high strength levels are primarily achieved by the precipitation of precursors of the η phase (MgZn₂) during heat treatment. Further increases in strength can arise by adding significant amounts of Cu, as in most modern aerospace alloys. This can alter precipitates into Mg(Zn,Cu)₂ [102] compounds.

### 7.7.1. The role of Fe and Si impurities in 7xxx series alloys

Historically, the widest application field of the 7xxx alloy series has been in structural components for the aircraft industry. This is due to their high strength-to-weight ratio and relatively low cost [913,1157]. Fracture-critical design concepts led to the development of special high-toughness alloys, such as 7150, 7175 and 7475 [1157]. The latter involves modification of the high-strength 7075 alloy to improve ductility. By limiting the total amount of Fe and Si content to max. 0.22 wt% in 7475 (compared to a limit of max. 0.9 wt% in 7075), both fracture toughness properties and strength were improved [913]. Reducing impurity elements, such as Fe and Si, is a common way to improve fracture toughness properties, not only in the aircraft industry. While impurity content is limited to 0.10–0.15 wt% Fe and 0.15–0.25 wt% Si in superior, high-toughness aerospace alloys, Fe and Si levels are typically limited to max. 0.40 wt% in wrought premium automotive alloys such as 7116 and 7129 (both max. 0.30 wt% Fe), used for bumpers. The amount of Fe and Si tends to increase the more frequently an Al-alloy is recycled [29], mainly because it is picked up by scrap handling equipment [1158]. An increase in Fe in recycled Al-alloys during automotive recycling is considered unavoidable, and presumably arises from the

<figure><img src="image_79.png" /><figcaption>**Fig. 77.** Schematic illustration of α- and β-particles during processing and the impact on surface properties of 6060 extrusions. Reprinted with permission from [1140].</figcaption></figure>
unintended mixing of steel parts in these alloys due to non-prompt sorting [1159]. Because Fe is highly soluble in molten Al [1257], abrasion from the steel machining tools used in melting production processes can increase the Fe content in Al products [1258]. Fe is in fact the most common impurity element in Al-alloys. Due to its low solubility in the solid state (≈ 0.05 wt% [1257]) it can form intermetallic particles, which can impair mechanical properties [1162] and corrosion resistance [102].

The presence of impurities may be unavoidable, and is to be expected in scrap material. Therefore we will address the effects of impurities in 7xxx series alloys and similar alloys. Several studies focus on the effects on microstructure caused by higher impurity content in 7xxx series alloys. Micropillar compression tests conducted on constituent particles in a 7075 alloy showed the completely brittle nature of Al<sub>7</sub>Cu<sub>2</sub>Fe particles, which had a higher compressive strength compared to Mg<sub>2</sub>Si particles and the 7075 alloy matrix [1163]. Microstructures of 7475 and 7081 alloys after potentiodynamic polarization tests in 3.5% NaCl solution revealed corrosion pits around Al<sub>7</sub>Cu<sub>2</sub>Fe particles, indicating the cathodic behavior of these particles relative to the matrix [1164]. Al<sub>7</sub>Cu<sub>2</sub>Fe particles were shown to be insoluble during heat treatment due to the high melting point [1165–1168]; subsequent mechanical processing leads to the cracking and re-allocation of these particles [1168]. In another study, the 7075 alloy, with 0.23 wt% Fe, was successfully produced by TRC. It exhibited a homogeneous distribution of small particles (≈ 1 µm) due to the higher solidification rate, compared to the same alloy fabricated with a permanent mold casting (PMC) technique, which generated a large number of finer particles (≈ 0.1 µm), but also a few much larger particles (>5 µm) (Fig. 78). While the small particles in the TRC alloy induced a fine-grained microstructure with enhanced ductility, the fine particles in the PMC alloy induced an inhomogeneous microstructure and decreased elongation at high temperatures during hot tensile deformation. Cracks were observed inside large Fe-bearing particles (>5 µm) and also at the interface between particles and matrix, lowering fracture elongation [1169].

The susceptibility of 7xxx series alloys to hot tearing has been shown in the context of DC casting. A DC-cast 7055 alloy exhibits a wider temperature range between solidus and liquidus than pure Al and an Al-4.5 wt%Cu alloy, leading to a higher temperature range in the sump. Due to the higher temperature range, cracks initiate and propagate more easily in 7xxx series alloys than in the other alloys investigated. Hot tearing occurred in the center of the ingot due to the formation of a shrinkage cavity [1170].

Several strategies can be used to overcome hot cracking in Al-Zn-Mg-Cu alloys during processing through SLM. Al<sub>3</sub>Zr particles have been shown to be particularly effective, serving as solidification nuclei and thus helping to refine grain size, thereby reducing hot cracking susceptibility. These particles can be introduced as Zr nanoparticles added to the Al powder [1171], in the form of zirconia powders mixed with the Al powders or added to the respective master alloy used to produce the alloy powder [1172]. Submicron Si and TiB<sub>2</sub> can also be added to the alloy. While the Al-Zn-Mg-Cu alloy exhibited macrocracks and poor fracture strength and elongation, the alloy with the addition of submicron Si and TiB<sub>2</sub> revealed no hot cracking and showed improved fracture toughness. Adding Si increased fluidity and reduced solidification shrinkage [1173].

Al-Si cast alloys (e.g. AlSi<sub>7</sub>Cu<sub>6</sub>Mg) are widely used in the automotive industry for engine block applications [1174]. Mixing them as aluminum cast scrap with 7xxx wrought alloys is not favorable due to the large amount of Si they contain, which will be detrimental due to the formation of coarse brittle Si-rich IMCs in the microstructure.

The detrimental effect of brittle IMCs on the mechanical properties of 7xxx series alloys can also be mitigated by severe plastic deformation. Large-strain-constrained plastic deformation breaks the intermetallics into smaller particles of smaller aspect ratio, and randomizes their distribution. All these effects may combine to improve the fracture properties of Al-alloys, as shown for friction stir processing [884]. In the processing by high-pressure torsion (HPT) of an alloy similar to 7449 supplemented with various levels of Fe up to 1.2 wt%, it was shown that a combination of tensile strength > 850 MPa and of ductility > 5% could be maintained independently of Fe content [1175]. In simpler Al-Fe binary alloys, processing by HPT can dissolve a significant portion of the Fe intermetallics into a supersaturated solid solution [1176]. Upon subsequent heat treatment, this solid solution decomposes into extremely fine Fe-rich precipitates, thus providing high strength, high thermal stability (due to the slow diffusion of Fe), and an absence of large intermetallic particles.

Nevertheless, the profound effects of intermetallics may be weakened by producing Al 7xxx wrought alloys with higher solidification rates or by modifying the composition to incorporate elements such as Mn into Fe-bearing intermetallics. Once coarse intermetallic particles are present in the 7xxx wrought microstructure, these not only tend to impair mechanical properties and corrosion resistance but may also generate difficulties in processing.

High-strength 7xxx series alloys are susceptible to hot tearing during extrusion due to large precipitated particles, which can raise the flow stress. An increase in the flow stress during extrusion can increase the actual temperature of the extrudate to above the solidus temperature, causing local melting and crack formation. The multiplication of the two factors (presence of large particles and local melting) may result in hot tearing. The applicable speed of extrusion is therefore determined by alloy composition and the microstructure, which can be modified by homogenization treatment. In the 7xxx series alloys, Al-Fe-Mn-Si particles located at grain boundaries and Al-Mg-Zn-Cu eutectic phases can lower the extrusion speed. Al-Fe-Mn-Si particles located at the grain boundaries (GB particles) are brittle and hard, and are therefore detrimental to mechanical properties and extrudability [1177]. The effectiveness of further processing, in terms of time consumption and the energy required for hot working and cooling, might therefore also be considered when choosing appropriate 7xxx recycling alloys with (preferably) higher local solidus temperatures and fewer large particles.

### 7.7.2. Recycling of 7xxx series alloys with and without pre-sorting

Das and Kaufman [1158] discussed the possible recycling of aerospace alloys, mainly the 2xxx and 7xxx alloy series. The pre-sorting of alloys within the 7xxx series and 2xxx series from aircraft sitting in graveyards may generate composition R7XXX, which is similar to the 7075 alloy after recycling (Table 23). Such an alloy may be usable in non-fracture-critical aircraft parts. No pre-sorting of 2xxx series alloys with 7xxx series alloys can result in composition R2 + 7XX, which is shown in Table 23 [1158]. Although the pre-
sorting of 2xxx series alloys and 7xxx series alloys should be technically achievable due to the higher Cu-content in 2xxx series alloys and higher Zn content in 7xxx series alloys [844], lack of proper pre-sorting (e.g. due to possible shredding inaccuracy) should also be addressed.

The Zn content in the recycled alloy without pre-sorting (3.0 wt%, R2 + 7XX) is expected to be approximately half of that in the resulting recycled alloy after pre-sorting (6.0 wt%, R7XXX).

### 7.7.3. The role of Zn and Mg in 7xxx series alloys

The 7xxx alloy series is the only Al-alloy series in which Zn is used as an intended alloying element in amounts of 1 to 8 wt% [1257]. This can explain a dilution of Zn after mixing 7xxx series alloys with Al-alloys of other series, as shown in the composition of R2 + 7XX compared to R7XXX. From composition R2 + 7XX it can also be observed that the Mg content is lower (1.8 wt%) than in composition R7XXX (2.5 wt%), because the Mg content of many 2xxx series alloys is lower compared to 7xxx series alloys [1158]. Since Zn in combination with Mg or Mg + Cu provides the high levels of strength in the 7xxx series, lower Zn/Mg ratios may influence the final properties of the product.

It has been shown that – among alloys with Zn + Mg content of ≈ 7 wt% – a decrease in the Zn/Mg weight ratio from 13 to 5 leads to an increase in both ultimate tensile strength and yield strength of about 30%. This was ascribed to a higher density of precipitates in the matrix [1178]. Electrochemical cyclic polarization tests in 3.5% NaCl solution revealed a higher corrosion potential and lower corrosion current density in low-Zn/Mg-ratio alloys than in high-Zn/Mg-ratio alloys, and hence higher corrosion resistance and higher stress corrosion cracking (SCC) resistance. Both are explained by the lower Zn content in the matrix and the narrower precipitate-free zones (PFZ) [1178]. Experimental Al-Zn-Mg-Cu alloys with a total Zn + Mg content of ≈ 4.8 wt% and Zn/Mg ratios between 2.3 and 0.9 exhibited the same peak hardness level after heat treatment. High-Zn/Mg-ratio alloys achieved peak hardness faster during heat treatment compared to low-Zn/Mg-ratio alloys. This was attributed to the nucleation of a larger number of GP II zones and their transformation to the η' phase in high-Zn/Mg-ratio alloys. However, after reaching peak-aged condition the rate of hardness reduction increased with higher-Zn/Mg ratios during overaging [1179]. The 7xxx series alloys are often used in a slightly overaged temper to achieve a trade-off between strength, corrosion resistance, and fracture toughness [1257]. Hence the Zn/Mg ratio should be considered during the production of recycled 7xxx series alloys and preferably kept in a tolerance range where similar applied heat treatments will lead to comparable material properties required for designated applications.

Zn may be removable from the melt by vaporization due to the high vapour pressure, and Mg can be removed by chlorination using Cl₂ gas [844], so it may be technically possible to decrease the Zn or Mg content to adjust the Zn/Mg ratio.

### 7.7.4. The role of Cu in 7xxx series alloys

The 7xxx alloy series can be divided into Cu-free alloys and alloys with Cu addition. The latter can have a significant influence on

<figure><img src="image_80.png" /><figcaption>**Fig. 78.** Distribution of second-phase particles of cold-rolled 7075 alloy samples produced by TRC (a,b) and PMC (c,d). Reprinted with permission from [1169].</figcaption></figure>
final strength and corrosion resistance. An increase in Cu content from 0 to 1.6 wt% in a 7xxx series alloy can generate an increase in tensile strength from 515 MPa to 574 MPa, presumably due to a greater amount of the main strengthening η' phase [1180]. Because they provide both mechanical strength and high corrosion resistance, Cu-rich 7xxx series alloys are a more appropriate choice for structural materials for aircraft than Cu-free alloys. However, it has been shown that Cu in intermetallic particles, the solid solution in the matrix and precipitates at grain boundaries can have differing effects on corrosion behavior, which might be adjustable by heat treatment. In under-aged condition the 7055 alloy showed a decrease in Zn in the matrix, leading to a more noble matrix and hence decreasing the galvanic coupling between Al<sub>7</sub>Cu<sub>2</sub>Fe particles and the matrix. In overaged temper condition, however, a higher pitting corrosion trend was observed, presumably due to the reduction of Cu in the matrix, which enhances the galvanic effect between Al<sub>7</sub>Cu<sub>2</sub>Fe particles and the matrix. While a higher Cu content in the matrix would therefore lead to higher pitting corrosion resistance, less Cu in grain boundary precipitates would increase the susceptibility to SCC and intergranular corrosion [1181]. Alloys with higher Cu and Zn content, such as 7050, show more S phase (Al<sub>2</sub>CuMg) to be dissolved in homogenization heat treatments. The presence of the S phase in combination with high heating rates can result in overheating, causing the formation of rosettes. An almost complete dissolution of the S phase should therefore be aimed for, since pre-melting could otherwise take place in the following hot working processes at temperatures below 480° C [1182]. Several studies focus on suitable homogenization treatments for 7xxx series alloys to dissolve phases such as the S phase. These have led to the development of multi-step homogenization treatments [1183,1184]. In addition, Cu in combination with Zn and Mg in the 7xxx series alloys generates enhanced susceptibility to weld cracking [1185]. Hence, mechanical and corrosion properties may to a certain extent be adjustable via specific heat treatments, and lead to a higher compositional tolerance range for possible Cu dilution or Cu addition during the recycling of Al scrap material in the 7xxx alloy series.

### 7.7.5. The role of transition elements in 7xxx series alloys

In 7xxx series alloys transition element additions such as Cr and Zr can be present in small quantities, primarily to inhibit recrystallization during processing and to maintain the resulting grain structure after solution heat treatment. Elongated grain structures in a 7xxx series alloy after hot rolling may increase resistance to SCC in extrusions, because the propagation of intergranular SCC cracks perpendicular to the elongated grains is difficult [1182]. It may be necessary to apply a two-step homogenization treatment to achieve an improved Al<sub>3</sub>Zr dispersoid distribution to almost completely inhibit recrystallization with < 0.09 wt% Zr [1186]. The potential accumulation of grain-refining elements such as Cr or Zr during recycling needs further investigation [1158]. However, removal of Ti and Zr might be possible with an addition of B to generate TiB<sub>2</sub> and ZrB<sub>2</sub> stable precipitates, which could be filtered [844].

### 7.7.6. "Green" 7xxx series alloys

While recycled high-strength 7xxx series alloys may not achieve the properties required for critical parts in the aircraft or automotive industries, they may be applicable in less critical applications such as consumer electronic devices, and easier to introduce. Because customers expect authentic materials, TVs are being developed with thinner profiles and the size of tablet computers is increasing. The use of Al-alloys could be increased further compared to plastic materials. Anodizing 7xxx series alloys does not generally produce the appearance desired. Nevertheless, a pure Al-clad material with a better surface finish around a 7xxx series core alloy bonded by hot rolling may be appropriate in consumer electronics [1187], and may have potential as a herald of "green" high-strength Al products even as recycled old scrap.

## 8. Future alloys, products and processes designed for high levels of scrap usage

### 8.1. Sustainable aluminum alloys and powder recycling procedures in additive manufacturing

#### 8.1.1. Introduction

Additive manufacturing (AM) comprises fabrication processes that are characterized by a layer-by-layer build-up of 3-D objects, whose geometrical information is provided in the form of computer-aided design models. In Powder Bed Fusion (PBF), a laser (L-PBF) or an electron beam (Electron Beam Melting (EBM)) is used to melt and consolidate a powder. Here we focus mainly on L-PBF, which is currently the most common technique in metal AM.

The geometrical freedom provided by L-PBF allows optimization of lightweight parts with better design and functionality than existing components. This in turn offers the potential to reduce energy consumption in the aerospace and automotive sectors [1188–1190]. On-demand production of parts reduces transportation in the supply chain [1188] and waste associated with large inventories [1188]. Because unmelted powder can be reused for subsequent building jobs, less waste material is produced during the

<caption>**Table 23**</caption>

The 7075 alloy and potential compositions in wt.% of recycled aircraft alloys, assuming both pre-sorting (R7XXX) and no pre-sorting (R2 + 7XX) [1188].

<table><thead><tr><th>Alloy</th><th>Al</th><th>Zn</th><th>Mg</th><th>Cu</th><th>Fe</th><th>Mn</th><th>Si</th><th>Others</th></tr></thead><tbody><tr><td>7075</td><td>≈ 90</td><td>5.6</td><td>2.5</td><td>1.6</td><td>max. 0.5</td><td>max. 0.4</td><td>max. 0.4</td><td></td></tr><tr><td>R7XXX</td><td>≈ 90</td><td>6.0</td><td>2.5</td><td>2.0</td><td>0.4</td><td>0.2</td><td>0.2</td><td>0.2</td></tr><tr><td>R2 + 7XX</td><td>≈ 92</td><td>3.0</td><td>1.8</td><td>3.0</td><td>0.4</td><td>0.4</td><td>0.4</td><td></td></tr></tbody></table>
layer-by-layer construction of the part [1188,1191]. This contrasts with traditional subtractive manufacturing, where material is selectively removed in order to produce a component. Particularly striking examples include components of complex geometries such as those used in aerospace applications, where buy-to-fly ratios (i.e. the ratio between the weight of the raw material and the weight of the final component) of > 95% are common [1192]. A decrease in the amount of waste material results in a reduced CO₂ footprint. The specific energy consumption of an efficient AlSi10Mg build in which extensive utilization of the available space on the building platform was realized was measured by Faludi et al. at 566 MJ/kg [1193]. In primary production, the embodied energy of bulk AlSi10Mg (currently the most-used aluminum alloy in additive manufacturing [1194]) is ~ 189 MJ/kg [1195], while atomization adds an additional ~ 8 MJ/kg to the embodied energy of the material [1195]. A non-significant proportion of the CO₂ footprint associated with parts produced by L-PBF therefore arises from steps preceding the AM process. This highlights the importance of powder recycling.

### 8.1.2. Powder recycling

Metal powders can be reused in PBF in a variety of ways [1196]. Used powder can be sieved and reused without mixing with other powders; this is currently the most commonly employed strategy documented in the scientific literature [1196]. Used powder can also be sieved and mixed with virgin powder or mixed with used powder of the same age [1197]. In the latter context Lutter-Günther et al. [1197] modelled the maximum number of powder reuse cycles as a function of powder batch size, build volume utilization, powder coating dosage factor and powder loss. The build volume utilization was found to have the greatest impact on the maximum life cycle, which varied between 1 and 117 cycles for all combinations of parameters tested [1197]. The average maximum number was calculated at 35 cycles for batch sizes of 100 – 1000 kg [1197]. This provides a reference value up to which powder recycling should be assessed for typical L-PBF applications [1196]. During the fabrication process, however, several mechanisms can affect the properties of metal powders [1198], leading to potential quality degradation and therefore to fewer achievable reuse cycles.

A well-documented phenomenon is the formation of laser spatter. The evaporation of material resulting from intense local interaction of the laser with the metal powder is associated with substantial recoil pressure, which can lead to melt pool instabilities and ejection of material from the melt pool [1199]. Spatter can also be formed if powder particles surrounding the melt pool are ejected due to the intense vapour flow or by the entrainment of particles within the inert gas flow present inside the building chamber [1200]. Individual spatter particles are typically spherical [1201]. If liquid spatter particles collide with other particles either during their flight or after contact with the powder bed, aggregates can form [1202]. In addition to the relatively coarse spatter particles, the partial evaporation of metal powder causes condensate to form, consisting of submicron particles deposited on the walls of the building chamber and on the surface of spatter particles [1202]. Apart from changes in granulometry and morphology, i.e. in the particle size distribution and morphology of the metal powders used, the chemical composition of spatter may differ from that of virgin powder. This was observed by Simonelli et al., who found surface oxides on laser spatter of an AlSi10Mg powder [1201]. While large agglomerates can be effectively filtered out during sieving, the majority of laser spatter particles, as well as most metal condensate, will be transferred to subsequent build jobs when recycled, potentially having negative effects on built parts. This is mainly because laser absorption and heat transfer, which govern the formation of melt pools [1203], are strongly dependent on the properties of the powders used. The absorptivity of metal powders is generally higher in comparison to flat surfaces due to multiple scattering [1204], with the absolute value for the absorption coefficient depending on the powder size distribution and the resulting packing density [1204]. Similarly, the thermal conductivity of metal powders is dependent on their size and shape [1205]. In the powder size ranges typical for L-PBF (<100 µm), densely spaced solid-to-solid contact points contribute significantly to thermal resistance [1206]. Conductivity therefore depends on the nature of contact points and voids between the powder particles. Similarly, surface oxides may affect both the thermal conductivity of the powder feedstock [1207] and the phase constitution of the final alloy due to the incorporation of oxides, thereby affecting reproducibility with an increasing number of powder reuse cycles.

Tradowsky et al. fabricated specimens of AlSi10Mg by L-PBF and subsequently analyzed them [1208]. After initial processing, the leftover powder was reused in a separate build job without prior sieving [1208]. The void content increased drastically, from 0.26% for a specimen made from virgin powder to 2.00% for a sample made from the recycled powder. The increased presence of irregularly shaped voids in the sample produced from the recycled powder was ascribed to the formation of significant amounts of oxides. Yield strength and ultimate tensile strength decreased by approximately 8%, while the elongation at fracture decreased by roughly 14% [1208].

In another study, Asgari et al. investigated the effect of one powder reuse cycle on the properties of AlSi10Mg powder and the properties of built parts [1210]. In contrast to the study by Tradowski et al. [1208], the powder was sieved after the first use cycle. No statistically significant difference in the average particle size was found between virgin and recycled powders [1210]. Moreover, the yield strength, ultimate tensile strength and elongation at fracture of a specimen made from recycled powder showed no significant variations from published values for samples made from virgin powders [1210], indicating that using a recycled powder did not negatively affect the quality of the built parts. Similarly, Hadadzadeh et al. investigated the effect of one powder reuse cycle on AlSi10Mg powder characteristics and the resulting part properties [1211], employing the same recycling strategy of sieving and reusing without mixing with virgin powder. Only minor variations were found in the average powder particle size (standard deviation ≈ 9 µm in the virgin powder and ≈ 10 µm in the recycled powder), solidification cell size and dislocation density. While the difference in yield strength between vertically-built virgin and recycled powder samples was found to be negligible, the latter exhibited a lower elongation to fracture [1211].

The experimental studies mentioned above only covered one reuse cycle. To quantify the effect of multiple reuse cycles on the mechanical properties of AlSi10Mg, Del Re et al. measured the powder properties and mechanical properties of L-PBF-processed parts for a total of 9 subsequent build jobs using the same powder batch, which was sieved after each build [1209] (see Fig. 79). Notably, the
chemical composition, apparent density, tap density and particle sphericity of the AlSi10Mg powder were found to remain approximately constant for the full duration of 8 reuse cycles [1209]. The oxygen content was not measured. The particle size distribution exhibited a shift towards smaller particle sizes, especially during the first cycles (see Fig. 79a and b) [1209]. This behavior was attributed to comparatively coarse particles being gradually removed from the powder batch due to sieving [1209]. The results of tensile tests are displayed in Fig. 79c and d. No significant effect of the powder recycling on elongation at fracture was observed. In contrast, a statistically significant reduction in yield strength with an increasing number of reuses was observed. However, the total extent to which the yield strength decreased over 8 reuse cycles was small. In addition to the decrease in strength, high cycle fatigue strength decreased by approximately 12% [1209]. According to the authors, the changes in mechanical properties observed still conform to the alloy specifications [1209]. The results of Del Re et al. [1209] regarding multiple reuse were largely confirmed in a separate study by Maamoun et al. [1212], who investigated the characteristics of AlSi10Mg powder subjected to 18 prior build jobs (see Fig. 80). Relative to the virgin powder, the recycled powder exhibited a shift towards smaller particle diameters. The overall degree to which the particle size distribution shifted was small, as values for the 10th (D10) and 90th (D90) percentiles only decreased from 11.8 and 54.1 µm for the virgin powder to 10.7 and 52.6 µm for the recycled powder, respectively.

The findings of the studies presented suggest that powder recycling is a viable solution for improving the eco-efficiency of additively manufactured components, provided that used powder is sieved after each build job. However, more research is needed to quantify the maximum numbers of reuse cycles up to which properties of build parts will continue to retain their respective specifications. Because certain applications, e.g. in the aerospace industry [1194], require mechanical properties beyond those offered by AlSi10Mg, it is also necessary to investigate the effects of powder recycling using different aluminum alloys, as recycling behavior may differ according to the chemical composition of the powder.

### 8.1.3. Sustainable aluminum alloys for additive manufacturing

The vast majority of aluminum powder is produced from primary aluminum [1213]. Apart from reusing powders, the CO₂ footprint associated with typical AM processing might be further reduced by employing secondary aluminum as the raw material. Due to the currently high price of spherical metal powder [1198] and the inability to process many commercial high-strength aluminum alloys [1194], comparatively little attention has so far been paid to developing sustainable aluminum alloys for AM.

In general, impurities can exert harmful effects on aluminum alloys in a variety of ways, some of which are universal among manufacturing methods. More specifically, the interaction between impurities and other solute species can cause drastic changes in phase equilibria. In high-strength Al-Cu-Mg-Ag alloys, for instance, the addition of only 0.2 wt% Si is sufficient to fully suppress the formation of the desirable Ω phase [1215]. Interactions between different solute species depend on the chemical composition and need to be considered for the various alloy systems individually. However, L-PBF provides cooling rates on the order of 10⁴ – 10⁶ K/s [1216], and some mechanisms by which impurities adversely affect the properties of aluminum alloys can be mitigated or even fully eliminated because rapid solidification is known to increase the effective solubilities of alloying elements in aluminum and decrease the size of second-phase particles [1217]. Clear examples of this are provided by studies on the AM of Fe-bearing aluminum alloys.

Suchy et al. [1009] compared the microstructure and mechanical properties of L-PBF-produced and as-cast AlSi9Cu3, an alloy frequently used in the production of die-cast engine blocks in the automotive industry [1218]. The microstructure of the as-cast alloy containing 0.65 wt% Fe was characterized by coarse eutectic Si and intermetallic Fe- and Cu-rich phases [1009]. The intermetallic phases in the L-PBF-processed material had a significantly finer dispersion, despite the higher Fe content of 1.01 wt% [1009]. This microstructural refinement was accompanied by a significant improvement in mechanical properties. The elongation at fracture increased from approximately 0.6% in the cast specimen to 5.3% in the sample produced by L-PBF. Moreover, the yield strength almost doubled. A similar improvement in mechanical properties was noted by Yamasaki et al. [1214], who produced AlSi12CuNiMg with a nominal Fe concentration of 0.1 wt% via both casting and L-PBF. Two modified alloys with nominal Fe concentrations of 3 and 5 wt% were also processed by L-PBF. All specimens were annealed at 300 °C [1214]. One group of alloys was heat-treated to the T6 condition prior to this step. High-temperature tensile tests were performed at 300 °C; the results are displayed in Fig. 81. Among the samples annealed at 300 °C, those produced by L-PBF exhibited enhanced strength and elongation at fracture at all Fe concentrations. Similarly, the elongation at fracture was drastically increased in the additively manufactured samples of the T6 group, indicating that Fe addition can be deployed to improve the strength-ductility relationship in L-PBF-processed Al-Si-based cast alloys.

Remarkably, an improvement in mechanical properties when Fe is added has not only been observed in alloys of the Al-Si system. An AM-related patent application published by AMAG casting GmbH and AUDI AG [1220] discloses the chemical compositions of commercial aluminum alloys modified by adding Mn and Fe, and presents the properties of L-PBF-processed materials. Adding 0.5 – 1 wt% Fe to Mn-modified wrought alloys of the 5xxx series (Al-4.9 Mg-0.6Mn-0.5Sc-0.5Zr-0.1Fe), the 6xxx series (Al-1.2 Mg-0.8Si-0.7Mn-0.2Fe), and the 7xxx series (Al-6Zn-1.6Cu-2.7 Mg-0.3Mn-0.1Fe) generated increased ultimate tensile strength, elongation at fracture, reduced porosity and hot cracking susceptibility in all alloys [1220].

Even higher Fe concentrations can be found in the rapid solidification alloy 8009 with the composition Al-8.5Fe-1.3 V-1.7Si (wt.%) [1219], originally developed by Skinner et al. [1221]. Fig. 82 provides a direct comparison between cast and L-PBF-produced 8009 [1219]. Despite the high Fe content, no coarse intermetallic phases can be observed in the additively manufactured alloy. It achieves an ultimate tensile strength of 454 MPa and an elongation at break of 7.2% [1222]. This provides a notable example of how high concentrations of otherwise deleterious alloying additions can be exploited to create high-strength alloys for AM.

The available literature suggests that many of the commercial aluminum alloys currently in use exhibit high impurity tolerance to Fe during additive manufacturing. This clearly illustrates the potential of using secondary aluminum as raw material in the production of aluminum-based powders.

Future research should aim at quantifying the effects of different levels and combinations of impurity elements (e.g. Fe and Mn) –
<figure><img src="image_81.png" /><figcaption>**Fig. 79.** Evolution of (a) particle size distribution in an AlSi10Mg powder; (b) average particle size; (c) tensile yield strength of AlSi10Mg in the as-built state; and (d) elongation at fracture with an increasing number of reuse cycles. The powder was sieved after each build job. Adapted with permission from [1209].</figcaption></figure>

not only on mechanical properties, but also on corrosion behavior and the processability of aluminum alloys. This will be a step towards determining the grade-specific maximum tolerable concentrations of these impurities.

## 8.2. Uni-alloys, crossover alloys and multi-purpose alloys

First, we will define what we mean by the terms uni-alloys, multi-purpose alloys and crossover alloys [48]. A uni-alloy addresses a single composition space and targets a reduced number of alloy compositions for different parts and components of a system made from it. In the simplest example, this can be the use of a single composition for can bodies and lids [31,299]. Its current extreme is a uni-alloy car, which is currently far from being realized. However, an interesting example that actually exists is car doors. Currently, the inner part of car doors is often made of non-hardenable, highly formable 5xxx wrought alloys [1223]. Unfortunately, these alloys are soft and their surfaces after forming is not decorative. Therefore wrought 6xxx alloys, which show good surface quality after being shaped but are more limited in formability compared to the 5xxx alloys, are often used for the outer part of the door [1223]. However, the combined recycling of these materials is problematic because the alloys cannot be recycled in the same scrap stream due to elemental incompatibilities (see section 7.5 and 7.6 for details). Although it is obvious that a uni-alloy strategy can reduce compositional complexity in future structures, it should be noted that in very simple cases, such as beverage cans, it does not always provide environmental benefits if certain boundary conditions apply [31]. However, in many complex structures and systems the strategy is indeed beneficial. A uni-alloying concept could remedy this situation and provide more sustainable, simplified recycling.

In this context it is very important to recognize that such concepts can only be realized if the properties of an alloy can be meaningfully changed via microstructure adjustment. To a small extent, this has been investigated in the area of 6xxx alloys in cars. Some 6xxx series alloys have been promoted as uni-alloy in the automotive sector (e.g. 6022 for outer and inner body panels [29,1224]). In reality, however, a mixture of alloys is still often used due to the varying demands of today's automotive designs and the lack of studies and development work on 6xxx-based automotive uni-alloys.

Uni-alloy concepts exist in the area of AlSiMg structural castings. These concepts mainly apply to the use of single compositions which provide a broad property range via adjustments to the heat treatment condition. For example, Müller et al. [1225] developed a high-pressure die-cast alloy for various car components classified as “crash-relevant”, “strength-relevant” and “high-strength”. At present different near-eutectic AlSi alloys are deployed in such components depending on the application, but in the future the new Si-
<figure><img src="image_82.png" /><figcaption>**Fig. 80.** SEM images showing the morphology of AlSi10Mg powder in the (a, b) virgin state and (c, d) after 18 prior uses. The powder was sieved after each build job. Reprinted with permission from [1212].</figcaption></figure>

reduced AlSi7Mg variant should meet all requirements. For this purpose, the heat treatment of the alloy was redesigned. Adjusting the solution heat treatment and increasing the quenching rate generates a high-strength condition. To be “crash-relevant”, where increased ductility is required, a share of the component’s strength is sacrificed by applying an overaging treatment. This procedure is further associated with an increase in necking elongation and a simultaneous improvement in the multiaxial deformability of the material, resulting in an improvement of the often application-critical self-pierce riveting properties [1225].

The aspect of setting different properties by modifying process parameters also applies to multi-purpose alloys. Multi-purpose alloys are particular alloys which can be deployed in several applications. In addition, one alloy composition can be subjected to different production methods (e.g. castings and wrought products are made from the same alloy). To a small extent multi-purpose alloys already exist today: for example, the wrought alloy 5182 overlaps closely with the casting alloy EN AC-AlMg5. Other examples include the use of both casting and wrought alloys in additive manufacturing, for example AlSi10Mg [1226] or 6061 [1227]. These have in common that due to a high solidification rate and the associated fine formation of IMCs, they tolerate a higher trace element content.

The term “crossover alloy” [48] implies that the compositional classifications usually in place for the different alloy classes are neglected for bridging properties. This applies to both the uni-alloy and the multi-purpose concepts. The aim of the crossover approach is to expand the property range to reduce the number of alloys required in complex systems [3,1228].

Production and design often require the use of several different alloy systems, limiting recyclability at the end of a product's service life [3]. One example is a complicatedly shaped component in a vehicle. Although it is possible to improve the formability of 6xxx alloys without loss of strength [1229-1231], this development direction has natural limits because strain hardening of these alloys cannot be easy modulated over a wide range. Hence, 6xxx alloys do not yet meet some of the requirements for complex forming operations (especially stretch forming), even though these alloys offer good age-hardening potential. Today's wrought 5xxx alloys offer high uniform elongation and work hardening due to their inherently high Mg content [1018], which is beneficial for complex forming operations. However, Lüders elongation, dynamic strain aging and low strength are disadvantages of these alloys [86,87].

One approach to solving this problem is to introduce additional alloying elements into 5xxx alloys to promote age hardening. The addition of small amounts of Cu to 5xxx wrought alloys to prevent softening due to recovery of automotive 5xxx alloys during paint
<figure><img src="image_83.png" /><figcaption>**Fig. 81.** True stress–strain curves of an AlSi12CuNiMg alloy with varying Fe content, produced by casting and L-PBF (denoted here as Selective Laser Melting (SLM)), and measured at 300 °C. Reprinted with permission from [1214].</figcaption></figure>

<figure><img src="image_84.png" /><figcaption>**Fig. 82.** SEM micrographs of (a) as-cast and (b) L-PBF-produced 8009. Adapted with permission from [1219].</figcaption></figure>

baking was first seen in Japan at the beginning of the 1990s [1232], and has attracted further interest since [1052,1109,1233]. This research may involve the first exploitation of the extensive work hardening of 5xxx alloys combined with the age-hardening potential of another alloy class (Cu from the 2xxx series) and the crossover of properties. As in 2xxx series alloys, the typical precipitation sequence in such crossover systems includes clusters, precursors and finally the equilibrium S phase (Al₂MgCu) [1233].

While the Cu additions to wrought 5xxx alloys investigated were generally rather small, other studies addressed Zn addition to 5xxx alloys of up to several percent. Compositionally the results can be regarded as crossovers of 5xxx and 7xxx series alloys. Their aim is not only increased strength, but also reduced serrated flow behavior and enhanced corrosion resistance. In recent years, an increasing trend towards adding Zn [1051,1234–1238] has been observed, although with an even stronger focus on the combined adding of Zn and Cu [87,1006,1239–1245]. Note that this concept has also been applied to casting alloys [1005].

It is now accepted that in equilibrium, the T phase (Mg₃₂(Al,Zn)₄₉) is formed in these systems. However, the precipitation sequence utilized in aging treatments is complex and still under debate. It depends on thermal history and composition. This is illustrated in Fig. 83.

Fig. 84 demonstrates the potential of this type of 5xxx to 7xxx crossover concept. The combination of high strain hardening caused by high Mg content and a rapid, superior age-hardening response within a short time, typically applied upon paint-baking in the automotive industry, proves to be feasible [1239]. Although the concept is promising, it has not yet reached industrial scale due to the recent nature of intensive research in this field.

Compositionally, crossovers between 6xxx and 7xxx series wrought alloys also exist. The main aim with these crossover alloys is to improve the age-hardening potential of 6xxx alloys via Zn addition and to add the aging characteristics of 7xxx series alloys to 6xxx alloys [405,831,861,862,1246–1249]. As can be seen from the literature on it, this research was not undertaken to reduce the number of alloys required via an expansion of the property profile, but the results may well serve this goal.

In general, the environmental benefits of the above-mentioned crossover alloys have not been studied in depth. There are also no studies on their recyclability, apart from the general argument – which should be noted by product designers – that reducing the