Journal Information
Vol. 9. Issue 1.
Pages 1005-1024 (January - February 2020)
Download PDF
More article options
Not available
Vol. 9. Issue 1.
Pages 1005-1024 (January - February 2020)
Original Article
DOI: 10.1016/j.jmrt.2019.11.040
Open Access
Mechanical properties and corrosion behavior of artificially aged Al-Mg-Si alloy
NOTICE Undefined index: totales (includes_ws_v2/modulos/cuerpo/info-item.php[202])
Aluru Praveen Sekhar, Ashim Bikash Mandal, Debdulal Das
Corresponding author
Department of Metallurgy and Materials Engineering, Indian Institute of Engineering Science and Technology, Shibpur, Howrah 711103, India
Article information
Full Text
Download PDF
Figures (15)
Show moreShow less
Tables (3)
Table 1. Summary of various ageing states of the selected Al-Mg-Si alloy.
Table 2. Summary of the mechanical properties of an Al-Mg-Si alloy artificially aged at 175°C for various durations.
Table 3. The dominant modes of corrosion attack at different states of ageing.
Show moreShow less

This article reports the influence of the state of artificial ageing on the hardness, tensile properties and corrosion behavior of an Al-Mg-Si alloy. Pitting corrosion in 0.1M NaOH solution and susceptibility to intergranular corrosion (IGC) have been systematically investigated, covering highly under-aged to highly over-aged states. In alkaline solution, corrosion rate increases with the progress of ageing; while for a given state of ageing, it initially increases with immersion time followed by rapid reduction since corrosion is controlled by the competition of formation and dissolution of Al(OH)3 plus Al2O3 film as determined by FESEM, EDS, and XRD characterizations. The IGC susceptibility of the selected alloy is found to be governed by the microstructure as determined by the state of ageing since it is controlled by the anodic dissolution of the precipitate free zones while closely neighboured grain boundary precipitates act as cathode. It is identified as slight to moderate pitting in highly under-aged, moderate to heavy pitting in under-aged, pitting with localized ICG in peak-aged, localized to uniform IGC in over-aged, and uniform IGC to etching in highly over-aged state. It is adjudged that apart from the inclusions like AlCrMnFeSi, the changes in the precipitation such as β” (Mg5Si6), β' (Mg9Si6) and β (Mg2Si) within the grains and at the grain boundaries with progress of artificial ageing play a pivotal role on mechanical properties and corrosion characteristics of the Al-Mg-Si alloy.

Mechanical properties
Intergranular corrosion
Full Text

Aluminum (Al) alloys of 6xxx series grade are the most extensively used extruded products around the globe [1]. Commercial 6000 series alloys generally consist of Mg, Si, Fe and Cu as the major alloying elements. Grade 6063 Al-alloy is one among them, which is the most economical and is used presently worldwide in a range of applications owing to its ability to form complex shapes with an excellent surface finish. In this alloy, Mg and Si are the main alloying elements, and these contribute to strengthening by intermetallic phases precipitated during proper heat treatments [2]. Although the alloy enhances its mechanical properties through precipitation, it can also have a detrimental effect on the corrosion properties [3]. In the recent past, several researches [4–7] have been directed to understand the role of precipitates on the corrosion characteristics of artificially aged Al-alloys.

Guan et al. [8] have investigated the influence of natural ageing on the pitting corrosion of 6005Al-T6 in 3.5wt.% NaCl solution for over a week to generate a database for the interactive corrosion risk management of high-speed railway lines. These authors have suggested that pitting corrosion occurs as a result of the dissolution of the Al-matrix around the cathodic precipitates, intergranular attack, and channelling. The density of pits has been observed to be high for the coarse grain material in comparison to their finer ones. Kairy et al. [9] have studied the effect of Cu and ageing conditions on the metastable pitting behavior of 6xxx series alloys. Based on potentiodynamic polarization as well as potentiostatic transient tests in 0.1M NaCl solution, it has been inferred that the metastable pitting rate decreases with increase in the Cu-content. Formation of larger precipitates during the process of ageing has caused a reduction in pitting potential, i.e., the resistance to pitting is maximum at the over-aged condition [9]. Prabhu et al. [10] have assessed the corrosion behavior of an Al-Mg-Si alloy in different concentrations of H3PO4 as well as NaOH solutions by electrochemical method. It has been reported that the corrosion rate is considerably higher in NaOH solution in comparison to the H3PO4 solution, and the corrosion rate exhibits an increasing trend with the concentration and temperature of the electrolyte. Several investigators [11–15] have reported that age hardenable Al-alloys are prone to pitting corrosion. De et al. [11] have shown that the extent of pitting corrosion is more significant in the over-aged alloy as compared to peak-aged ones. Gadpale et al. [12] have indicated that corrosion resistance of 2014 Al-alloy degrades with the increase of temperature and time of ageing.

Pitting is frequently encountered as precarious forms of corrosion in Al-alloys. In aggressive environments, these alloys are prone to pitting corrosion with severity ranging from highly uniform to extremely local [16]. Pitting corrosion is driven by the concept of passivity and occurs when a passive film breaks down under the presence of any electrolyte resulting in localized cavity formation. Pits are generally nucleated by the adsorption of aggressive anion on the surfaces having chemical or physical heterogeneities such as inclusions, intermetallic particles, flaws, and mechanical damages [16–18]. Besides pitting, intergranular corrosion (IGC) is also the commonly occurred one in many wrought 6xxx Al-alloys primarily influenced by the chemical compositions [19] and heat treatments [20]. The IGC susceptibility of 6xxx Al-alloys is known to be influenced by the ratio of Mg/Si, Cu content, the depletion of Si and Cu atoms, precipitation of Si and Cu phases at the grain boundaries, and the preferential dissolution of Mg2Si phase [21,22].

Guillaumin and Mankowski [23] have earlier studied the effect of two-step over-ageing (T78) and other intermediate ageing states (T78-1 and T78-2) on the localized corrosion of Al 6056 in 1M NaCl solution. The corrosion mechanisms of the alloy in T78, T78-1 and T78-2 states are found to be similar to that of peak-aged state. The coarse intermetallic particles formed during the over-ageing are found to be the nucleation sites for the corrosion. These authors also advocated that the alloy in over-aged condition exhibits high resistance to IGC than the peak-aged state because of the lower potential difference between the matrix and the grain boundary as a result of larger precipitates that are formed throughout the matrix. Similar kind of ageing treatments is employed recently by Wang et al. [24] to improve the intergranular corrosion resistance of Al-Mg-Si-Cu alloy in comparison to T6 treatment while retaining its strength.

Most of the research on the corrosion behavior of 6xxx Al-alloys [20,22,25–27] is primarily focused on the effects of Cu content [25,28,29] and Mg/Si ratio [3,30,31]. The role of microstructure on corrosion behavior of Cu-free Al-Mg-Si alloy covering a wide range of ageing states has not been well understood. The present study aims to reveal the influence of artificial ageing covering highly under-aged to highly over-aged states on the susceptibility to IGC and corrosion characteristics in an alkaline (0.1M NaOH) media of Al-Mg-Si alloy with negligible Cu. For this purpose, the selected Al-Mg-Si alloy has been subjected to artificial ageing at 175°C for a time duration ranging from one hour to over two weeks. Corroded surfaces of the investigated specimens have been examined via optical microscopy, FESEM, EDS, 3D optical profilometer and XRD techniques to achieve insight on the micromechanisms of corrosion. Measurements of hardness and tensile properties apart from the assessment of precipitation state are integral parts of the present research scheme.

2Experimental procedures2.1Material

A commercial Al-Mg-Si alloy available in the hot extruded condition with 16mm diameter is used for the present examination. An atomic emission spectroscopic analysis is performed to acquire the chemical composition (0.5Mg, 0.435Si, 0.178Fe, 0.001Cu, 0.082Mn, 0.006Cr, 0.067Zn, 0.023Ti and Al in balance) of the selected alloy (all in wt.%). The obtained composition in comparison with the ASM standards [32] confirms that the alloy procured belongs to AA6063 grade.

2.2Heat treatment and microstructural characterizations

The initial objective of this investigation is to understand the ageing response of an Al-Mg-Si alloy based on the evolution of microstructure, measurements of Vickers hardness and tensile properties. Specimens of 10mm length are sectioned for a microstructure as well as for Vickers hardness indentations. For the tensile tests, specimens having gauge dimensions of length, 25mm and diameter, 6mm are prepared via a mini-lathe machine. After few principal assessments, all the specimens are exposed to solution treatments in PID controlled muffle furnace at the selected temperature of 525±2°C for 2h followed by rapid quenching (quench rate, 10°C s−1) in an ice bath to form a solid solution in supersaturated state. These specimens are immediately subjected to age-hardening treatments at a temperature of 175±2°C in a woven furnace for different time intervals starting from 1 to 504h to achieve different ageing states. For the ease of discussion, the various states of ageing are divided into five segments; namely, highly under-aged (HUA), under-aged (UA), peak-aged (PA), over-aged (OA) and highly over-aged (HOA). The limits of ageing time (tA) for these segments are defined here by considering the obtained ageing curve. Hardness value of 65 HV2 is used to demarcate between HUA and UA as well as OA and HOA states. These are summarized in Table 1. Furthermore, it may be mentioned here that a particular specimen is denoted by the state of ageing where the value of tA is also noted as subscript. For example, HUA2 refers to the highly under-aged state obtained by isothermal (TA=175°C) ageing for tA=2h.

Table 1.

Summary of various ageing states of the selected Al-Mg-Si alloy.

Abbreviations  Ageing State  Ageing Time (h) 
AQ  As Quenched 
HUA  Highly Under-Aged  > 0 to ≤ 2 
UA  Under-Aged  > 2 to ≤ 8 
PA  Peak-Aged 
OA  Over-Aged  > 8 to ≤ 250 
HOA  Highly Over-Aged  > 250 to ≤ 504 

After the heat treatments, specimens are made flat and polished using different grades of emery papers (grit sizes: 220, 400, 800, 1200) and diamond pastes (6, 3, 1 and 0.25μm) to achieve a fine scratch-free surface. Whereas, for tensile specimens, the gauge portion is polished in a similar method before performing any test. Microstructural analysis in this study is primarily carried out to reveal the grain morphology of the alloy under different ageing states. The polished specimens are etched using a modified Keller’s reagent by changing the concentrations as required for differently aged samples. Optical micrographs of the etched specimens are recorded at different locations in various magnifications using the optical microscope (Carl Zeiss). The grain sizes of the microstructures are measured using Image J software following a linear intercept method.

2.3Mechanical testing

Vickers hardness indentations are carried out at random locations using the micro-hardness tester (Leica, Germany) under 2 Kgf load for indentation time of 15s. At least ten readings are taken on each specimen to report the average value with standard error. The tensile tests are performed uniaxially at room temperature (25°C) using Instron 8862 universal testing machine controlled via Bluehill® software. The crosshead speed of 1.92mm/s is selected to perform the tensile test, which corresponds to a strain rate of 10−3 s−1. At least two tests are performed for each ageing condition, and their average values are reported.

2.4Corrosion tests

To investigate the influence of the state of ageing on the corrosion resistance of an Al-Mg-Si alloy, specimens of 15×15×4mm3 dimensions for immersion testing, as well as 30×15×4mm3 for IGC testing are prepared from the stock material. Heat treatments are carried out on the prepared specimens at selected temperatures and durations. Afterwards, the surface of specimens is finely polished following standard metallographic techniques. Before corrosion tests, specimens are thoroughly cleaned twice, both in acetone as well as in ethanol using an ultrasonic cleaner to remove any impurities that are caught during their preparation.

Analytical grade sodium hydroxide (NaOH) pellets (low chloride) are procured for the immersion corrosion tests. Solution bath of 500ml volume is prepared with 0.1 molar concentrations, which correspond to 2g of NaOH and double-distilled water. Corrosion tests are performed at room temperature by dipping the entire specimen in solution which is tied to a nylon thread for different immersion times (ti) of 1, 3, 6, 12, 24, 72 and 168h. After performing the tests for specific immersion times, specimens are retrieved from the solution, and then the surface is gently wiped with a soft tissue. Cleaning process of the corroded specimens is initiated by washing them in 70 % HNO3 solution for 2min and then in distilled water to remove the corrosion products while retaining the topography of the corroded surface [33]. The procedure mentioned above follows the standard practice ISO 8407:2009 [34]. Alloys resistance to immersion corrosion at different ageing conditions is assessed from the weight loss data obtained by measuring the weight of specimen before and after the corrosion tests using a high precision (10−4 mg) microbalance (ABS 220-4, KERN & Sohn, Gmbh, Germany). A minimum of two tests have been performed for a particular ageing condition for a given ti, and their average values are reported. Corrosion rate has been determined based on the weight loss of specimens following [35]:

where the corrosion rate is expressed in mils per year, WL is the weight loss in mg, ρ is the density of specimen which is 2.698g/cm3 for the selected Al-Mg-Si alloy, A is the exposed surface area (here, 6.78cm2) and ti is immersion time (h). Further, the topography of the corroded surfaces is characterized using 3D optical profilometer (Contour GT-K, Germany) to measure the changes in the average surface roughness (Ra) that is occurred as a result of pitting.

The IGC testing is carried out as per the British standard (BS-ISO 11846) method B [36]. This process comprises of sequence of steps starting with the etching of the specimens in 10wt.% NaOH solution for 5min at room temperature followed by desmutting in 30 % HNO3 for 2min before immersing them for 24h in 30g NaCl +10ml conc. HCL+distilled water solution. Precisely after the specified time of 24h, the samples are withdrawn from the solution and cleaned them in distilled water and ethanol before drying in cold air. IGC susceptibility of the tested alloy at different ageing states is evaluated from the measurements of the average and the maximum corrosion depths. For this purpose, corroded specimens are sectioned in a transverse direction having a width of 10mm using an abrasive cutter. Two of these sectioned pieces are cold mounted and mirror polished for measuring the corrosion depth under an optical microscope, and the middle portion is retained for morphological characterizations.

Field emission scanning electron microscopy (FESEM) equipped with energy-dispersive X-ray spectroscopy (EDS) is used to characterize the morphology of some corroded surfaces tested in 0.1M NaOH solution as well after the IGC testing. After scrutinizing the recorded micrographs, the prevailing mode of corrosion is noted for each ageing condition. Characteristics of the corroded surfaces (with and without removing the corrosion products) are further evaluated by the X-ray diffraction technique (Model: Bruker D8 Advanced, UK) using Cu-Kα diffractometer (λ=1.5418Å) with an operating voltage of 60 KV and operating current of 30mA. The data is collected within the scan range of 25°-100° at a scan rate of 0.02°/s; PAN Analytical X-pert High Score software is used to analyze the corrosion products and phases present in the investigated alloy.

3Results and discussion3.1Ageing characteristics

The artificial ageing response of the selected Al-Mg-Si alloy is evaluated based on the evolution of microstructure and the measurements of Vickers hardness values. The microstructures of different ageing states (HUA2, UA4, PA8, OA72 and HOA336) are shown in Fig. 1. Etching of the chosen specimens has revealed a fine equiaxed grains with their grain sizes ranging from about 35–75μm. The grain size of HOA336 specimen is measured to be slightly higher in comparison to those specimens of other ageing states. The average grain size is measured as 52±5μm. It is a well-known fact that the age hardening phenomena is a result of the formation of metastable precipitates. Fig. 2(a) illustrates the Vickers hardness values of Al-Mg-Si alloy against different ageing times from 1 to 504h heat-treated at a temperature of 175°C. Results in Fig. 2(a) of the selected alloy show the typical response of any age hardenable alloy [37]. Vickers hardness value of the as-quenched (AQ) sample after solutionizing at near eutectic temperature is measured as 43.7±1.11 HV2. Artificial ageing treatment of the alloy exhibits a higher level of hardness values in comparison with the as-quenched sample. After 1h of ageing, the hardness value of the sample increases to 46.65±0.25 from 43.7±1.11 HV2. The marginal increase in hardness is attributed to the quenched-in vacancy assisted formation of solute (Mg and Si) clusters [38,39]. With increase of tA, hardness value increases gradually and reaches a maximum value of 89.46±0.25 HV2 at 8h. The increment in hardness level is due to the restraining of the dislocation movement by GP zones [40,41], and semi-coherent β” (Mg5Si6) needle-shape nanometer-size ranged precipitates [42] which precipitates with the progress of ageing. Further continuation of ageing treatment has shown a drop in the level of hardness values from 89.46±0.25 HV2 at 8h to 52.6±0.59 HV2 after 504h of ageing (Fig. 2(a)). The observed trend in hardness is an outcome of the coarsening of precipitates like β' (Mg9Si5) [43], U1 (MgAl2Si2) [44] and U2 (Mg4Al4Si4) [45] apart from precipitation of equilibrium β (Mg2Si) and its coarsening [46,47]. The ageing response of 6063 Al-alloy at same ageing temperature for different ageing times are studied earlier by He et al. [48] and obtained a similar age-hardening response although the reported peak hardness value is 83 HV2 achieved at 12h of ageing.

Fig. 1.

Representative optical microstructures of Al-Mg-Si alloy specimens artificially aged at fixed temperature (TA=175°C) for the durations (tA) of (a) 2h (highly under-aged, HUA), (b) 4h (under-aged, UA), (c) 8h (peak-aged, PA), (d) 72h (over-aged, OA) and (e) 336h (highly over-aged, HOA).

Fig. 2.

Mechanical properties of Al-Mg-Si alloy artificially aged at TA=175°C against different times of ageing are presented. (a) Vickers hardness, (b) Engineering stress-strain curves of some selected specimens, (c) yield strength and ultimate tensile strength, and (d) uniform and total elongation properties.

3.2Tensile properties

Engineering stress-strain curves of some selected specimens are depicted in Fig. 2(b) to illustrate the effects of different state of ageing on tensile behavior of the selected alloy. It is evident from Fig. 2(a) that tensile response varies widely with the tA that determines the state of ageing. The tensile properties such as yield strength (YS), ultimate tensile strength (UTS), uniform elongation (UE) and total elongation (TE) of the selected alloy aged for various times are presented in Figs. 2(c and d); the dashed lines in these figures denote the corresponding strength and elongation values of AQ specimen. Artificial ageing results in significant variation in the strength and elongation as compared to AQ sample. One can note that with an increase in time of ageing from HUA1 to PA8 states the YS and UTS of specimens gradually increase from 156±4.13 to 264.3±4.26MPa and 191.7±5.06 to 285.7±6.3MPa which corresponds to the percentage of increase by 69.5 % and 49 %, respectively (Table 2). In contrast, the UE and TE values decrease from 14.47±0.1 to 7.2±0.05 % and 26.32±0.8 to 17.62±0.65 % corresponding to the percentage of decrease by 50 % and 33 %, respectively. The alloy achieves a maximum YS and UTS values at PA condition with the lower TE values (Fig. 2(c and d)). This observation can be attributed to the hindrance of dislocation movements within the matrix by the formation of transient precipitates in a sequential manner through diffusion phenomena [38]. As the ageing is progressed further, i.e. from PA8 to HOA504 states, the YS and UTS values reduce from 264.3±4.26 to 103.3±7.41MPa and 285.7±6.3 to 155.5MPa, respectively; while, the TE value increase from 17.62±0.65 to 18.7±0.75 % (Table 2); as expected for any typical age hardenable alloy. The declination of hardness from PA8 to HOA504 can be attributed to the dynamic change in the strengthening phenomena from shearing of precipitates for PA state to the formation of Orowan loops around the equilibrium precipitates in HOA states [49].

Table 2.

Summary of the mechanical properties of an Al-Mg-Si alloy artificially aged at 175°C for various durations.

Ageing time, tA (h)  VHN  YS (MPa)  UTS (MPa)  UE (%)  TE (%) 
43.17±1.11  94.1±2.5  183.3±3.4  17.53±0.61  45.51±1.05 
46.65±0.25  155.9±3.8  191.7±3.5  14.47±0.61  26.32±0.99 
1.5  51.96±0.25  151.5±2.8  202.9±4.3  11.94±0.58  22.52±0.79 
61.09±0.56  172.2±4.5  226.9±5.1  10.22±0.48  25.63±0.68 
2.5  68.20±0.41  181.3±4.8  220.9±5.3  9.00±0.36  19.08±0.77 
73.69±0.36  176.3±5.4  234.2±5.5  8.73±0.35  19.54±0.57 
81.26±0.55  227.3±5.3  266.0±5.7  8.38±0.33  18.35±0.59 
86.40±0.63  246.5±6.8  278.1±6.6  8.41±0.34  18.36±0.55 
89.46±0.25  264.3±7.4  285.7±6.9  7.20±0.29  17.62±0.55 
12  86.06±0.27  243.6±5.1  259.6±7.1  6.43±0.26  19.57±0.59 
24  84.26±0.24  216.7±6.7  245.2±6.2  6.60±0.26  18.96±0.57 
48  82.91±0.50  227.3±6.0  249.3±7.0  6.18±0.25  18.09±0.54 
72  81.20±0.40  216.5±6.1  241.6±7.1  6.23±0.25  18.92±0.47 
96  77.91±0.73  207.5±6.8  238.5±6.2  6.29±0.24  19.24±0.58 
120  72.98±0.18  192.3±6.2  229.5±6.0  6.16±0.22  19.54±0.79 
168  68.42±0.13  174.4±5.2  222.6±5.6  6.10±0.28  20.49±0.61 
336  62.97±0.09  137.6±4.1  193.5±4.8  5.51±0.26  19.74±0.59 
504  52.60±0.59  103.3±2.8  155.5±3.8  5.51±0.31  18.70±0.90 

VHN: Vickers hardness; YS: Yield strength; UTS; Ultimate tensile strength; UE: Uniform elongation; TE: Total elongation.

3.3Fracture mechanisms

To have further insight into the effects of ageing parameters on the tensile behavior of Al-Mg-Si alloy, some selected tensile-fractured surfaces have been examined under SEM coupled with EDS facility. The representative SEM micrographs of the tensile fracture surfaces of HUA2 and PA8 specimens are shown in Fig. 3. Macroscopic views of the fractured surfaces of both HUA2 (Fig. 3(a)) and PA8 (Fig. 3(b)) specimens show typical cup-and-cone fracture indicating the mode of failure as ductile. The degree of deformation before fracture is higher for HUA2 alloy as compared to PA8 alloy indicating higher ductility of former; this corroborates well with the obtained results (Table 2).

Fig. 3.

Representative fracture surface morphologies of (a and c) HUA2 and (b, d and f) PA8 Al-Mg-Si alloy specimens. (e) EDS analysis corresponds to the elongated intermetallic particle as marked by square box in (d).


Microscopic views of the fracture surfaces corresponding to the tensile and shear failure zones of HUA2 and PA8 specimens are presented in Figs. 3(c, d and f). The fractured specimens reveal typical transgranular dimples in the central tensile failure zone. Both the specimens show equiaxed dimples of different sizes; much finer dimples are revealed in HUA2 alloy (Fig. 3(c)) indicating higher ductility in comparison with the PA8 specimen (Fig. 3(d)). The formation of dimples under tensile loading is a consequence of the damage accumulation caused by the initiation, growth and coalescence of micro-voids [50,51]. Locations pointed out with red arrows are the ones where the coalescence of closely spaced voids occurs as marked by rectangular box in Fig. 3(c). The initiation of micro-voids generally occurs at the precipitate-matrix and/or inclusion-matrix interfaces [50]. The presence of such particles is evidenced in Fig. 3(d), which is further confirmed by EDS analyses (Fig. 3(e)). Elongated dimples close to the specimen ridges (Fig. 3(f)) are indicative of shear fracture at the final stage of failure.

3.4Corrosion response in alkaline solution

Five widely different states of ageing achieved via isothermal treatment have been selected to study the influence of immersion time (ti) on corrosion behavior in alkaline (0.1M NaOH) solution. Ageing times (tA) have been selected, taking into account the ageing curve of Al-Mg-Si alloy as obtained at a temperature of 175°C (seeFig. 2(a)). The selected tA values are 2h, 4h, 8h, 72h and 336h that develop highly under-aged (HUA), under-aged (UA), peak-aged (PA), over-aged (OA) and highly over-aged (HOA) states, respectively (Table 1).

Fig. 4 illustrates the variations of the corrosion loss and corrosion rate with immersion time (ti) for differently aged Al-Mg-Si alloy. The obtained results signify that corrosion loss increases with increasing ti, and the difference of corrosion loss amongst the various states of ageing becomes more evident only at longer ti (Fig. 4(a)). However, the dissimilarity of corrosion behavior with the state of ageing is obvious in the estimated corrosion rate curve (Fig. 4(b)). At any given ti, the corrosion rate increases in the order of HUA2, UA4, PA8, OA72 and HOA336. For any particular state of ageing, the corrosion rate increases rapidly with ti before reaching a maximum or stable corrosion rate; this is followed by fast reduction of corrosion rate at longer (≥ 24h) ti (Fig. 4(b)). The time to reach the maximum or stable corrosion rate is found to be higher for specimens having relatively lower initial corrosion rate; i.e., longer ti (>12h) for HUA2 and shorter ti (< 3h) for HOA336 alloy (Fig. 4(b)). For a similar alloy and under the identical (0.1M NaOH) electrolyte, Kisasoiz [52] has reported that the corrosion rate is higher for artificially aged alloy as compared to the naturally aged one. It has been further emphasized that the corrosion rate is initially higher but it reduces substantially with increase in the duration of exposure in alkaline solution [52], which corroborates with the obtained results (Fig. 4). The reduction in the corrosion rate of an alloy is due to the generation of the barrier layer at the solution-metal interface, which decelerates the dissolution process [53,54].

Fig. 4.

The variations of (a) corrosion loss and (b) corrosion rate with immersion time (ti) of diffrently aged Al-Mg-Si alloy specimens in 0.1M NaOH solution. Various states of ageing are achieved via isothermal (TA=175°C) ageing: highly under-aged (HUA2), under-aged (UA4), peak-aged (PA8), over-aged (OA72) and highly over-aged (HOA336).


The topographies of the corroded surfaces have been examined with the help of 3D optical profilometer, and their average surface roughness (Ra) values have been measured. The effects of immersion time (ti) for a selected state of ageing (here, PA alloy, i.e., tA=8h), and the ageing time (tA) at a particular ti (=24h) on the estimated Ra values are shown in Figs. 5(a) and 5(b), respectively. These results are further illustrated using 2D and 3D views of some selected surface topographies as presented in Figs. 6 and 7. In these figures, typical surface height plots are also depicted to highlight the characteristics of the pits generated during immersion in the alkaline solution. It can be seen from the results in Fig. 5(a) that Ra value increases monotonically with ti. It is evident from Fig. 6(a) that the specimen immersed for 1h has only few pits, and most of the surface is unreacted; the corresponding Ra value is measured as 0.54μm. The width and depth of pits are increased with ti (Fig. 6), and the pits are quite large at longer ti (cf Fig. 6(c)). The Ra value is found to increase from 0.54 to 5.72μm with an increase of ti from 1 to 72h. The increase in the dimension of the pits (Fig. 6(d)) with ti has been in good agreement with the rise in the corrosion loss versus ti (Fig. 5(a)).

Fig. 5.

Illustrating (a) the influence of immersion time (ti) on the peak-aged (PA8) Al-Mg-Si alloy, and (b) the effects of state of ageing on the variation of the average surface roughness (Ra) values after the immersion in 0.1M NaOH solution for fixed (ti=24h) duration. Various states of ageing are achieved via isothermal (TA=175°C) ageing: highly under-aged (HUA2), under-aged (UA4), peak-aged (PA8), over-aged (OA72) and highly over-aged (HOA336).

Fig. 6.

The representative 3D (right column) and corresponding 2D (left column) optical profilometer images of the corroded surfaces along with the average surface roughness (Ra) values of peaked-aged (PA8) alloy specimens immersed in 0.1M NaOH solution for duration (ti) of (a) 1, (b) 12 and (c) 72h. (d) surface height versus distance along the lines marked on the 3D images illustrating the variations in the depth and width of the generated pits with ti.

Fig. 7.

The representative 3D (right column) and corresponding 2D (left column) optical profilometer images of corroded surfaces along with the average surface roughness (Ra) values for differently aged Al-Mg-Si alloy specimens after immersion in 0.1M NaOH solution for fixed duration (ti) of 24h. (d) shows the surface height versus distance along the lines marked on the 3D images illustrating the variations in the depth and width of the generated pits with the state of ageing. Various states of ageing are achieved via isothermal (TA=175°C) ageing: highly under-aged (HUA2), peak-aged (PA8), and highly over-aged (HOA336).


Fig. 5(b) compares the roughness values of differently aged specimens immersed for the same duration (ti=24h) in the alkaline solution. The Ra value is found to increase with tA (i.e., in the order of HUA2, UA4, PA8, OA72 and HOA336 states) due to the higher surface degradation by pitting corrosion as evidenced from images shown in Fig. 7. It can also be seen that the Ra value of HUA2 alloy is around 2.13μm, and the same is increased to 4.27μm for an HOA336 alloy. For same immersion condition, typical values of width and depth of largest pit in HUA2 alloy are 9μm and 0.1mm, respectively; these are increased to 32μm and 0.4mm, respectively, for an HOA336 alloy (Fig. 7(d)). These observations are in excellent agreement with the recorded variation of corrosion rate versus tA (Fig. 4(b)).

Morphologies of the corroded surfaces have been investigated with the help of FESEM coupled with EDS microanalyses. Fig. 8 compares the features of the corroded surfaces of HUA2, PA8 and HOA336 specimens immersed for a fixed duration (ti=24h). These specimens have been examined immediately after the completion of immersion tests, i.e., without soaking the specimens in 70 % HNO3 solution to remove the corrosion products which is followed before measurement of weight as per the recommendation of the ISO 8407:2009 standard [34]. The acid treatment has not been carried out purposefully to examine the features of the generated barrier layer. Whereas, another set of FESEM images in Fig. 9 represent the morphology of the surface beneath the barrier layer; i.e., these specimens have been examined after the standard acid treatment to remove the corrosion product/barrier layer. It may also be noted that images in Fig. 9 correspond to the ti of 1 and 168h for a particular state of ageing (here, PA) and hence, illustrate the effect of exposure duration on the surface morphology.

Fig. 8.

Typical FESEM micrographs of corroded surfaces before acid cleaning (i.e., without removal of corrosion products) of (a) HUA2, (b) PA8, and (c) HOA336 Al-Mg-Si alloy specimens after immersion in 0.1M NaOH solution for fixed duration (ti) of 24h.

Fig. 9.

Typical FESEM micrographs of corroded surfaces after removal of corrosion products by acid cleaning of PA8 alloy immersed in 0.1M NaOH solution for duration (ti) of (a) 1 and (b) 168h.


Micrographs in Figs. 8 and 9 indicate that the localized pitting as the prevailing mode of corrosion in alkaline solution. It is obvious from the images in Fig. 8 that the density of the pitting enhances in the order of HUA2, PA8 and HOA336 alloys. The corroded surface of HUA2 specimen reveals cavity-like regions with the occasional presence of corrosion products (Fig. 8(a)). The corroded surfaces of PA8 (Fig. 8(b)) and HOA336 (Fig. 8(c)) alloys exhibit the development of relatively compacted film; the thickness of the film in the latter is, however, more as compared to the former one. Considering the fact that these features have been developed at a fixed ti, the increase in the corrosion products in the order of HUA2, PA8 and HOA336 reflect the variation of corrosion rate as shown in Fig. 4(b). The PA specimen corroded for 1h show only few regions of the surface have been attacked, and most of it is still nearly unaffected (Fig. 9(a)). In contrast, the specimen immersed for 168h exhibits near-uniform pitting corrosion over the entire surface (Fig. 9(b)). The uniform pitting is expected since the dissolution of Al-alloy in a strong alkaline solution is continuous and, the corrosion rapidly propagates over the entire surface at longer exposure times [52,53]. Comparison of images in Fig. 9 further reveals that the size of the pits and their density rise with increasing ti indicating a greater magnitude of corrosion loss as observed in Fig. 4(a) for a particular state of ageing. The competitive film formation and its dissolution process regulate the variations in the corrosion rate of the Al-alloys [54].

Pitting corrosion occurs due to the localized breakdown of the passive film resulting in the rapid dissolution of the metal [55]. The pits commonly occur at the physical or chemical heterogeneity of the surface and tend to be more reactive initially at the most vulnerable sites [56,57]. In the case of age-hardenable Al-alloys, the common sites of chemical heterogeneity are precipitates that eventually control the rate of pitting corrosion [58]. The strong influences of the size of precipitates and their number density on the corrosion behavior have been reported earlier for 2xxx [59], 6xxx [9], and 7xxx [60,61] series age-hardenable alloys. For instance, El-Menshawy et al. [62] have shown that with rise in tA, the volume fraction of the Cu-containing Q (Al4Mg8Si7Cu2) phase and β” (Mg2Si) phases increase in 6061 alloy which, in turn, amplify the cathodic reaction rate leading to the rapid dissolution of the matrix in NaCl solution. Similarly, the evolution of higher amounts of precipitates such as β” (Mg5Si6) and β (Mg2Si) with increasing tA is found to cause greater pitting corrosion leading to higher corrosion rate in the order of HUA2, UA4, PA8, OA72 and HOA336 alloys (Fig. 4). Identification of Fe-rich inclusion at the centre of a pit (Fig. 10) also infers inclusion associated initiation of pitting corrosion [58].

Fig. 10.

(a and b) FESEM micrographs of corroded surface of HOA336 alloy immersed in 0.1M NaOH solution for ti=24h, (c to g) EDS maps illustrating the distribution of various elements, and (h) EDS profile with semi-quantitative results corresponds to the area marked in (b). Note the presence of Fe-rich intermetallic which is found to be depleted in Mg but enriched in Si.


XRD studies of some selected corroded samples, both with and without acid cleaning, have been performed to identify the nature of the corrosion products. Fig. 11 depicts the XRD line profiles of HUA2, PA8 and HOA336 samples immersed in 0.1M NaOH solution for ti=24h. Analyses of XRD patterns helps to identify the presence of Al (JCPDS 00-001-1180) matrix, Mg2Si (JCPDS 00-035-0773) precipitate and Fe-rich intermetallics such as Al0.7Fe3Si0.3 (JCPDS 00-045-1204) in addition to the corrosion products like Al(OH)3 (JCPDS 00-038-0376), Al2O3 (JCPDS 01-075-0783), Fe3O4 (JCPDS 00-026-1136) and SiO2 (JCPDS 01-082-1646) oxides. Formation of these corrosion products has been reported earlier by Soliman et al. [63] on similar kind of alloy immersed in NaOH solution. The obtained results allow to conclude the formation of Al(OH)3 on the surface of differently aged Al-Mg-Si alloy specimens during immersion in the 0.1M NaOH solution. Generation of such film decreases the contact between the unreacted material and the electrolyte resulting in the rapid reduction of the corrosion rate at longer ti (Fig. 4(b)) when the reaction with the virgin surface is nearly completed.

Fig. 11.

XRD line profiles of acid cleaned corroded HUA2 and PA8 specimens, as well as both with and without acid, cleaned corroded HOA specimens. Corrosion has been carried out in 0.1M NaOH solution for 24h. AC: Acid cleaned; WOAC: without acid cleaned.


In alkaline solution, the corrosion involves the electrochemical dissolution of Al [53,64], via both anodic and cathodic processes occurring simultaneously on the metal-electrolyte interface [63,65]. To understand the underlying mechanism of the corrosion, it is important to explore which partial anodic and partial cathodic reactions are involved, and which of them prevails in the gross corrosion reaction. Moreover, the characteristics of the generated film of the corrosion products should also be considered. In the presence of oxide film at the metal-electrolyte interface, the anodic Al dissolution proceeds by direct metal dissolution reaction following the movement of Al ions through the film, as well as, by indirect metal dissolution reaction through repeated formation and dissolution of oxide film [10,66,67]. The corrosion reaction can be expressed as:

In summary, it can be inferred that the corrosion reaction of the Al-Mg-Si alloy in the selected alkaline (0.1M NaOH) solution on the whole proceeds via a partial anodic reaction comprising of two sub-steps - the electrochemical formation of Al(OH)3 film and its chemical dissolution, and a partial cathodic reaction associated with water reduction reaction.

3.5Intergranular corrosion behavior

Most of the reports in the open literature [26,29,30,68] related to the IGC susceptibility of 6xxx alloys are primarily limited to Cu containing Al-Mg-Si alloys. Understanding the IGC susceptibility of Al-Mg-Si alloys having extremely low or free Cu is yet to be crystallized. The IGC susceptibility has been evaluated based on the measurements of average and the maximum corrosion depths. The plots of average and the maximum corrosion depths measured using the optical and/or FESEM micrographs of the cross-sectional views of the IGC specimens is shown in Fig. 12. The markings and digits embedded on the images in Fig. 12 are close to the obtained values of the corresponding maximum corrosion depth. In addition, the prevailing modes of corrosion for each ageing condition are identified via scrutinizing the recorded FESEM micrographs, and these are summarized in Table 3. Representative images of some corroded surfaces are depicted in Fig. 13. It is evident from the cross sectional images in Fig. 12 that the values of both average and maximum corrosion depths increase with the progress of artificial ageing; i.e., the susceptibility to IGC rises with increase in tA. Besides, the images in Figs. 12 and13 indicate that the mode of corrosion changes with progress of artificial ageing. The investigated alloy is found to be prone to localized pitting in the under-ageing regime, and it changes to pitting plus IGC at around peak-ageing condition. Mode of corrosion is primarily IGC in the over-ageing regime. These observations are attributed to the alterations in the characteristics of the precipitation (i.e., size, shape, amount and distribution) with the state of artificial ageing [62]. In general, the pitting after its initiation is known to grow into various shapes which are classified into isotropic and anisotropic group [69]. In the present study, various shapes and sizes of pits are observed (Fig. 12). For instance, shallow elliptical surface and subsurface pits are observed for tA=1h specimen; whereas, narrow but the deep pit is found in tA (=3h) specimen. Considering the morphology of pitting and the IGC, it can be adjudged that the characteristics of intermetallic particles play a vital role in determining the mode of corrosion since they act either as a cathode or anode resulting in their dissolution or that of the Al-matrix [8].

Fig. 12.

The estimated average and the maximum corrosion depths of IGC tested Al-Mg-Si alloy samples artificially aged for varying time (tA) at TA=175°C. Some selected cross sectional optical images illustrate the variation of IGC susceptibility.

Table 3.

The dominant modes of corrosion attack at different states of ageing.

Ageing time (h)  Ageing state  Dominant mode of corrosion 
Highly under-aged  Slight pitting 
Highly under-aged  Moderate pitting 
Under-aged  Moderate pitting 
Under-aged  Heavy pitting 
Peak-aged  Pitting and localized IGC 
12  Over-aged  Pitting and localized IGC 
24  Over-aged  Localized IGC 
96  Over-aged  Uniform IGC 
168  Over-aged  Uniform IGC 
336  Highly over-aged  Etching 
Fig. 13.

Representative FESEM micrographs of intergranular corrosion tested specimens illustrating the influence of different states of ageing on the predominant modes of corrosion. Here, various states of ageing are obtained by isothermal ageing at fixed temperature (TA=175°C) for the durations (tA) of 1 (highly under-aged, HUA1), 3 (under-aged, UA3), 8 (peak-aged, PA8), 24 (over-aged, OA24), 168 (over-aged, OA168) and 336h (highly over-aged, HOA336).


One can notice from a set of FESEM images in Fig. 13 that the modes of corrosion vary with isothermal ageing at TA=175°C. At the early stage of ageing (i.e. HUA), investigated alloy experiences slight pitting corrosion (Fig. 13(a)) which is triggered by the presence of either surface defects or a few numbers of solute clusters [26]. The increase of atomic clusters at higher tA results in moderate pitting corrosion as evidenced by the development of much deeper and denser pits in the UA3 alloy (Fig. 13(b)). The dominant modes of corrosion attack in IGC specimens of PA alloy are found to be pitting and localized IGC (Figs. 13(c)). It may be noted that the pits are much smaller in size in the PA sample as compared to the UA specimen (Fig. 13(b) vis-a-vis Fig. 13(c)). The localized ICG is the predominant mode of IGC in the OA (tA=24h) alloy although few pits are also observed (Fig. 13(d)). With increase in tA in the OA regime, localized IGC is transformed into uniform IGC (Fig. 13(e)). In contrast, HOA (tA=336h) alloy exhibits etching of the surface (Fig. 13(f)) having characteristics of shallow and dense pits closely packed over the entire surface [70,71]. The UA alloy is resistive to IGC and exhibits pitting corrosion owing to the homogenous distribution of solute atoms throughout the matrix, i.e., at both grain boundaries and grains [72]. The formation of precipitates such as βꞌꞌ (Mg5Si6) at the grain boundaries and the depletion of solute atoms specifically Si around these introduce the localized IGC in PA alloy [20]. The presence of discontinuous β (Mg2Si) precipitates of larger size and high volume fractions are primarily causing the severe grain boundary corrosion via crystallographic tunnelling in an OA alloy [31,71].

The IGC occurs due to severe selective dissolution at the grain boundaries or in the vicinities of the grain boundaries without an appreciable attack on the grain interiors. Susceptibility to IGC in Al-alloys is controlled by the segregation of alloying elements or preferential precipitation of second-phase particles rich in alloying elements at the grain boundaries causing the development of solute or precipitate free zones (PFZ) adjacent to the grain boundaries. These regions are prone to corrosion attack while the grains are cathodically protected to some extent [73]. In general, 6xxx Al-alloys show minor susceptibility to IGC, particularly if the Mg/Si ratio is balanced as required for the formation of Mg2Si [74]. However, an excess amount of Si is often added in the Al-Mg-Si alloys to improve mechanical property [68,75] as the case in the chosen alloy in which the Mg/Si is 1.16. The excess Si increases the susceptibility to IGC since it preferentially segregates at the grain boundaries making these more cathodic compared to Si depleted adjacent regions which are anodic in nature [76–79].

The observed alterations in the corrosion behavior of Al-Mg-Si alloy with the state of ageing can be explained by considering the changes in the microstructural characteristics such as the evolution and distribution of precipitates within the grain as well as at the grain boundaries. Fig. 14 shows the schematic representation of the surface morphology of differently aged alloy subjected to IGC testing. The schematic representations of the microstructure of UA, PA and OA alloys are based on the concerned precipitation states reported elsewhere [80]. Except for a few solute clusters or GP zones, the matrix and grain boundary of UA specimen is free from precipitates (Fig. 14(a)). The lack of precipitation does not provide any favourable conditions for the occurrence of IGC. In PA alloy, the localized IGC has occurred as a result of the coupling between the discontinuous grain boundary precipitation and matrix involving cathodic like AlFeSi inclusions or anodic such as β” (Mg5Si6) precipitate and the precipitate free zone (PFZ) [20,68]. In contrast, almost continuous grain boundary precipitation of equilibrium (Mg2Si) phase and the development adjacent PFZ zones has resulted uniform IGC in case of OA alloy. In case of Cu-containing Al-Mg-Si alloys, it has been reported that the evolution of Q-phase (Al4Cu2Mg8Si7) as the grain boundary precipitates increases the susceptibility to IGC since the precipitate forms the micro-galvanic coupling with the less noble PFZs [25,68,81]. In addition, intermetallic dispersoids generally found in the grain boundary provide additional driving force for IGC and stimulate it’s propagation [72,81] since these are cathodic with regard to the matrix [82]. The presence of an intermetallic particle (AlCrMnFeSi) at the grain boundary region has been confirmed by EDS analyses (Fig. 15).

Fig. 14.

Schematic representation illustrating the role of microstructure as defined by the states of ageing on the intergranular corrosion mechanisms. UA: Under-aged; PA: Peak-aged; OA: Over-aged.

Fig. 15.

EDS mapping and spectra of an intermetallic particle present in the grain boundary of typical corroded IGC tested specimen. Note that intermetallic is rich in Fe, Cr, Mn and Si.


In summary, it can be stated that the mechanisms involved in the intergranular attack of the investigated Al-Mg-Si alloy with its progress of ageing from HUA to HOA are: pit initiation → pit propagation → pitting+IGC → IGC → IGC propagation → etching.


Artificially aged 6xxx Al-Mg-Si alloys are employed in a spectrum of applications because they are cost effective and provide wide range of properties. The present study is directed to investigate the influence of the state of ageing on the mechanical properties and corrosion behavior of an Al-Mg-Si alloy. Artificial ageing has been carried out at the temperature of 175°C for various time intervals ranging from 1h to 336h, and the corrosion performance is evaluated via immersion in 0.1M NaOH solution and the intergranular tests. The major conclusions are:

  • The selected Al-Mg-Si alloy with little amount of Mg (0.5wt.%) and Si (0.43wt.%) exhibits strong age hardening response. Yield strength of the selected alloy increases from 94±5.32MPa in as-quenched state to 264.3±4.26MPa in peak-aged condition. The variations in the hardness and tensile properties with ageing time are discussed.

  • Irrespective of the state of ageing, the corrosion rate of the Al-Mg-Si alloy in the alkaline (0.1M NaOH) solution increases rapidly with immersion time (ti) before reaching either a maximum or a stable value; which is followed by a rapid reduction of corrosion rate at higher (≥24h) ti due to the formation of aluminium hydroxide film. The corrosion process of the Al-Mg-Si alloy proceeds by means of a partial anodic as well as a partial cathodic reactions.

  • IGC susceptibility of the selected alloy is governed by microstructure as determined by the state of ageing. It is identified as slight to moderate pitting in highly under-aged, moderate to heavy pitting in under-aged, pitting with localized ICG in peak-aged, localized to uniform IGC in over-aged, and uniform IGC to etching in highly over-aged state. The IGC susceptibility is controlled by the anodic dissolution of the precipitate free zones and closely neighboured grain boundary precipitates of AlCrMnFeSi acting as cathodes specifically in the highly over-aged alloy.

  • The variation of corrosion behavior with the state of artificial ageing is related to the formation of the micro-galvanic coupling between the anodic precipitate free zones and the cathodic matrix grains with precipitates such as β'' (Mg5Si6), β' (Mg9Si6) and β (Mg2Si), and /or grain boundary precipitates of β (Mg2Si) and Fe-rich inclusions as evolved with the progress of ageing.

Conflict of interest

The author declares no conflict of interest.


Authors would like to thank Mr. Kuna Pradhan for his support in carrying out the experiments. Authors also wish to gratefully acknowledge the financial assistance received from the Center of Excellence for Microstructurally Designed Advanced Materials Development, TEQIP-II, IIEST Shibpur for carrying out some part of the present work.

Y. Wu, H. Liao.
Corrosion behavior of extruded near Eutectic Al-Si-Mg and 6063 alloys.
J Mater Sci Technol, 29 (2013), pp. 380-386
H. Zhan, J.M.C. Mol, F. Hannour, L. Zhuang, H. Terryn, J.H.W. De Wit.
The influence of copper content on intergranular corrosion of model AlMgSi(Cu) alloys.
Mater Corros, 59 (2008), pp. 670-675
S.K. Kairy, P.A. Rometsch, K. Diao, J.F. Nie, C.H.J. Davies, N. Birbilis.
Exploring the electrochemistry of 6xxx series aluminium alloys as a function of Si to Mg ratio, Cu content, ageing conditions and microstructure.
Electrochim Acta, 190 (2016), pp. 92-103
F. Song, X. Zhang, S. Liu, Q. Tan, D. Li.
The effect of quench rate and overageing temper on the corrosion behaviour of AA7050.
Corros Sci, 78 (2014), pp. 276-286
C. Brito, T. Vida, E. Freitas, N. Cheung, J.E. Spinelli, A. Garcia.
Cellular/dendritic arrays and intermetallic phases affecting corrosion and mechanical resistances of an Al-Mg-Si alloy.
J Alloys Compd, 673 (2016), pp. 220-230
Y. Zheng, B. Luo, Z. Bai, J. Wang, Y. Yin.
Study of the Precipitation Hardening Behaviour and Intergranular Corrosion of Al-Mg-Si Alloys with Differing Si Contents.
Metals (Basel), 7 (2017), pp. 387
K. Gopala Krishna, G. Das, K. Venkateswarlu, K.C. Hari Kumar.
Studies on Aging and Corrosion Properties of Cryorolled Al–Zn–Mg–Cu (AA7075) Alloy.
Trans Indian Inst Met, 70 (2017), pp. 817-825
Q. Guan, J. Sun, W. Wang, J. Gao, C. Zou, J. Wang, et al.
Pitting corrosion of natural aged Al–Mg–Si extrusion profile.
Materials (Basel), 12 (2019), pp. 1081
S.K. Kairy, P.A. Rometsch, C.H.J. Davies, N. Birbilis.
The influence of copper additions and aging on the microstructure and metastable pitting of al-mg-si alloys.
Corrosion, 71 (2015), pp. 1304-1307
P. Deepa, R. Padmalatha.
Corrosion behaviour of 6063 aluminium alloy in acidic and in alkaline media.
Arab J Chem, 10 (2017), pp. S2234-44
S. De, S. Palit Sagar, S. Dey, A. Prakash, et al.
Quantification of pitting in two tempers of 7075 aluminium alloy by non-destructive evaluation.
Corros Sci, 52 (2010), pp. 1818-1823
V. Gadpale, P.N. Banjare, M.K. Manoj.
Effect of ageing time and temperature on corrosion behaviour of aluminum alloy 2014.
IOP Conf Ser Mater Sci Eng, 338 (2018),
M. Navaser, M. Atapour.
Effect of friction stir processing on pitting corrosion and intergranular attack of 7075 aluminum alloy.
J Mater Sci Technol, 33 (2017), pp. 155-165
Y. Yan, L. Peguet, O. Gharbi, A. Deschamps, C.R. Hutchinson, S.K. Kairy, et al.
On the corrosion, electrochemistry and microstructure of Al-Cu-Li alloy AA2050 as a function of ageing.
Materialia, 1 (2018), pp. 25-36
M. Cao, L. Liu, Z. Yu, L. Fan, Y. Li, F. Wang.
Electrochemical corrosion behavior of 2A02 Al alloy under an accelerated simulation marine atmospheric environment.
J Mater Sci Technol, 35 (2019), pp. 651-659
E.N. Codaro, R.Z. Nakazato, A.L. Horovistiz, L.M.F. Ribeiro, R.B. Ribeiro, L.R.O. Hein.
An image processing method for morphology characterization and pitting corrosion evaluation.
Mater Sci Eng A, 334 (2002), pp. 298-306
S. Dey, M.K. Gunjan, I. Chattoraj.
Effect of temper on the distribution of pits in AA7075 alloys.
Corros Sci, 50 (2008), pp. 2895-2901
S. Dey, S.K. Das, A. Basumallick, I. Chattoraj.
The effect of pitting on fatigue lives of peak-aged and overaged 7075 aluminum alloys.
Metall Mater Trans A Phys Metall Mater Sci, 41 (2010), pp. 3297-3307
S.K. Kairy, P.A. Rometsch, C.H.J. Davies, N. Birbilis.
On the Intergranular Corrosion and Hardness Evolution of 6xxx Series Al Alloys as a Function of Si:Mg Ratio, Cu Content, and Aging Condition.
Corrosion, 73 (2017), pp. 1280-1295
C. Schnatterer, C. Altenbach, D. Zander.
The effect of simulated in-service heat impact on the microstructure and corrosion properties of a high Cu containing Al-Mg-Si alloy.
Mater Corros, 70 (2019), pp. 1205-1213
F.F. Chen, I. Cole, A.E. Hughes, A.M. Glenn, E. Sapper, J. Osborne.
Microstructure characterisation and reconstruction of intermetallic particles.
Mater Corros, 65 (2014), pp. 664-669
D. Zander, C. Schnatterer, C. Altenbach, V. Chaineux.
Microstructural impact on intergranular corrosion and the mechanical properties of industrial drawn 6056 aluminum wires.
Mater Des, 83 (2015), pp. 49-59
V. Guillaumin, G. Mankowski.
Influence of Overaging Treatment on Localized Corrosion of Al 6056.
Corrosion, 56 (2000), pp. 12-23
Z. Wang, H. Li, F. Miao, W. Sun, B. Fang, R. Song, et al.
Improving the intergranular corrosion resistance of Al-Mg-Si-Cu alloys without strength loss by a two-step aging treatment.
Mater Sci Eng A, 590 (2014), pp. 267-273
G. Svenningsen, M.H. Larsen, J.C. Walmsley, J.H. Nordlien, K. Nisancioglu.
Effect of artificial aging on intergranular corrosion of extruded AlMgSi alloy with small Cu content.
Corros Sci, 48 (2006), pp. 1528-1543
H. Li, P. Zhao, Z. Wang, Q. Mao, B. Fang, R. Song, et al.
The intergranular corrosion susceptibility of a heavily overaged Al-Mg-Si-Cu alloy.
Corros Sci, 107 (2016), pp. 113-122
M. Cornejo, T. Hentschel, D. Koschel, C. Matthies, L. Peguet, M. Rosefort, et al.
Intergranular corrosion testing of 6000 aluminum alloys.
Mater Corros, 69 (2018), pp. 626-633
M.H. Larsen, J.C. Walmsley, O. Lunder, R.H. Mathiesen, K. Nisancioglu.
Intergranular Corrosion of Copper-Containing AA6xxx AlMgSi Aluminum Alloys.
J Electrochem Soc, 155 (2008), pp. C550
S.K. Kairy, T. Alam, P.A. Rometsch, C.H.J. Davies, R. Banerjee, N. Birbilis.
Understanding the origins of intergranular corrosion in copper-containing Al-Mg-Si alloys.
Metall Mater Trans A Phys Metall Mater Sci, 47 (2016), pp. 985-989
Y. Zou, Q. Liu, Z. Jia, Y. Xing, L. Ding, X. Wang.
The intergranular corrosion behavior of 6000-series alloys with different Mg/Si and Cu content.
Appl Surf Sci, 405 (2017), pp. 489-496
W.J. Liang, P.A. Rometsch, L.F. Cao, N. Birbilis.
General aspects related to the corrosion of 6xxx series aluminium alloys: Exploring the influence of Mg/Si ratio and Cu.
Corros Sci, 76 (2013), pp. 119-128
W. Cubberly, H. Barker, D. Benjamin.
ASM handbook.
9th ed., ASM International, (1979),
D.S. Kharitonov, C. Örnek, P.M. Claesson, J. Sommertune, I.M. Zharskii, I.I. Kurilo, et al.
Corrosion Inhibition of Aluminum Alloy AA6063-T5 by Vanadates: Microstructure Characterization and Corrosion Analysis.
J Electrochem Soc, 165 (2018), pp. C116-26
ISO 8407-2009. Corrosion of metals and alloys: Removal of corrosion products from corrosion test specimens n.d.
M. Starostin, G.E. Shter, G.S. Grader.
Corrosion of aluminum alloys Al 6061 and Al 2024 in ammonium nitrate-urea solution.
Mater Corros, 67 (2016), pp. 387-395
Determination of resistance to IGC of solution heat‐treatable aluminium alloys, Standard BS 11486:1995.
British Standard Institution, (1995),
D.M. Jiang, B.D. Hong, T.C. Lei, D.A. Downham, G.W. Lorimer.
Fracture Behaviour of Aluminium Alloy 6063.
Mater Sci Technol, 7 (1991), pp. 1010-1014
R.A. Siddiqui, H.A. Abdullah, K.R. Al-Belushi.
Influence of aging parameters on the mechanical properties of 6063 aluminium alloy.
J Mater Process Technol, 102 (2000), pp. 234-240
S.K. Panigrahi, R. Jayaganthan, V. Chawla.
Effect of cryorolling on microstructure of Al-Mg-Si alloy.
Mater Lett, 62 (2008), pp. 2626-2629
C.D. Marioara, S.J. Andersen, H.W. Zandbergen, R. Holmestad.
The influence of alloy composition on precipitates of the Al-Mg-Si system.
Metall Mater Trans A Phys Metall Mater Sci, 36 (2005), pp. 691-702
M.W. Zandbergen, Q. Xu, A. Cerezo, G.D.W. Smith.
Study of precipitation in Al-Mg-Si alloys by Atom Probe Tomography I. Microstructural changes as a function of ageing temperature.
Acta Mater, 101 (2015), pp. 136-148
S. Nandy, K. Kumar Ray, D. Das.
Process model to predict yield strength of AA6063 alloy.
Mater Sci Eng A, 644 (2015), pp. 413-424
R. Vissers, M.A. van Huis, J. Jansen, H.W. Zandbergen, C.D. Marioara, S.J. Andersen.
The crystal structure of the β′ phase in Al-Mg-Si alloys.
Acta Mater, 55 (2007), pp. 3815-3823
S.J. Andersen, C.D. Marioara, R. Vissers, A. Frøseth, H.W. Zandbergen.
The structural relation between precipitates in Al-Mg-Si alloys, the Al-matrix and diamond silicon, with emphasis on the trigonal phase U1-MgAl2Si2.
Mater Sci Eng A, 444 (2007), pp. 157-169
S.J. Andersen, C.D. Marioara, A. Frøseth, R. Vissers, H.W. Zandbergen.
Crystal structure of the orthorhombic U2-Al4Mg4 Si4 precipitate in the Al-Mg-Si alloy system and its relation to the β′ and β″ phases.
Mater Sci Eng A, 390 (2005), pp. 127-138
M. Jacobs.
The structure of the metastable precipitates formed during ageing of an Al-Mg-Si alloy.
L. Ding, Z. Jia, J.F. Nie, Y. Weng, L. Cao, H. Chen, et al.
The structural and compositional evolution of precipitates in Al-Mg-Si-Cu alloy.
Acta Mater, 145 (2018), pp. 437-450
C.L. He, X.D. Meng, G.F. Ma, J.M. Wang, Z.F. Du, D.L. Zhao.
Hardness and corrosion behavior of aged 6063 Al alloy.
Adv Mater Res, 1046 (2014), pp. 54-57
S. Nandy, M.A. Bakkar, D. Das.
Influence of ageing on mechanical properties of 6063 Al alloy.
Mater Today Proc, 2 (2015), pp. 1234-1242
J.P. Bandstra, D.A. Koss.
On the influence of void clusters on void growth and coalescence during ductile fracture.
Acta Mater, 56 (2008), pp. 4429-4439
D. Broek.
The role of inclusions in ductile fracture and fracture toughness.
Eng Fract Mech, 5 (1973), pp. 55-66
A. Kisasoz.
Corrosion behavior of alloy AA6063-T4 in HCl and NaOH solutions.
Mater Test, 60 (2018), pp. 478-482
M.R. Tabrizi, S.B. Lyon, G.E. Thompson, J.M. Ferguson.
The long-term corrosion of aluminium alkaline media.
Corros Sci, 32 (1991), pp. 733-742
R.S. Alwitt, L.C. Archibald.
Some observations on the hydrous oxide film on aluminium immersed in warm water.
Corros Sci, 13 (1973), pp. 687-688
G.S. Frankel.
Pitting corrosion of metals a review of critical factors.
J Electrochem Soc, 145 (1998), pp. 2186-2198
J.A. Richardson, G.C. Wood.
A study of the pitting corrosion of Al byscanning electron microscopy.
Corros Sci, 10 (1970), pp. 313-323
F. Eckermann, T. Suter, P.J. Uggowitzer, A. Afseth, P. Schmutz.
The influence of MgSi particle reactivity and dissolution processes on corrosion in Al-Mg-Si alloys.
Electrochim Acta, 54 (2008), pp. 844-855
R. Ly, K.T. Hartwig, H. Castaneda.
Influence of dynamic recrystallization and shear banding on the localized corrosion of severely deformed Al–Mg–Si alloy.
Materialia, 4 (2018), pp. 457-465
K.D. Ralston, N. Birbilis, M. Weyland, C.R. Hutchinson.
The effect of precipitate size on the yield strength-pitting corrosion correlation in Al-Cu-Mg alloys.
Acta Mater, 58 (2010), pp. 5941-5948
N. Birbilis, M.K. Cavanaugh, L. Kovarik, R.G. Buchheit.
Nano-scale dissolution phenomena in Al-Cu-Mg alloys.
Electrochem commun, 10 (2008), pp. 32-37
R.K. Gupta, A. Deschamps, M.K. Cavanaugh, S.P. Lynch, N. Birbilis.
Relating the early evolution of microstructure with the electrochemical response and mechanical performance of a Cu-rich and Cu-lean 7xxx aluminum alloy.
J Electrochem Soc, 159 (2012), pp. 492-502
K. El-Menshawy, A.W.A. El-Sayed, M.E. El-Bedawy, H.A. Ahmed, et al.
Effect of aging time at low aging temperatures on the corrosion of aluminum alloy 6061.
Corros Sci, 54 (2012), pp. 167-173
H.N. Soliman.
Influence of 8-hydroxyquinoline addition on the corrosion behavior of commercial Al and Al-HO411 alloys in NaOH aqueous media.
Corros Sci, 53 (2011), pp. 2994-3006
H.H. Uhlig, R.W. Revie.
Corrosion and its control.
third ed., Wiley, (1991),
I. Boukerche, S. Djerad, L. Benmansour, L. Tifouti, K. Saleh.
Degradability of aluminum in acidic and alkaline solutions.
Corros Sci, 78 (2014), pp. 343-352
R.K. Hart.
The formation of films on aluminium immersed in water.
Trans Faraday Soc, 53 (1957), pp. 1020-1027
I.B. Obot, N.O. Obi-Egbedi, S.A. Umoren, E.E. Ebenso.
Synergistic and antagonistic effects of anions and ipomoea invulcrata as green corrosion inhibitor for aluminium dissolution in acidic medium.
Int J Electrochem Sci, 5 (2010), pp. 994-1007
M.H. Larsen, J.C. Walmsley, O. Lunder, K. Nisancioglu.
Effect of excess silicon and small copper content on intergranular corrosion of 6000-series aluminum alloys.
J Electrochem Soc, 157 (2010), pp. C61-8
K. Srinivasa Rao, K. Prasad Rao.
Pitting Corrosion of Heat-Treatable Aluminium Alloys and Welds: A Review.
Trans Indian Inst Met, 57 (2004), pp. 593-610
G. Svenningsen, J.E. Lein, A. Bjørgum, J.H. Nordlien, Y. Yu, K. Nisancioglu.
Effect of low copper content and heat treatment on intergranular corrosion of model AlMgSi alloys.
Corros Sci, 48 (2006), pp. 226-242
J.F. Li, Z.Q. Zheng, S.C. Li, W.J. Chen, W.D. Ren, X.S. Zhao.
Simulation study on function mechanism of some precipitates in localized corrosion of Al alloys.
Corros Sci, 49 (2007), pp. 2436-2449
A. Shi, B.A. Shaw, E. Sikora.
The role of grain boundary regions in the localized corrosion of a copper-free 6111-like aluminum alloy.
Corrosion, 61 (2005), pp. 534-547
J.R. Galvele, Micheli SMDEDE, B. Aires.
Mechanism of Intergranular corrosion of Al-Cu alloys*.
Corrosion, 10 (1970), pp. 795-807
F.L. Zeng, Z.L. Wei, J.F. Li, C.X. Li, X. Tan, Z. Zhang, et al.
Corrosion mechanism associated with Mg 2 Si and Si particles in Al-Mg-Si alloys.
Trans Nonferrous Met Soc China (English Ed, 21 (2011), pp. 2559-2567
A.K. Gupta, D.J. Lloyd, S.A. Court.
Precipitation hardening in Al-Mg-Si alloys with and without excess Si.
Mater Sci Eng A, 316 (2001), pp. 11-17
S. Saikawa, G. Aoshima, S. Ikeno, K. Morita, N. Sunayama, K. Komai.
Microstructure and mechanical properties of an Al-Zn-Mg-Cu alloy produced by gravity casting process.
Arch Metall Mater, 60 (2015), pp. 871-874
W. Sha, A.E. Long.
Quantification of overaging hardening kinetics of aluminum alloys.
Metall Mater Trans A Phys Metall Mater Sci, 35 (2004), pp. 2172-2174
D. Wang, Z.Y. Ma.
Effect of pre-strain on microstructure and stress corrosion cracking of over-aged 7050 aluminum alloy.
J Alloys Compd, 469 (2009), pp. 445-450
M. Merklein, J. Degner.
Influence of pre-strain and simulated paint-bake on mechanical properties of high strength aluminum alloy AA7020.
Appl Mech Mater, 805 (2015), pp. 115-122
A.P. Sekhar, D. Das.
Corrosion behavior of under‐, peak‐, and over‐aged 6063 alloy: a comparative study.
J. Li, F. Li, X. Ma, J. Li, S. Liang.
Effect of grain boundary characteristic on intergranular corrosion and mechanical properties of severely sheared Al-Zn-Mg-Cu alloy.
Mater Sci Eng A, 732 (2018), pp. 53-62
S. Kumari, S. Wenner, J.C. Walmsley, O. Lunder, K. Nisancioglu.
Progress in understanding initiation of intergranular corrosion on AA6005 aluminum alloy with low copper content.
J Electrochem Soc, 166 (2019), pp. C3114-23
Copyright © 2019. The Authors
Journal of Materials Research and Technology

Subscribe to our newsletter

Article options
Cookies policy
To improve our services and products, we use cookies (own or third parties authorized) to show advertising related to client preferences through the analyses of navigation customer behavior. Continuing navigation will be considered as acceptance of this use. You can change the settings or obtain more information by clicking here.