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Original Article
DOI: 10.1016/j.jmrt.2018.12.014
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Available online 19 February 2019
Mechanical and electrochemical characteristics of solutionized AA 6061, AA6013 and AA 5086 aluminum alloys
E. F. Abo Zeid
Physics Department, Faculty of Science, Assiut University, Assiut 71516, Egypt
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Tables (3)
Table 1. Chemical composition of the studied alloys in (wt.%).
Table 2. The change in the crystallite size and microhardness for all investigated alloys.
Table 3. corrosion current density (ImA/cm2) and corrosion voltage (Emv) of three studied alloys.
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The current study was conducted to investigate the effect of age hardening on three different aluminum alloys: AA 6061, AA6013 and AA 5086. Different characterizing and testing techniques were utilized in this study. The XRD results for 6xxx series demonstrated the Mg2Si formation with different orientations after natural and artificial aging (NAT and AAT), while MgZn2 phase was depicted in AA 5086 alloy. The friction coefficient of all naturally aged Al alloys at room temperature (RT) was found to be higher than the friction coefficient of the artificially aged Al alloys at 175°C. For group series 6xxx the corrosion resistance was decreased after the natural aging at RT, while the corrosion resistance was improved after artificial age treatment at 175°C for 30min (AAT). The reduction in the friction coefficient and the growing in the corrosion resistance of the Al alloys after artificial age hardening make them potentially be used in many industrial sectors.

Age hardening
Mg2Si phase
MgZn2 phase
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Aluminum alloys are utilized in diversity of cast and wrought forms and in different conditions of heat treatment. AA 6061 and 6013 aluminum alloys are considered to be aged hardenable alloys and AA 5086 is accountable to be a strained hardenable aluminum alloy. These alloys exhibit decent ductility, higher strength to weight ratio, and good corrosion characteristics [1–3]. In furthermost requests of Al-Mg based (5xxx) alloys, the work hardening and solution treatment strengthening are significant contributors, and small Cu insertions can enhance the mechanical properties owed to the precipitation hardening. These alloys are employed in many industrial branches including but not limited; beverage cans [1], as car body panels to lessen the weight and enhance fuel economy and emissions [2]. 5xxx with small amount of Cu are talented candidates for these requests due to their outstanding formability, good strength and the benefits of precipitation hardening during paint baking due to Cu additions [3–5]. On the other hand, the induced impurity of Si in most 5xxx series alloys and in most wrought Al alloys besides the 6xxx series lead to the formation of Mg2Si particles which are improved the corrosion resistance of these alloys [6]. The impact of Mg2Si phase on the corrosion performance of Al–Mg alloys were examined [7–11]. It was demonstrated, the strength of pure magnesium is low for utmost industrial applications. In this regard, some elements are alloying to magnesium to enhance the alloy mechanical and corrosion characteristics [11]. The metastable β″ phase was reported to be the prevailing intermediary phase that presented in Al–Mg–Si and in some Al–Mg–Si–Cu based alloys at early stages of aging. The β″ phase is characterized by needle shaped with long axis alongside (100) of the Al matrix and its crystal structure is grounded to the monoclinic system [8]. After aging treatment, some of the needle shaped β″ precipitates are substituted by rod shaped phase of β′. In Al–Mg–Si–Cu alloys, lath shaped precipitates are initiated during over-aging which is called B′ [9]. This lath shaped phase of the equilibrium Q phase is designated as Q′ phase [10]. Contrasting to β′-β transition which includes an alteration in crystal structure from hexagonal to cubic CaF2 structure, the Q′ phase sustains the same crystal structure and morphology as Q form through the over-aged conditions and only its size increases. It was demonstrated, the corrosion behavior is influenced by grain size and grain morphology, constituent particles, existence of precipitates, and grain boundary distribution. The grain shape and orientation have a gigantic role in the penetration depth of an inter-granular corrosion (IGC) attack [12]. Magnesium-rich and iron-rich precipitates are considered the main intermetallics presented in 5086 alloy. Different precipitates are doped in alloy matrix and play a significant role in preventing localized corrosion attack. Mg is solutionized to these alloys to achieve high strength via solid solution hardening. The Magnesium-rich intermetallic particles comprise principally the β-phase (Al3Mg2) for 5xxx series and (Mg2Si) for 6xxx series. It is well branded that the Al–Mg alloys are susceptible to IGC if the electrochemically active β-phase (Al3Mg2) precipitates at the grain boundaries [11–15]. Moreover, Mg2Si intermetallics reveal anodic performance and exhibit fractional degeneration with discrete alloying owed to elective leaching of magnesium [16,17]. On the other hand, iron-rich intermetallics exhibit extra noble potential relative to aluminum [18,19]. The aim of the present work is to investigate the effect of heat treatment on the microstructure, mechanical and corrosion behavior of different two series (6xxx and 5xxx) of Al alloys.

2Experimental procedure

The AA6061, AA 6013 and AA5086 were supplied by Alcoa Inc. The chemical composition of the studied alloys in (wt. %) is presented in Table 1.

Table 1.

Chemical composition of the studied alloys in (wt.%).

Sample no.  Si  Fe  Cu  Mn  Mg  Cr  Zn  Ti  Al 
AA6061  0.60  0.35  0.30  0.07  0.20  0.10  0.02  Bal. 
AA6013  0.80  0.25  0.85  0.50  0.05  0.12  0.05  Bal. 
AA5086  0.40  0.50  0.10  0.45  0.15  0.25  0.07  Bal. 

Disk like shape samples of 1mm thick and 4mm in diameter were machined for the heat treatment and experimental investigations. Solution heat treatment for all samples was performed in a circulating air furnace at 530°C for 30min, and then quenching was carried out immediately in ice water. Diverse techniques were employed to examine the as-quenched aluminum matrices (ASQ), the natural aged Al alloys at room temperature for long term 1.2×105min (NAT) and the artificial aged Al alloys at 175°C for 30min (AAT). Philips-PW1710 XRD diffractometer with Cu Kα radiation of λ=1.541838Å was used to identify the phase constitution in the Al alloys. The XRD scan was utilized between 30° and 90°, with scan rate of 1°/min and step interval of 0.02°. The surface morphologies of the mentioned Al alloys were performed by Joel-JSM-5400 LV (Japan) scanning electron microscope (SEM). Microhardness tester with load of 100 gmf was laboring to evaluate the Vickers's microhardness. The friction measurements were accomplished in air atmosphere using ball-on-disk tribometer. The corrosion investigations were achieved by Gill AC device using potentiodynamic method in Ringer's solution.

3Results3.1Microstructure characterization

Fig. 1(a–c) indicates SEM micrographs of AA6061, AA6013 and AA5086 aluminum alloys solutionized at 530°C for 30min and quenched in ice water (ASQ). Quenching the metal from elevated temperature (super-saturation) prevented the precipitation of the strengthening phases. Where, the main purpose of the quenching is to obtain an Al matrix supersaturated with solute atoms and vacancies which are retained in solution. It was observed that, there is no formation of precipitates during this condition of heat treatment. Therefore, the Al matrix would contain Mg-Si-Cu in case of 6xxx and Mg-Mn-Zn in case of 5xxx alloys survival in a super saturated solid solution at room temperature [20]. Solution treated alloys undergo natural aging at room temperature (NAT). The formation of small solute clusters was approved by SEM micrographs (Fig. 2a–c) on the beginning of this hardening process, which continues by GP Zones development.

Fig. 1.

SEM micrographs of (a) AA5086, (b) AA6061 and (c) AA6013 alloys solutionized at 530°C for 30min and quenched in Ice water (ASQ).

Fig. 2.

SEM micrographs of (a) AA5086, (b) AA6061 and (c) AA6013 naturally aged treated samples (NAT) at RT immediately after ice quenching.


It was observed that, the natural aging has an adverse effect on artificial aging of the studied alloys. Adding Cu to Al–Mg–Si balanced alloy decreases the solubility of the stable phase Mg2Si, which resulted in increasing the super-saturation of the Mg2Si phase of the given alloy composition [9]. Also, the migration of Mg and Si atoms is reduced by the Cu addition, which retards the formation of Mg–Si clusters at room temperature [13]. A refinement of the microstructure occurs as a results of Cu addition, and consequently the artificial aging kinetics is promoted due to the formation of the Q′ phase in addition to the phases which could be precipitated in alloys with lower Cu content, i.e. β″ and β′ phase. SEM micrographs in Fig. 2(a–c) indicates the formation of finer precipitates from co-clusters and vacancy clusters as a precursor precipitates for the metastable phases β″ and β′ and/or Q′ phase as a result of the natural aging for the studied alloys. Pogatscher et al. [21] reported that, after the dissolution of Mg clusters during the natural aging the co-clusters are formed which is agreed with the presented results.

The coexistence of majority of needle and some road shaped precipitates from β″ and β′ and/or Q′ respectively was observed after artificially aged of the samples at 175°C for 30min (AAT) (Fig. 3a–c).

Fig. 3.

SEM micrographs of (a) AA5086, (b) AA6061 and (c) AA6013 alloys aged at 175°C for 30min and directly quenched in ice water (AAT).


Fig. 4(a–c) clarifies the phase constitutions of the investigated alloys as characterized by XRD. For 6061 and 6013 series, new phase of Mg2Si was detected with preferred orientations of (420), (422) and (511) after natural aging treatment of the alloys at room temperature and artificial aging treatment at 175°C (Fig. 4a, b) which matches with ICDD-card no. 04-017-6811 [22]. Moreover, η(MgZn2) and Al3Mg2 phases were observed in AA5086 alloy after natural and artificial aging (Fig. 4c) which matches with ICDD-card no. 04-017-1426 and ICDD-card no. 00-040-0903 [23,24]. Further, a narrow peak shift toward higher angle was observed for all alloys after the natural and artificial aging. This peak shift indicates the thermally activated nature for these precipitated phases and increase in the critical nucleation size for the most important hardening phase precipitates. Table 2 clarifies the change in the crystallite size of all investigated alloys which are calculated based on Scherrer equation.

Fig. 4.

The phase constitutions of (a) AA6061 alloy, (b) AA6013 alloy and (c) AA5086 alloy as characterized by XRD.

Table 2.

The change in the crystallite size and microhardness for all investigated alloys.

Sample  AA6061AA6013AA5086
Heat treatment  ASQ  NAT  AAT  ASQ  NAT  AAT  ASQ  NAT  AAT 
Average crystallite size (nm)  26.15  42.90  42.90  20.60  50.50  50.20  27.04  53.94  47.70 
Average microhardness (HV100)  71  63  153  74  65  160  70  60  140 

It was observed from the table that the crystallite size increases for all investigated alloys after natural and artificial aging. The presence of narrow peak shift and the increase in the crystallite size are ascribed to the clusters formation caused by natural aging for long term which in turn reduces the initial solute super-saturation as in the case of artificial aging [21].

3.2Mechanical behavior

Table 2 represents the Vickers microhardness measurements for all investigated alloys. It was observed from the table that the microhardenss is decreased after natural aging treatment (NAT) at RT for all alloys. On the other hand the microhardness is increased after age hardening at 175°C for all Al alloys artificial aging (AAT). This behavior is ascribed to lower concentration of vacancies and solute super saturated atoms during the natural aging at RT which have adverse effect on precipitation hardening in the material. During the NAT the formation of co-clusters governs by the vacancy diffusion mechanism [25]. Moreover, during the age hardening process (artificial aging) the particle size of nanostructure precipitates increases which resulting in increasing the microhardness. The density of movable vacancies which result after the quenching process controlling at most in the artificial aging effect on the microhardness [21–29].

Fig. 5(a–c) clarifies the friction coefficient of all the investigated alloys. For AA 6061, the friction coefficient of the artificially aged alloy at 175°C (AAT) is less than the naturally aged alloy at RT (NAT) and both two values are less than the as-Quenched one (ASQ). For AA 6013 one can observe that the friction coefficient is increased after naturally aged (NAT) then it decreased after artificially aging at 175°C (AAT). Moreover, the two values of the friction coefficient are higher than the ASQ one.

Fig. 5.

The friction coefficient of (a) AA6061 alloy, (b) AA6013 alloy and (c) AA5086 alloy.


For AA 5086, the friction coefficient is increased after naturally aged treatment and after that it decreased after artificially aged treatment and recorded value lower than the as-quenched one. From the presented results of friction coefficient, two major outcomes can be obtained: The first one is the friction coefficient of all artificially aged alloys are less than the naturally aged one. The second one, comparing the friction coefficient for all artificially aged alloys, the friction coefficient for AA 5086 is less than the friction coefficient for both AA 6013 and 6061.

It should be mentioned that the chemical composition of the surface is responsible for decreasing the coefficient of friction. In farthest requests of Al-Mg based (5xxx) alloys, the work hardening and solution treatment strengthening clearly affected the surface chemical composition and surface properties. Moreover, small Cu insertions can enhance the mechanical properties owed to the precipitation hardening. 5xxx alloys with small amount of Cu has good formability, good strength and the benefits of precipitation hardening which in turn decrease the friction coefficient [3–5].

3.3Electrochemical behavior

Fig. 6(a–c) displays the potentiodynamic polarization curves of all investigated alloys. For group series AA6061 and 6013 the corrosion potentials shifted to more negative values for the alloys naturally aged at RT in comparison with the as-quenched alloys. Moreover, the corrosion current is high, indicating the decrease in the corrosion resistance. After aging the alloys at 175°C for 30min (AAT), the corrosion potentials shifted to more positive values and a decrease in the corrosion current was observed. This signified to the enhancement in the corrosion resistance. On the other hand, for AA5086, artificial aging of the alloy at 175°C resulted in decreasing the corrosion resistance. This behavior is clarified by high corrosion current and low corrosion potential. Comparing the potentiodynamic curves for all investigated alloys artificially aged at 175°C for 30min (AAT) clarifies that the 6061 and 6013 group series have the best corrosion resistance in comparison with other alloy as indicated in Table 3.

Fig. 6.

The potentiodynamic polarization curves of (a) AA6061 alloy, (b) AA6013 alloy and (c) AA5086 alloy.

Table 3.

corrosion current density (ImA/cm2) and corrosion voltage (Emv) of three studied alloys.

E corr. (mv)  I corr. (mA/cm2Heat treatment  Sample 

The presented impurity of Si in most 6xxx series lead to the formation of Mg2Si particles which are improve the corrosion resistance of these alloys [6]. It was reported, the insertion of Si during selective dissolution of Mg caused the formation of Mg(OH)2 and SiO2·nH2O hydroxide that improve the corrosion resistance by acting as an extra barrier for pitting corrosion [7–11].

To discuss the difference in the properties between the solution treatment (ASQ) and aged treatment of the investigated alloys (AAT and NAT), the precipitation sequence in the studied alloys shouted be presented as follows:

The Precipitation Sequence in the studied alloys AA6061 alloy: SSSS – GP-zones – β″ (Mg2Si) – β′ (Mg2Si) – β(Mg2Si) phase, and for AA6013 alloy: SSSS – GP-zones – β″ (Mg2Si) – β′ (Mg2Si) and/or Q′ – β(Mg2Si)+Q phase. And for AA5086 alloy: SSSS – GP-zones – β″ (Mg2Al3) – β′ (Mg2Al3) – β(Mg2Al3) phase [30–36].

The 530°C is the temperature in which most of Al alloys take the supersaturated solid solution. This means that the dispersed atoms have homogeneous shape with Al matrix. Moreover, the supersaturated solid solution phase has more vacancy atoms. While aging the alloys at 175°C will result in forming β″ (Mg2Si). This phase is considered as a semi-coherent phase which is characterized by low vacancy and the atoms is more bonded to Al matrix. In this regard the hardness, wear and corrosion resistance is decreased after solution treatment of aluminum alloys at 530°C (ASQ) and it increased after aging the alloys at 175°C for 30min (AAT).


The effect of age hardening of three different solution treated aluminum alloys (AA6061, AA6013 and AA5086) was studied. The solution treatment was performed at a temperature of 530°C (ASQ), while the natural age hardening is performed at RT for two weeks and the artificial age hardening is performed at a temperature of 175°C (AAT) for 30min. It was found that, the mechanical and electrochemical properties of all (naturally and artificially) aged alloys are bitter than the solutionized alloys (ASQ). The AA5068 alloy has a wear characteristic better than AA6013 and 6061, while AA6061 and 6013 has a corrosion resistance better than AA5068. In this regard, AA5068 can be used in mechanical applications while AA6013 and 6061 can be used in corrosive media after artificial aging at the optimum temperature 175°C for 30min.

Conflict of interest

The author declares no conflict of interest.

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