Corrosion Behaviour of Recycled Aluminium AlSi9Cu3(Fe) Machining Chips by Hot Extrusion and Thixoforming

Corrosion Behaviour of Recycled Aluminium AlSi9Cu3(Fe) Machining Chips by Hot Extrusion and Thixoforming

3.1. Materials Characterisation

Figure 2 shows the microstructure of the tested samples. RS shows a characteristic dendritic microstructure of primary α-Al surrounded by an acicular morphology of the eutectic phase (composed of Si and IMCs) (Figure 2a). However, the plastic deformation of the machining chips changes the microstructure of the sample, whereby the dendritic microstructure disappears completely, and the crystal grains are considerably fragmented. These changes can be seen in DHES (Figure 2b), which has a very uniform microstructure with refined crystal grains and IMCs.
However, the aim is to obtain a TFS directly from RS machining chip scrap (without remelting), and the confirmation of the qualitatively obtained sample is the globular (spherical) microstructure. As can be seen, TFS has a round, globular microstructure with an average globule size of about 50 µm (Figure 2c).
It is well known that the microstructure of various Al-Si alloys contains a matrix of dendritic α-Al surrounded by a eutectic (or polyhedral) Si phase and various secondary phases or IMCs, which are additionally enriched with Fe and Cu [37,38,39,40,41]. Such a complex structure has a significant influence on the mechanical and corrosive properties of the tested material. To better analyse the microstructure and determine the distribution and composition of the IMCs, the investigated samples were examined using SEM/EDS analysis. Figure 3 shows the surface texture of RS, DHES, and TFS at different magnifications (with the indicated area where the EDS analysis was performed; ochre rectangle, Table 2).
As can be seen in Figure 3, separate Si phases form on the RS around the dendritic α-Al (dark areas, spectrum 1) in the form of large coarse clusters of “flecks” and needle-like formations (white formations, spectrum 2), then Cu-rich phases (wrinkled structure, spectrum 3) and needle-like IMC formations, which are rich in Fe, Mn, Cr, and Si (greyish structures, spectrum 4).

In DHES, the eutectic Si and other IMCs are highly refined. DHES is produced directly from machining chips of RS, whereby the material undergoes several stages of plastic deformation. In the first step, i.e., the machining of RS ingots, chips are produced, and the material is subjected to severe deformation (the crystal grains and the intermetallic phases are cut into smaller pieces). The chips are then compacted in the second step in the form of a roller, which causes further deformation of the material and refinement of the crystal grains and coarse Si phases, as well as other IMCs. Finally, the DHE process produces a chip-based semifinished product (or DHES). It can be observed that the eutectic Si and IMCs are highly comminuted and uniformly distributed in the α-Al matrix, and the overall microstructure of the alloy is extremely homogeneous.

By reheating the DHES in the semisolid region and thixoforming in the SSMP tool, such a highly deformed structure is transformed into a globular TFS structure, which was confirmed by the development of circular boundaries between the individual α-Al phases (observed at a magnification of 1000×). It can also be observed that the α-Al phases have a diameter of ≈50 µm and that polyhedral and refined IMCs (instead of needles) accumulate along their boundaries.

3.2. Corrosion Behaviour

Electrochemical methods, i.e., OCP, EIS, and PD, were used to investigate the corrosion resistance of RS in 0.5 M NaCl at different stages of its recycling (plastic deformations through DHE and TF steps).

Figure 4 shows the change in the OCP curve during exposure of the test samples to NaCl solution for 1 h. Although the OCP curves are similar for all samples, RS shows certain differences. During the first 10 min, the RS potential rises sharply from −1.0 V to a value of about −0.9 V, after which it remains more or less constant with slight fluctuations. These changes can be attributed to the forming of a surface oxide layer in contact with the aqueous solution. The plastic deformation causes a potential shift in the positive direction. DHES and TFS behave similarly and very quickly reach a stable OCP value (−0.86 V for DHES and −0.88 V for TFS). Since the established OCP value represents the balance between the formation and dissolution of the surface passive film during exposure to the electrolyte solution, it is obvious that plastic deformation has a positive influence on material stability under spontaneous corrosion conditions.

Due to its high sensitivity and non-destructive nature, EIS is an indispensable method for characterising the phase boundary between electrode and electrolyte. When analysing the results obtained (EIS spectra), the phase boundary is described by the electrical equivalent circuit (EEC), which can provide additional information on the mechanism of the corrosion process.

Figure 5 shows the Nyquist and Bode diagrams recorded after 60 min of OCP stabilisation of RS, DHES, and TFS in NaCl solution. The Nyquist diagrams show the existence of two depressed capacitive loops for all investigated samples (Figure 5a). The size and diameter of the capacitive semicircles are significantly larger for DHES and TFS than for RS, indicating increased corrosion resistance. In the Bode plots (Figure 5b), the capacitive properties of all samples are highlighted in a wide frequency range (at f Z|), the maximum phase angles, and the peak area are larger for DHES and TFS compared to RS, confirming the favourable influence of plastic deformation on the corrosion resistance of the Al alloy.
Figure 5 shows the best-fit curves with the assumed EEC (Figure 6; χ2 −3), which consists of two-time constants and is often used to describe the corrosion behaviour of Al alloys subjected to plastic deformation [44,45,46].
The proposed EEC consists of resistance elements (R) and constant phase elements (CPE) as follows: electrolyte resistance (Rel ≈ 6 Ω cm2), surface layer resistance (R1), charge transfer resistance (R2), CPE of the surface layer (Q1), and CPE of the double layer (Q2). According to the size of the coefficients n1 and n2, the CPEs in the EEC represent the deviation of the capacitance from the ideal behaviour [47]. The calculated EEC parameters for the examined samples are listed in Table 3.

For RS, the values of R1 and R2 are 1.36 kΩ cm2 and 3.35 kΩ cm2, respectively, due to the formation of a surface oxide layer. In comparison, the samples that were subjected to plastic deformation have significantly higher values for the parameters R1 and R2. For example, the R1 values for DHES and TFS are 3.21 and 3.02 kΩ cm2, while the R2 values for the same samples are 10.52 and 9.41 kΩ cm2, which is about three times higher than for RS.

In parallel (Table 3), the capacitance of the surface layer (Q1) is lower than the capacitance of the double layer (Q2) for all samples. With plastic deformation, the capacitance of both layers (surface and double layer) decreased. Since the capacitance is inversely proportional to the thickness, the direction of the mentioned changes indicates that the resistance and thickness of the oxide layers on the observed samples increase in this order: RS
The results obtained show that the refined and homogeneous microstructure contributes positively to the formation of a stable oxide film on DHES and TFS and increases their corrosion resistance in an aggressive environment. Interestingly, slightly worse oxide films are formed on TFS (lower R1 and R2 values, higher Q1 and Q2 values) than on DHES, which can be attributed to the accumulation of polyhedral IMCs along the α-Al phase boundaries (Figure 3).
The PD curves of RS, DHES, and TFS in 0.5 M NaCl solution over a wider potential range are shown in Figure 7, while the corrosion parameters (the corrosion potential, Ecorr, the corrosion current density, icorr, the pitting potential, Epit) determined by analysing the polarisation curves are listed in Table 4.
The starting point for analysing the results obtained is the inflexion point on the PD curve, where the cathodic current changes to the anodic current and is determined by Ecorr (indicated by an arrow in Figure 7) and the corresponding value of icorr.
In general, Ecorr becomes nobler with increasing plastic deformation, while icorr decreases (Table 4). Although the influence of plastic deformation of RS (i.e., chips obtained by machining RS) on the Ecorr shift is very small (≈40 mV), icorr decreases by 4–5 times (from 3.50 to 0.67 µA cm−2 for DHES and 0.98 µA cm−2 for TFS).
In addition, pitting corrosion occurs in all samples, as indicated by the increase in current at the breakpoint in the anodic branch of the PD curve (indicated by an arrow in Figure 7). The potential of the predetermined breaking point (also known as pitting potential, Epit) shifts to the anode side as the degree of material deformation increases. In addition, material deformation (by DHE and TF) leads to a larger potential difference (ΔE) between Epit and EcorrE = EpitEcorr) is an indicator of the pitting initiation rate; the higher it is, the slower the pitting initiation rate [48,49,50].
According to the literature, the natural oxide film formed on the surfaces of various Al-Si alloys has the properties of an n-type semiconductor [29,30,51], which favours the adsorption of Cl ions and the formation of soluble chlorides and causes pitting corrosion. Higher ΔE with lower icorr content (Table 4) shows that oxide films with better protective properties (higher resistance and thickness) are formed on deformed samples, which is consistent with the results of the EIS measurements. Based on the results of the electrochemical measurements, it can be generally stated that the corrosion resistance of the tested materials increases in the following sequence: RS

3.3. Surface Analysis after Corrosion Measurements

After reaching the upper limit of anodic polarisation (of −0.4 V), the RS, DHES, and TFS surfaces were examined by optical microscopy and SEM/EDS analysis. It should be emphasised that the corrosion products were not removed from the surface of the samples after the PD measurements.

The optical micrographs (Figure 8) show clear differences in the corrosion damage on the examined samples, with the largest damage on the surface of RS and the smallest on the surface of DHES. This condition is a consequence of the fact that a finer structure (resulting from deformation) increases the corrosion stability of the material.
To better determine the morphology and investigate the influence of the microstructure (i.e., different stages of plastic deformation by DHE and subsequent TF) on corrosion during recycling of the alloy EN AC AlSi9Cu3(Fe), a SEM/EDS analysis was performed. Figure 9 shows the SEM images of the analysed samples at different magnifications (magnification areas are marked with rectangles), while Figure 10, Figure 11 and Figure 12 show the EDS mapping analysis of the corroded surfaces. As expected, anodic polarisation causes the highest damage to the RS at the same magnification. The damage is distributed over the entire surface and is mostly located near the separated Si phases and the IMCs. At higher magnifications, a layer of corrosion products (oxides) is more clearly visible above the damaged metal surface, which was confirmed by the EDS analysis (Figure 10).
The centre of Figure 10 shows the observed SEM image of RS, while the smaller images show the distribution of the individual elements on the surface. As can be seen, the surface of RS exhibits a complex multiphase microstructure consisting of α-Al phases, and numerous intermetallic phases with isolated areas of high Si, Cu, Fe, Mn, and Cr content are observed. The colours corresponding to these elements overlap in different ways (and in different proportions) at certain locations, indicating the formation of different IMCs. Cu appears mainly as an independent phase in α-Al, while the distribution of Fe completely overlaps with the distribution of Mn and partially with Cr. The distribution of Si is not close to any element (except in some places where it coincides with Fe, Mn, and Cr), which indicates the presence of eutectic Si particles. So, according to the results obtained and in agreement with the literature, the formation of Al2Cu and various Fe-containing phases (AlFeSi, Al(Fe,Mn,Cr)Si) with Si is possible [37,38,39,40,41]. The greatest surface damage has been observed in the vicinity of intermetallic phases rich in Cu and Fe, which is associated with a higher oxygen content (i.e., the formation of Al2O3 due to the Al dissolution).
Aluminium alloys are subject to pitting corrosion and intergranular corrosion in aggressive environments, and the conditions for their formation are similar from an electrochemical point of view. As can be seen in Figure 3, the intermetallic RS phases (Cu- and Fe-rich phases and eutectic Si particles) are locally distributed around the dendritic α-Al phase. These phases are more noble than α-Al and can form micro galvanic cells. Due to the differences in electrochemical potential, the interface between these phases is susceptible to pitting corrosion and further intergranular corrosion. The intermetallic phases are more noble than aluminium and support the cathode reaction (hydrogen evolution or oxygen reduction) in the corrosion process and accelerate the corrosion of the surrounding α-Al matrix [31,38,39,41]. According to Monticelli et al. [41], the most efficient cathodic areas in AlSi9Cu3(Fe) alloy are the Al2Cu phase, nobler than the α-Al matrix by about +400 mV. The Fe-rich phases are about +300 mV nobler than the matrix. The potential difference between the eutectic Si particles and the matrix is about +250 mV, but their cathodic activity is lower due to the high overpotential for the reduction reaction.
Pitting corrosion therefore primarily attacks aluminium near the Cu- and Fe-rich phases, and then the Si phases (dangerous combination of large cathode/small anode) form an oxide layer at the edges of the pits that seals the hole. Due to the autocatalytic process (with an increase in acidity and chloride ion content in the pits), corrosion progresses along the grain boundary and penetrates into the deeper layers of the material [49,52,53,54], which can also be seen in Figure 9 and Figure 10.
It was found that the localised corrosion correlates more strongly at the boundaries of die-parting line zones than with neighbouring ones. This is probably due to the continuous network formed by the morphology of the cathodic Al2Cu phase and/or the higher initial levels of Cu and Cr (and thus higher amounts of Cu- and Fe-rich (with Cr) IMCs). The studies of the impact of intermetallic particle size and distribution on pitting corrosion have shown that the presence of the isolated intermetallic particles was inadequate to initiate a stable pitting activity on the alloy surface [41,55,56,57]. Conversely, the existence of these particles in groups or clusters, that form a continuous network in the material (which is also the case in this work), can promote the development of more active pitting.
The morphology of the corroded surfaces of DHES and TFS is also shown in Figure 9. Compared to RS, these samples (at the same magnification) show a less damaged surface (pits or cracks) and a more uniform oxygen distribution, which is particularly pronounced in DHES (Figure 9 and Figure 11). The SEM/EDS results are consistent with the electrochemical measurements, which showed that DHES and TFS have a nobler OCP, a lower corrosion current, and a more stable oxide film than RS.
During the formation of DHES, the RS machining chips undergo a series of deformations (mechanical cutting, compaction, extrusion at elevated temperature with dimensional changes), during which the eutectic Si and intermetallic phases are significantly refined and uniformly redistributed into the fine-grained α-Al matrix so that the resulting DHES microstructure is homogeneous and refined (Figure 3). In this way, a larger number of galvanic microcells are formed between ultrafine Si particles (or IMCs) and α-Al grains, resulting in a favourable small cathode/large anode combination and reducing the susceptibility to pitting corrosion [28]. It should not be forgotten that this process generates a larger number of grain boundaries, a greater proportion of unbalanced grain boundaries, dislocations, and residual stresses in the machined sample. According to the literature, the aforementioned phenomena in the different plastic deformations provide favourable conditions for the formation of stable, compact, and denser oxide films that increase the corrosion resistance of Al [58,59] and Al–Si alloys [28,29,30,31,38,60] in NaCl solution, which was confirmed by SEM/EDS for the DHES (Figure 9 and Figure 11) (in addition to the electrochemical tests). As can be seen in Figure 2 and Figure 3, Si and IMCs are finely dispersed and uniformly distributed in the fine-grained α-Al phase. After anodic polarisation, some small pits can be seen, while the oxide layer largely covers the surface of the sample.
Reheating DHES in a semisolid temperature range and thixoforming in the SSMP tool achieves the desired globular (spherical) TFS structure (Figure 2), which is confirmed by the development of clear circular boundaries between individual α-Al phases (Figure 3). Polyhedral IMC was also observed to accumulate along the boundaries of the α-Al phases, which influences the slightly poorer corrosion resistance of TFS compared to DHES. As can be seen from Figure 9 and Figure 12, the anodic polarisation of TFS leads to the occurrence of pitting corrosion near areas rich in Cu, Fe, or Si and IMCs. A larger amount of oxygen was observed along the damaged parts of the surface.
Electrochemical measurements and SEM/EDS analysis showed that all samples are subject to corrosion attack and that the percentage of surface affected by corrosion is highest in RS. Indeed, RS has a very inhomogeneous structure with a large proportion of cathodic phases, i.e., coarse (flecked and needle-like) Si and IMCs in a less noble α-Al matrix. This RS structure contributes to the development and progression of pitting corrosion by creating a dangerous combination of large cathode/small anode. The plastic deformation reduces the size of all particles and thus the ratio of the surface area of the noble IMCs and Si particles (cathode) relative to the less noble α-Al phase (anode), which significantly improves the material’s resistance to pitting corrosion. In other words, reducing the ratio between the surface area of the cathode and the anode (Ac/Aa) facilitates the passivation of the material and improves the corrosion resistance [28], which increases for the tested materials in the following order: RS

However, information on actual damage (i.e., the depth, width, and volume of pits) can be obtained by recording the surface topography with an optical profiler. To obtain more accurate data, layers of corrosion products were removed from the surface immediately before the measurement. The non-contact 3D profilometer analysis technique is an ideal, user-friendly method for maximising surface investigation when pitting analysis is required, and it also offers the advantages of combined 2D and 3D capability.

The shape, size, and density of the pits were strongly influenced by the type of plastic deformation of the EN AC AlSi9Cu3(Fe) alloy, e.g., as-cast only, DHE, or TF conditions. The anodic polarisation of RS led to the development of large and rough surface damage of different depths and widths which occurred around the cathodic phases (Si and IMCs) according to SEM/EDS analysis. In addition, there is a large amount of damage on the surface with an average depth of 70–100 μm and a width of about 200 μm, which is interconnected with large and rough damage (Figure 13a). Additional analyses showed that the volume of the pits is 1,464,000 μm3 (position 1) and 1,377,000 μm3 (position 2), while the total surface damage is 41%. On the other hand, individual extremely deep (up to 250 μm) but somewhat narrower (up to 100 μm) pits can also be observed (Figure 13b; most likely formed by dissolution in the α-Al matrix with separation of cathodic Cu or Si phases). Many shallow and narrow pits can be observed under such coarse damage. The volume of the deepest pit (position 3) is 1,683,000 μm3, and the surface damage is 20.84%.
Electrochemical measurements and SEM/EDS analysis showed that the surface of DHES is covered with an oxide layer that adheres well to the surface and remains stable even after anodic polarisation. However, after removal of the oxide layer, a surface with a large number of tiny pits can be seen. Figure 14 shows that the pits have an average width of ≈10 μm and a maximum depth of ≈13 μm. The surface is 27.77% damaged overall. Additionally, the marked part of the surface (almost half of the DHES damaged sections) has a pit volume of 10,600 μm3, which is two orders of magnitude smaller than that of RS. The results obtained indirectly confirm the highly separated and refined microstructure of DHES and as mentioned above, the corrosion stability of the sample.
Subsequent heating of DHES in the semisolid temperature range transforms the highly deformed microstructure into a globular structure, and refined IMCs accumulates at the grain boundary of the α-Al phase. The topography of two different segments of the corroded surface for TFS is shown in Figure 15.
A series of tiny round pits with a maximum depth of up to 3 µm and a width of up to 5 µm can be seen along the grain boundary (Figure 15a). The volume of such pits is about 155 μm3 (position 1), and the total surface damage is 10.6%. In addition, larger corrosion areas can be observed with a width of around 200 µm, which, according to the image, were probably formed by the coalescence of individual smaller pits (Figure 15b; at the locations of the accumulation of polyhedral clusters rich in Si, Cu, and IMCs along the α-Al grain boundary). At the same time, individual pits with a depth of up to ≈60 µm and a width of around 100 µm as well as narrow breakthroughs in the sample depth of ≈60 µm and ≈120 µm were detected. In this case, the volume of the pits marked in Figure 15b is 293,500 μm3 (position 2) and 119,600 μm3 (position 3), which is an order of magnitude less compared to RS, while the total surface damage is 14.64%.

There are strong correlations between the microstructure, the corrosion rate, and the corrosion morphology of the analysed samples. The measurements carried out in this work indicate the importance of optical profilometry analysis in revealing the above relationships.

The aforementioned analysis confirmed that the groups or clusters of cathodic phases (i.e., coarse (flecked and needle-like) Si and Cu and Fe intermetallic phases), forming a continuous network in the material, can stimulate the active dissolution and significant damage of RS (pits are connected, spread, and penetrate into the depth of the sample). It was also found that the presence of isolated cathodic intermetallic particles, as well as the finely chopped microstructure, is not sufficient to develop a stable pitting activity on the alloy surface [41,55,56,57].

Plastic deformation of RS machining chips results in the formation of a homogeneous and refined DHES microstructure. In this case, profilometry analysis suggests changes in corrosion morphology. Extremely shallow and small damage is evenly distributed over the surface, indicating that the DHES is subject to general corrosion. Reheating DHES in a semisolid temperature range and thixoforming achieves the desired globular TFS structure, as well as the accumulation of polyhedral IMCs along the boundaries of spherical α-Al phases, which according to profilometry analysis (in addition to the general form of corrosion) causes the appearance and development of local corrosion.

As described in the introduction, secondary aluminium production is necessary because of its economic and ecological benefits. The direct conversion of aluminium chips into semifinished products can eliminate the cost of the remelting process and reduce CO2 emissions. The idea presented in this paper can therefore be described as an innovative process for recycling semisolid aluminium materials that consumes less energy and emits fewer greenhouse gasses into the atmosphere than conventional recycling. The fact that the semifinished products obtained through recycling have a significantly higher corrosion resistance is an additional advantage for the ecological acceptance of the process presented as the cost and environmental impact of corrosion are reduced and the sustainability of aluminium materials is improved. Namely, when analysing the effects of corrosion on society, three factors are usually considered: financial costs, environmental pollution, and direct effects that cause damage or death. For instance, the most recent corrosion research conducted in the United States (by NACE International and the United States. Federal Highway Administration) found that the direct annual cost of metal corrosion in the country is USD 276 billion, equivalent to 3.1% of GDP [61]. Therefore, the presented method helps to reduce the cost and environmental impact of corrosion and improve the sustainability of materials.

Further research is planned to ensure that the process of recycling semisolid aluminium materials is fully justified from a technical point of view. Metallographic analyses will be carried out to determine the average globule size and the shape factor of the samples. In addition, the mechanical properties (e.g., compressive strength, hardness, tensile strength) and the electrical conductivity of the samples will be tested.

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