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Improvements in mechanical and stress corrosion cracking properties in Al-alloy 7075 via retrogression and reaging
Materials Science and Engineering A 485 (2008) 468–475
Improvements in mechanical and stress corrosion cracking properties in Al-alloy 7075 via retrogression and reaging Y. Reda, R. Abdel-Karim ∗ , I. Elmahallawi Department of Metallurgy, Faculty of Engineering, Cairo University, Giza, Egypt Received 4 September 2006; received in revised form 4 August 2007; accepted 7 August 2007
Abstract Aluminum alloy 7075 has been widely studied, due to its excellent mechanical properties developed by age hardening and their extensive uses in the aircraft structure. Tensile test specimens were tested in various preaging conditions (100, 120, 140 ◦ C) and various retrogression temperatures T7 (160, 180, 200 ◦ C) in order to evaluate the effect of heating cycles on the mechanical as well as corrosion properties. The optimum condition was preaging at T6 120 ◦ C and retrogressing at T7 200 ◦ C, which gave the highest hardness and tensile properties. The corrosion behavior of three-point-loaded specimens after 7 days of immersion in 3.5% NaCl solution, using modified electrochemical cell, was studied. The most resisting condition for SCC was preaging at T6 = 100 ◦ C and retrogression at T7 = 160 ◦ C for 250 min. Specimens retrogressed at 200 ◦ C for 8 min showed low resistance to stress corrosion cracking. © 2007 Elsevier B.V. All rights reserved. Keywords: Aluminum alloy 7075; Retrogression; Reaging; SCC
1. Introduction Aluminum alloy 7075 is widely used for aircraft structural materials because it is a high strength and a low density [1]. Because of the Navys unique service requirements, this alloy is subjected to aggressive conditions where it often encounter salt water spray and/or salt fog environments. Since commercial purity aluminum alloys contain numerous constituent particles that have electrochemical potentials different from that of the matrix and corrosion pits can readily develop at these particles. Once corrosion pits are formed they act as stress concentration sites leading to stress corrosion cracking [2]. In commercial aluminum alloys, pitting corrosion has been found to occur at intermetallic constituent particles. According to Gao et al. [3], two types of particles were identified. Type A particles were anodic with respect to the matrix and tend to dissolve themselves, while type C particles were cathodic to the matrix and tended to promote dissolution of the adjacent matrix. Types A particles were those with Al, Mg, and Zn and type C particles were those with Fe, Cu, Mn.
The conventional method of solving the low corrosion resistance problem has been to overage the material (T73). Consequently a strength loss of 10–15% was inevitable. Retrogression and reaging was advised in order to overcome this problem. This method consists of retrogression the T6 structure at a high temperature within the two-phase field, then reaging at the original T6 condition [1]. Cina and Ranish [4] utilized temperatures T6 = 120 ◦ C for preaging treatment, and temperature T7 = 200 ◦ C for the retrogression process, then reaging the alloy at low temperature similar to the T6 aging temperature and time. Retrogression and reaging results in an optimum combination of corrosion resistance and mechanical properties. The objectives of this study were to quantify the effect of RRA treatments with preaging temperature ≤140 ◦ C and low retrogression temperature ≤200 ◦ C on the strength and SCC resistance of alloy 7075, by applying a modified electrochemical potentiodynamic testing using three-point-loaded tensile specimens. The effects of heat treatment cycles as well as the applied stress were studied. 2. Experimental techniques
∗
Corresponding author. Tel.: +2 02 0106690945. E-mail address: [email protected] (R. Abdel-Karim).
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.08.025
The material used in this investigation was Aluminum alloy (7075-T0 ), delivered in the form of a 0.5 mm thick plate of the chemical analysis illustrated in Table 1.
Y. Reda et al. / Materials Science and Engineering A 485 (2008) 468–475
469
Table 1 Chemical analysis of Al-alloy 7075 (wt.%) Zn
Mg
Cu
Mn
Si
Fe
Cr
Ti
Total other
Al
5.7
2.4
1.6
0.3
0.4
0.5
0.2
0.2
0.15
Rem
Fig. 3. Modified electrochemical cell showing the rotating wheel forcing the fiber pen.
the following equation: σuts = Fig. 1. Tensile specimen.
All specimens were machined according to requirements of ASTM B557 [5], for both tensile and corrosion testing (Fig. 1). Tensile specimens taken from as received condition, were solution treated by heating to 470 ◦ C for 30 min; water quenching, then were artificially preaged at 100, 120 ◦ C as well as 140 ◦ C for 24 h, and retrogressed at different temperatures namely; 160, 180 and 200 ◦ C. All specimens were finally reaged at 120 ◦ C for 24 h. The heat treatment duration was chosen according to previous data [6] provided: retrogression at 160 ◦ C for 250 min, 180 ◦ C for 34 min and 200 ◦ C for 8 min, which gave the best values of mechanical properties. The Brinell Hardness (HB5) was measured by using Wilson hardness testing machine, having a hardened spherical ball of diameter 1/16 mm. A load of 15 was used as per requirements for non-ferrous metal. The reported results are an average of six readings. The tensile specimens for corrosion testing were three-pointloaded using a special teflon frame (Fig. 2). A motor driven rotating wheel (Fig. 3) was forcing a fiber pin leading to a specimen’s displacement L = 4 mm. The stress applied was equivalent to σ = 178 MPa, which was calculated according to
Fig. 2. Three-point-loaded specimen for SCC testing.
6EtY h2
(1)
where E is the modulus of elasticity (66 GPa), t thickness of specimen (0.5 mm) Y displacement (4 mm) and h is the distance between fixed point (35 mm). The corrosion cell was composed of test specimen, auxiliary graphite electrode and a saturated calomel reference electrode. The loading frame and test specimen were immersed in 3.5% NaCl solution for 7 days before applying electrochemical testing. All the measurements took place using a scan range from −200 to1000 mV with a scan rate 0.5 mV/s, using a power supply Meinsberg MS6 galvanostat/potentiostat, under stress and non-stress conditions. Evaluation of corrosion rate in mm/year was obtained using Tafel Extrapolation technique. 3. Results 3.1. Microstructure characterization The effect of preaging and retrogression cycles on precipitation morphology of aluminum alloy 7075 is illustrated in Fig. 4a–l using SEM. Generally, the structure of this alloy consists of a mixture of ␣ phase (S.S of Zn in Al) and intermetallic second phase.These areas appear by using HNO3 etchant solution as white and dark area, respectively. At T6 = 100 ◦ C, Fig. 4a shows constituent particles aligned in the direction of working, with large areas of ␣ phase. Theses constituents primarily evolve from the presence of Fe, Mn and Si particles. They are irregularly shaped and can be formed during alloy solidification whilst not being appreciably dissolved during subsequent thermomechanical processing. Rolling and extrusion tends to break up and align constituent particles into bands within the alloy. Often constituents are found in colonies made up of several intermetallic crystals. Typical examples include Al3 Fe and Al7 Cu2 Fe [7]. Such particles do not represent a significant aspect in the development of mechanical properties in high strength Al alloys.
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After retrogression, the micrographs (Fig. 4b–d) reveal the formation of fine intermetallic particles, distributed inside the grains, with the phases formed at the grain boundary being coarser and less spread. Retrogression at 200 ◦ C, led to the highest phase intensity among the other two retrogression temperatures 160 and 180 ◦ C. Similarly, Fig. 4e shows the rolling direction with some fine constituents by preaging at 120 ◦ C. By increasing the retrogres-
sion temperature, the intensity of both grains and grain boundary intermetallic phases increases (Fig. 4f–h). Max intensity of intermetallic particles is revealed by retrogression at 200 ◦ C (Fig. 4h). Clusters of non-uniform intermetallics are formed due to preaging at 140 ◦ C and retrogression at 160 ◦ C. The microstructures corresponding to retrogression at 180 and 200 ◦ C are illustrated in Fig. 4k and l, respectively, where the secondary intermetallics was spread all over the grains, with coarser par-
Fig. 4. Microstructure of alloy 7075 after preaging and retrogression.
Y. Reda et al. / Materials Science and Engineering A 485 (2008) 468–475
471
Fig. 4. (Continued ).
ticles at the grain boundaries, upon increasing the retrogression temperature. 3.2. Mechanical properties Table 2 illustrates the change in hardness values for the studied alloy after preaging at different temperatures (100, 120 and
140 ◦ C) for 24 h, followed by retrogression at different temperatures and reaging at 120 ◦ C for 24 h. The hardness values were 116, 124, 116 for preaging at 100, 120 and 140 ◦ C, respectively. Increasing the preaging temperature at 120 and 140 ◦ C, did not yield a significant change in hardness. Preaging at 120 ◦ C for 24 h yielded a comparable hardness (130 HB5) to other treat-
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Table 2 Effect of retrogression and reaging on mechanical properties of alloy 7075 Condition
Hardness (HB5)
Tensile strength (MPa)
100 ◦ C
T7 160 ◦ C T7 180 ◦ C T7 200 ◦ C
116 102 128 114
546.19 469.3 531.19 622.43
T6 120 ◦ C T7 160 ◦ C T7 180 ◦ C T7 200 ◦ C
95 124 128 130
558.26 672.89 688.84 698
T6 140 ◦ C T7 160 ◦ C T7 180 ◦ C T7 200 ◦ C
116 126 124 126
468.06 431.02 526.25 547.65
T6
ment, only when followed by retrogression at 200 ◦ C for 8 min. Considering the influence of heating on tensile strength, preaging at 120 ◦ C gave the highest tensile properties among all preaging conditions Upon preaging at 120 ◦ C the ultimate tensile strength approached 558.26 Mpa, compared with 546.19 and 468.06 Mpa for T6 = 100 and 140 ◦ C, respectively. Higher retrogression temperature led to higher tensile values for all preaging ranges. Preaging at T6 = 120 ◦ C followed by T7 = 200 ◦ C resulted in the highest tensile properties (698 Mpa). The effect of heating cycle on strength was more significant than its effect on hardness. 3.3. Corrosion testing The results of the potentiodynamic polarization tests of all specimens are illustrated in Fig. 5a–f and Table 3. Generally, stressed specimens showed lower corrosion resistance than unstressed specimens. Under no stress conditions and at T6 = 100 ◦ C and T7 = 160 ◦ C (Fig. 5a), the polarization curve starts with a gradual anodic dissolution in the range of current density from 1.2 × 10−3 to 1.2 mA/cm2 , with a potential range from −200 to 0 mV, then
it reaches a stability region at current density 1.2 mA/cm2 , for a potential range from 0 to 130 mV. A sudden increase in the current density from 1.2 to1.2 × 100.5 mA/cm2 is observed at a potential 130 mV, followed by another steady state condition at potential 130–500 mV. One breakdown potentials was observed at 130 mV. The corrosion potential approached −200 mV. Due to retrogression at 180 ◦ C, the curve starts with a gradual anodic dissolution in the range of current density from 0.2 to 20 mA/cm2 , in a potential range from −770 to −680 mV, followed by a sudden increase in the current density from 20 to 70 mA/cm2 at a potential range from −680 to −380 mV. Only one break down potential was registered at −680 mV. At higher retrogression temperature (200 ◦ C), general continuous anodic dissolution is detected. The corrosion potential remained at the same active value as for 180 ◦ C. In case of T7 = 160 ◦ C (Fig. 5b), the curve starts with a gradual anodic dissolution in the range of current density from 10−1.5 to 10−0.5 mA/cm2 in the potential range from −170 to 250 mV (breakdown potential), followed by an irregular increase in the current density from 10−0.5 to10 mA/cm2 . The potential range changes from 250 to 440 mV in this region. The combination of preaging at 100 ◦ C and retrogression at 160 ◦ C lead to the low corrosion rate (0.029 mm/y) with less active corrosion potential (−200 mV), compared to all other combinations of preaging at 100 ◦ C and retrogression at different temperatures. A high corrosion rate (0.619 mm/y) with the most negative corrosion potential (∼−770 mV) was observed under heat treatment conditions T6 = 100 ◦ C and T7 = 200 ◦ C. Fig. 5c and d compares the corrosion behavior of unstressed and stressed specimens after preaging at 120 ◦ C and retrogression at 160, 180, and 200 ◦ C. For T7 = 160 ◦ C, the polarization curve is also composed of three regions under the no stress condition. It starts with a gradual anodic dissolution in the range of current density from 0.004 to 1.2 mA/cm2 with increasing the potential from −350 to −40 mV. At region 2, a sudden increase in the current density from 1.2 to 12 mA/cm2 at almost constant potential value of −40 mV. At region 3, a gradual anodic dissolution as the potential raise from −40 to 0 mV was formed. While at T7 = 180 ◦ C, the curve starts with a gradual anodic dis-
Table 3 Corrosion data of aluminum alloy 7075 in 3.5% NaCl solution (mm/y) Condition
No Stress
Stress
Corrosion potential (mV)
Corrosion rate (mm/y)
Corrosion potential, (mV)
Corrosion rate (mm/y)
T7 = 160 ◦ C T7 = 180 ◦ C T7 = 200 ◦ C
−200 −300 −800
0.019 0.045 0.278
−200 −770 −768
0.029 0.197 0.619
T6 = 120 ◦ C T7 = 160 ◦ C T7 = 180 ◦ C T7 = 200 ◦ C
−290 −310 −380
0.021 0.028 0.041
−200 −300 −800
0.198 0.311 1.327
T6 = 140 ◦ C T7 = 160 ◦ C T7 = 180 ◦ C T7 = 200 ◦ C
−420 −310 −350
0.046 0.028 0.033
−210 −405 −760
0.051 0.053 1.749
T6
= 100 ◦ C
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Fig. 5. Corrosion behavior of alloy 7075 tested in 3.5% NaCl solution after different preaging and retrogression treatments under stress and no stress conditions.
solution in the range of current density from 0.3 to 1 mA/cm2 with a potential range from −380 to 40 mV, then another anodic dissolution with a higher density from 1 to 1.2 × 102 with a potential range from 40 to 150 mV. The break down potentials E1 and E2 were detected at 40 and 150 mV. Also, for T7 200 ◦ C under no stress condition, the polarization curve is composed of three anodic dissolution regions with increasing slopes. Region 1 Region 2 Region 3
−380 to −200 mV −200 to 0 mV 0 to 200 mV
Anodic dissolution Anodic dissolution Steady state
This reflects two break down potentials occurring at −200 and 0 mV.Under stressed condition, the anodic current is sharply
increasing and no stability state regions are observed reflecting a continuous metal dissolution (Fig. 5d). The corrosion potential values for unstressed specimens were −290, −350, and −380 mV for 160, 180, 200 ◦ C, respectively. Meanwhile, for stressed specimens the corrosion potential was −200, −300 and −800 mV for previously mentioned retrogression temperatures. Increasing the preaging temperature to 140 ◦ C, specimens showed continuous anodic behavior for all retrogression temperatures (Fig. 5e). Under stressed conditions, Fig. 5f, at T7 = 160 ◦ C, the curve starts with a gradual anodic dissolution in the range of current density from 0.07 to 0.7 mA/cm2 with a potential range from −180 to −40 mV, followed by a steady state region with current density equal to 0.7 mA/cm2 , at a potential range from −40 to 40 mV. Then another steady state appeared with increase in current density from 0.7 to 3 mA/cm2 with a
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potential range change from 40 to 300 mV.For specimen subjected to T7 = 180 ◦ C, the curve starts with a anodic dissolution in the range of current density from 0.07 to 0.3 mA/cm2 with a potential range from −420 to −350 mV, then an irregular increase in the current density from 0.3 to 40 mA/cm2 , where the poten-tial changes from −350 to 90 mV. For specimen subjected to T7 = 200 ◦ C, a continuous sharp anodic dissolution is observed.funcited It can be concluded, that retrogression at low temperature (160 ◦ C) is responsible for the low corrosion rate and consequently a good corrosion resistance. A high corrosion rate accompanied with more active corrosion potential was recorded under the due to retrogression at T7 200 ◦ C. Under no stress conditions, the difference between the two breakdown potentials observed in some cases, increases with increasing the retrogression temperature. 4. Discussion High strength Al alloys such as AA 7075 are commonly used in aircraft structure applications. The addition of Cu greatly improves the mechanical strength of this alloy by precipitation hardening. The peak strength is developed through the T6 condition, while the T7 produces superior corrosion resistance. In general, these temper specific properties are controlled by the precipitation sequence: Solid solution → Guinier Preston (GP) zones → η → η equilibrium phase G-P zones are metastable, coherent solute clusters of Zn, Mg and Cu. The metastable η , MgZn2 or more correctly Mg(ZnCuAl)2 appear as discrete platelet particles that are semi coherent with the matrix, which is known to populate within the grains, and η is pseudostable, non-coherent of the same phase appearing as rods or plates, which is known to populate the grain boundary [8]. During retrogression, the hardness and yield strength of alloy 7075 decreases rapidly as the Guinier-Preston (GP) zones dissolve. Stage 2 is a transient recovery period where the hardness increases as the remaining η grows to a near optimum size and distribution. The hardness reaches a maximum and then begins to decrease as this η coarsen excessively and starts transforming to η [8]. According to Park [1], the high strength of the aluminum is believed to arise from the high concentration of the fine η particles within the matrix, while the η phase hardly contributes to it. With X-ray photoelectron spectroscopy (XPS) analysis, it was found that Al and Cu were rich in the grain boundary of AA 7055 Li, and inferred that it was θ (Al2 Cu) phase [9]. From Table 2, the specimens with T7 temper at temperature <200 ◦ C had lower hardness and strength compared with the T6 temper. But theses values exceeded that of T6 temper upon retrogression at 200 ◦ C. The maximum tensile strength (698 MPa) accompanied by maximum hardness (130 HB5) due to preaging at 120 ◦ C and retrogression at 200 ◦ C. The microstructure
was characterized by high intermetallic particles intensity both at grain boundaries and inside the grains. The stress corrosion resistance of the alloy is generally thought to be controlled by the microstructure near to grain boundaries (the size and spacing of precipitates, free precipitation zones, and solute concentration gradients). The effect of these parameters is however not fully understood, but it is well known that the susceptibility of these alloys to stress corrosion cracking decreases with ageing time [7]. Stressed specimens subjected to preaging at 100 ◦ C and retrogression at 160 ◦ C, proved to have the highest corrosion resistance. The microstructure under this condition was composed of ␣ matrix and fine intermetallic particles with low intensity. On the other hand, the specimens under these conditions exhibited low hardness and tensile values. The hardness and ultimate tensile strength values were 102 HB5 and 469.3 Mpa, respectively. According to Zielinski et al. [11], T6 temperature demonstrated the best mechanical properties and the highest susceptibility to stress corrosion cracking. The increase in retrogression temperature and decrease in retrogression time, positively influenced the resistance to SCC. According to Ferrer et al. [6] various retrogression and reaging tempers <200 ◦ C produced strength similar to that of T6 with improved SCC. The RRA temper with retrogression at 160 ◦ C for 660 min produced the greatest improvement in SCC, with only 4% reduction in strength below T6 . Specimens subjected to preaging at 120 ◦ C followed by retrogression at 200 ◦ C, showed an active corrosion behavior. The corrosion potential and corrosion rate were −800 mV and 1.32 mm/y, respectively. Under this condition, alloy 7075 proved to have the highest tensile strength (698 Mpa) and the highest hardness value (130 HB5). The reason for the enhanced SCC in RRA treated 7xxx series alloys, is that RRA promotes coarse precipitation of the equilibrium phase η in the grains and subgrain boundaries, while maintaining a fine distribution of η in the grain interiors. These coarse intermetallic particles then act as hydrogen trapping sites, locally reducing the hydrogen concentration in the matrix around the grain boundary [2]. The most resistant microstructure possess a great number of very small particles. The relative resistance to SCC is related to the hydrogen trapping by numerous interphase boundaries [10]. By studying the electrochemical polarization test results illustrated in Fig. 5, phenomena of two breakdown potential were only detected for T6 = 120 ◦ C. Meanwhile, only one break down potential was detected for all other T7 temper conditions. Specimens retrogressed at 200 ◦ C, showed continuous dissolution behavior. Maitra and English [12] have conducted electrochemical measurements on alloy AA 7075 and AA 7150 in the T6 temper in deareated NaCl solution. Two breakdown potentials have been found in potentiodynamic polarization scans. It was suggested that for AA 7075, the active break down potential was associated with IGC, and the noble breakdown potential with pitting in the matrix. They attributed the IGC susceptibility of the T6 temper to Mg and Zn solute segregation or enrichment in the grin bound-
Y. Reda et al. / Materials Science and Engineering A 485 (2008) 468–475
ary regions. The breakdown potential in AA 7075 is the result of preferential dissolution or dealloying of fine Mg(ZnCuAl)2 hardening precipitates in the matrix. Since these fine hardening precipitates are highly dispersed coherently or semicoherently in the Al solid solution containing Zn, Mg, and Cu, the preferential dissolution of theses fine precipitates may also cause the reaction of the Al solid solution in the matrix. This combined dissolution of fine hardening precipitates and surrounding Al solid solution results in the formation of an Al(OH)3 product layer, which limit the depth of attack, making it a transient process. When the applied potential is above the second breakdown potential, stable pits form on the matrix by breakdown of the passive film formed between E1 and E2 . Pitting corrosion usually occurs in the Al matrix near Cu or Fe- containing intermetallic particles owing to galvanic interaction with the Al matrix. IGC is generally believed along grain boundaries [13]. Only one breakdown potential was found for AA 7075 in the T7 temper. It is possible that the T7 temper decreases the critical current density for stable pitting [12]. 5. Conclusions As a result of the work took place, it was found that: 1. Preaging at T6 120 ◦ C and retrogressing at T7 200 ◦ C resulted in the highest hardness and tensile properties. 2. Preaging at T6 120 ◦ C and retrogressing at T7 200 ◦ C lead to deterioration of corrosion properties of alloy 7075. The
475
coarse precipitates formed during retrogression and reaging act as hydrogen trapping sites, locally reducing the hydrogen concentration in the matrix around the grain boundary. 3. The optimum condition for SCC was aging at T6 100 ◦ C and retrogression at T7 160 ◦ C. The relative resistance to SCC is related to the hydrogen trapping by numerous interphase boundaries. References [1] J.K. Park, Mater. Sci. Eng. A103 (1988) 223–231. [2] F. Viana, A.M.P. Pinto, H.M.C. Santos, A.B. Lopes, J. Mater. Prop. Technol. 92–93 (1999) 54–59. [3] M. Gao, C.R. Feng, et al., Met. Mater. Eng. Trans. A 29A (1998) 1145–1155. [4] Yasser Reda, Randa Abdel-Karim, Iman Elmahallawi, Cina, Ranish. Reducing stress corrosion cracking in aluminum Alloys, US Patent 3,856,584. [5] ASTM B597 preparation and use of bent beam stress corrosion test specimens, ASTM G39. [6] C.P. Ferrer, M.G. Koul, B.J. Connolly, A.L. Moran, Corrosion 59 (6) (2003) 520–528. [7] N. Birbilis, M.K. Cavanaugh, R.G. Buchheit, et al., Proceedings Symposium Applications of Materials Science to Military System, Materials Science and Technology “05, Pisttsburgh, PA, 2005. [8] J.M. Papazian, Mater. Sci. Eng. 79 (1986) 97–104. [9] G.-W. Zhu, D. Li, P.-Y. Liu, J.-H. Liu, Mater. Forum 28 (2004) 805–810. [10] J.S. Robinson, S.D. Whelan, R.L. Cudd, J. Mater. Sci. Technol. 15 (1999) 717–724. [11] A. Zielinski, M. Warmuzek, et al., Adv. Mater. Sci. 2 (1 (2)) (2002) 33–42. [12] S. Maitra, G.C. English, Metall. Mater. Trans A 12 (1981) 535. [13] Q. Meng, G.S. Frankel, JES 151 (5) (2004) B271–B283.
Improvements in mechanical and stress corrosion cracking properties in Al-alloy 7075 via retrogression and reaging Y. Reda, R. Abdel-Karim ∗ , I. Elmahallawi Department of Metallurgy, Faculty of Engineering, Cairo University, Giza, Egypt Received 4 September 2006; received in revised form 4 August 2007; accepted 7 August 2007
Abstract Aluminum alloy 7075 has been widely studied, due to its excellent mechanical properties developed by age hardening and their extensive uses in the aircraft structure. Tensile test specimens were tested in various preaging conditions (100, 120, 140 ◦ C) and various retrogression temperatures T7 (160, 180, 200 ◦ C) in order to evaluate the effect of heating cycles on the mechanical as well as corrosion properties. The optimum condition was preaging at T6 120 ◦ C and retrogressing at T7 200 ◦ C, which gave the highest hardness and tensile properties. The corrosion behavior of three-point-loaded specimens after 7 days of immersion in 3.5% NaCl solution, using modified electrochemical cell, was studied. The most resisting condition for SCC was preaging at T6 = 100 ◦ C and retrogression at T7 = 160 ◦ C for 250 min. Specimens retrogressed at 200 ◦ C for 8 min showed low resistance to stress corrosion cracking. © 2007 Elsevier B.V. All rights reserved. Keywords: Aluminum alloy 7075; Retrogression; Reaging; SCC
1. Introduction Aluminum alloy 7075 is widely used for aircraft structural materials because it is a high strength and a low density [1]. Because of the Navys unique service requirements, this alloy is subjected to aggressive conditions where it often encounter salt water spray and/or salt fog environments. Since commercial purity aluminum alloys contain numerous constituent particles that have electrochemical potentials different from that of the matrix and corrosion pits can readily develop at these particles. Once corrosion pits are formed they act as stress concentration sites leading to stress corrosion cracking [2]. In commercial aluminum alloys, pitting corrosion has been found to occur at intermetallic constituent particles. According to Gao et al. [3], two types of particles were identified. Type A particles were anodic with respect to the matrix and tend to dissolve themselves, while type C particles were cathodic to the matrix and tended to promote dissolution of the adjacent matrix. Types A particles were those with Al, Mg, and Zn and type C particles were those with Fe, Cu, Mn.
The conventional method of solving the low corrosion resistance problem has been to overage the material (T73). Consequently a strength loss of 10–15% was inevitable. Retrogression and reaging was advised in order to overcome this problem. This method consists of retrogression the T6 structure at a high temperature within the two-phase field, then reaging at the original T6 condition [1]. Cina and Ranish [4] utilized temperatures T6 = 120 ◦ C for preaging treatment, and temperature T7 = 200 ◦ C for the retrogression process, then reaging the alloy at low temperature similar to the T6 aging temperature and time. Retrogression and reaging results in an optimum combination of corrosion resistance and mechanical properties. The objectives of this study were to quantify the effect of RRA treatments with preaging temperature ≤140 ◦ C and low retrogression temperature ≤200 ◦ C on the strength and SCC resistance of alloy 7075, by applying a modified electrochemical potentiodynamic testing using three-point-loaded tensile specimens. The effects of heat treatment cycles as well as the applied stress were studied. 2. Experimental techniques
∗
Corresponding author. Tel.: +2 02 0106690945. E-mail address: [email protected] (R. Abdel-Karim).
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.08.025
The material used in this investigation was Aluminum alloy (7075-T0 ), delivered in the form of a 0.5 mm thick plate of the chemical analysis illustrated in Table 1.
Y. Reda et al. / Materials Science and Engineering A 485 (2008) 468–475
469
Table 1 Chemical analysis of Al-alloy 7075 (wt.%) Zn
Mg
Cu
Mn
Si
Fe
Cr
Ti
Total other
Al
5.7
2.4
1.6
0.3
0.4
0.5
0.2
0.2
0.15
Rem
Fig. 3. Modified electrochemical cell showing the rotating wheel forcing the fiber pen.
the following equation: σuts = Fig. 1. Tensile specimen.
All specimens were machined according to requirements of ASTM B557 [5], for both tensile and corrosion testing (Fig. 1). Tensile specimens taken from as received condition, were solution treated by heating to 470 ◦ C for 30 min; water quenching, then were artificially preaged at 100, 120 ◦ C as well as 140 ◦ C for 24 h, and retrogressed at different temperatures namely; 160, 180 and 200 ◦ C. All specimens were finally reaged at 120 ◦ C for 24 h. The heat treatment duration was chosen according to previous data [6] provided: retrogression at 160 ◦ C for 250 min, 180 ◦ C for 34 min and 200 ◦ C for 8 min, which gave the best values of mechanical properties. The Brinell Hardness (HB5) was measured by using Wilson hardness testing machine, having a hardened spherical ball of diameter 1/16 mm. A load of 15 was used as per requirements for non-ferrous metal. The reported results are an average of six readings. The tensile specimens for corrosion testing were three-pointloaded using a special teflon frame (Fig. 2). A motor driven rotating wheel (Fig. 3) was forcing a fiber pin leading to a specimen’s displacement L = 4 mm. The stress applied was equivalent to σ = 178 MPa, which was calculated according to
Fig. 2. Three-point-loaded specimen for SCC testing.
6EtY h2
(1)
where E is the modulus of elasticity (66 GPa), t thickness of specimen (0.5 mm) Y displacement (4 mm) and h is the distance between fixed point (35 mm). The corrosion cell was composed of test specimen, auxiliary graphite electrode and a saturated calomel reference electrode. The loading frame and test specimen were immersed in 3.5% NaCl solution for 7 days before applying electrochemical testing. All the measurements took place using a scan range from −200 to1000 mV with a scan rate 0.5 mV/s, using a power supply Meinsberg MS6 galvanostat/potentiostat, under stress and non-stress conditions. Evaluation of corrosion rate in mm/year was obtained using Tafel Extrapolation technique. 3. Results 3.1. Microstructure characterization The effect of preaging and retrogression cycles on precipitation morphology of aluminum alloy 7075 is illustrated in Fig. 4a–l using SEM. Generally, the structure of this alloy consists of a mixture of ␣ phase (S.S of Zn in Al) and intermetallic second phase.These areas appear by using HNO3 etchant solution as white and dark area, respectively. At T6 = 100 ◦ C, Fig. 4a shows constituent particles aligned in the direction of working, with large areas of ␣ phase. Theses constituents primarily evolve from the presence of Fe, Mn and Si particles. They are irregularly shaped and can be formed during alloy solidification whilst not being appreciably dissolved during subsequent thermomechanical processing. Rolling and extrusion tends to break up and align constituent particles into bands within the alloy. Often constituents are found in colonies made up of several intermetallic crystals. Typical examples include Al3 Fe and Al7 Cu2 Fe [7]. Such particles do not represent a significant aspect in the development of mechanical properties in high strength Al alloys.
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After retrogression, the micrographs (Fig. 4b–d) reveal the formation of fine intermetallic particles, distributed inside the grains, with the phases formed at the grain boundary being coarser and less spread. Retrogression at 200 ◦ C, led to the highest phase intensity among the other two retrogression temperatures 160 and 180 ◦ C. Similarly, Fig. 4e shows the rolling direction with some fine constituents by preaging at 120 ◦ C. By increasing the retrogres-
sion temperature, the intensity of both grains and grain boundary intermetallic phases increases (Fig. 4f–h). Max intensity of intermetallic particles is revealed by retrogression at 200 ◦ C (Fig. 4h). Clusters of non-uniform intermetallics are formed due to preaging at 140 ◦ C and retrogression at 160 ◦ C. The microstructures corresponding to retrogression at 180 and 200 ◦ C are illustrated in Fig. 4k and l, respectively, where the secondary intermetallics was spread all over the grains, with coarser par-
Fig. 4. Microstructure of alloy 7075 after preaging and retrogression.
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Fig. 4. (Continued ).
ticles at the grain boundaries, upon increasing the retrogression temperature. 3.2. Mechanical properties Table 2 illustrates the change in hardness values for the studied alloy after preaging at different temperatures (100, 120 and
140 ◦ C) for 24 h, followed by retrogression at different temperatures and reaging at 120 ◦ C for 24 h. The hardness values were 116, 124, 116 for preaging at 100, 120 and 140 ◦ C, respectively. Increasing the preaging temperature at 120 and 140 ◦ C, did not yield a significant change in hardness. Preaging at 120 ◦ C for 24 h yielded a comparable hardness (130 HB5) to other treat-
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Table 2 Effect of retrogression and reaging on mechanical properties of alloy 7075 Condition
Hardness (HB5)
Tensile strength (MPa)
100 ◦ C
T7 160 ◦ C T7 180 ◦ C T7 200 ◦ C
116 102 128 114
546.19 469.3 531.19 622.43
T6 120 ◦ C T7 160 ◦ C T7 180 ◦ C T7 200 ◦ C
95 124 128 130
558.26 672.89 688.84 698
T6 140 ◦ C T7 160 ◦ C T7 180 ◦ C T7 200 ◦ C
116 126 124 126
468.06 431.02 526.25 547.65
T6
ment, only when followed by retrogression at 200 ◦ C for 8 min. Considering the influence of heating on tensile strength, preaging at 120 ◦ C gave the highest tensile properties among all preaging conditions Upon preaging at 120 ◦ C the ultimate tensile strength approached 558.26 Mpa, compared with 546.19 and 468.06 Mpa for T6 = 100 and 140 ◦ C, respectively. Higher retrogression temperature led to higher tensile values for all preaging ranges. Preaging at T6 = 120 ◦ C followed by T7 = 200 ◦ C resulted in the highest tensile properties (698 Mpa). The effect of heating cycle on strength was more significant than its effect on hardness. 3.3. Corrosion testing The results of the potentiodynamic polarization tests of all specimens are illustrated in Fig. 5a–f and Table 3. Generally, stressed specimens showed lower corrosion resistance than unstressed specimens. Under no stress conditions and at T6 = 100 ◦ C and T7 = 160 ◦ C (Fig. 5a), the polarization curve starts with a gradual anodic dissolution in the range of current density from 1.2 × 10−3 to 1.2 mA/cm2 , with a potential range from −200 to 0 mV, then
it reaches a stability region at current density 1.2 mA/cm2 , for a potential range from 0 to 130 mV. A sudden increase in the current density from 1.2 to1.2 × 100.5 mA/cm2 is observed at a potential 130 mV, followed by another steady state condition at potential 130–500 mV. One breakdown potentials was observed at 130 mV. The corrosion potential approached −200 mV. Due to retrogression at 180 ◦ C, the curve starts with a gradual anodic dissolution in the range of current density from 0.2 to 20 mA/cm2 , in a potential range from −770 to −680 mV, followed by a sudden increase in the current density from 20 to 70 mA/cm2 at a potential range from −680 to −380 mV. Only one break down potential was registered at −680 mV. At higher retrogression temperature (200 ◦ C), general continuous anodic dissolution is detected. The corrosion potential remained at the same active value as for 180 ◦ C. In case of T7 = 160 ◦ C (Fig. 5b), the curve starts with a gradual anodic dissolution in the range of current density from 10−1.5 to 10−0.5 mA/cm2 in the potential range from −170 to 250 mV (breakdown potential), followed by an irregular increase in the current density from 10−0.5 to10 mA/cm2 . The potential range changes from 250 to 440 mV in this region. The combination of preaging at 100 ◦ C and retrogression at 160 ◦ C lead to the low corrosion rate (0.029 mm/y) with less active corrosion potential (−200 mV), compared to all other combinations of preaging at 100 ◦ C and retrogression at different temperatures. A high corrosion rate (0.619 mm/y) with the most negative corrosion potential (∼−770 mV) was observed under heat treatment conditions T6 = 100 ◦ C and T7 = 200 ◦ C. Fig. 5c and d compares the corrosion behavior of unstressed and stressed specimens after preaging at 120 ◦ C and retrogression at 160, 180, and 200 ◦ C. For T7 = 160 ◦ C, the polarization curve is also composed of three regions under the no stress condition. It starts with a gradual anodic dissolution in the range of current density from 0.004 to 1.2 mA/cm2 with increasing the potential from −350 to −40 mV. At region 2, a sudden increase in the current density from 1.2 to 12 mA/cm2 at almost constant potential value of −40 mV. At region 3, a gradual anodic dissolution as the potential raise from −40 to 0 mV was formed. While at T7 = 180 ◦ C, the curve starts with a gradual anodic dis-
Table 3 Corrosion data of aluminum alloy 7075 in 3.5% NaCl solution (mm/y) Condition
No Stress
Stress
Corrosion potential (mV)
Corrosion rate (mm/y)
Corrosion potential, (mV)
Corrosion rate (mm/y)
T7 = 160 ◦ C T7 = 180 ◦ C T7 = 200 ◦ C
−200 −300 −800
0.019 0.045 0.278
−200 −770 −768
0.029 0.197 0.619
T6 = 120 ◦ C T7 = 160 ◦ C T7 = 180 ◦ C T7 = 200 ◦ C
−290 −310 −380
0.021 0.028 0.041
−200 −300 −800
0.198 0.311 1.327
T6 = 140 ◦ C T7 = 160 ◦ C T7 = 180 ◦ C T7 = 200 ◦ C
−420 −310 −350
0.046 0.028 0.033
−210 −405 −760
0.051 0.053 1.749
T6
= 100 ◦ C
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Fig. 5. Corrosion behavior of alloy 7075 tested in 3.5% NaCl solution after different preaging and retrogression treatments under stress and no stress conditions.
solution in the range of current density from 0.3 to 1 mA/cm2 with a potential range from −380 to 40 mV, then another anodic dissolution with a higher density from 1 to 1.2 × 102 with a potential range from 40 to 150 mV. The break down potentials E1 and E2 were detected at 40 and 150 mV. Also, for T7 200 ◦ C under no stress condition, the polarization curve is composed of three anodic dissolution regions with increasing slopes. Region 1 Region 2 Region 3
−380 to −200 mV −200 to 0 mV 0 to 200 mV
Anodic dissolution Anodic dissolution Steady state
This reflects two break down potentials occurring at −200 and 0 mV.Under stressed condition, the anodic current is sharply
increasing and no stability state regions are observed reflecting a continuous metal dissolution (Fig. 5d). The corrosion potential values for unstressed specimens were −290, −350, and −380 mV for 160, 180, 200 ◦ C, respectively. Meanwhile, for stressed specimens the corrosion potential was −200, −300 and −800 mV for previously mentioned retrogression temperatures. Increasing the preaging temperature to 140 ◦ C, specimens showed continuous anodic behavior for all retrogression temperatures (Fig. 5e). Under stressed conditions, Fig. 5f, at T7 = 160 ◦ C, the curve starts with a gradual anodic dissolution in the range of current density from 0.07 to 0.7 mA/cm2 with a potential range from −180 to −40 mV, followed by a steady state region with current density equal to 0.7 mA/cm2 , at a potential range from −40 to 40 mV. Then another steady state appeared with increase in current density from 0.7 to 3 mA/cm2 with a
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potential range change from 40 to 300 mV.For specimen subjected to T7 = 180 ◦ C, the curve starts with a anodic dissolution in the range of current density from 0.07 to 0.3 mA/cm2 with a potential range from −420 to −350 mV, then an irregular increase in the current density from 0.3 to 40 mA/cm2 , where the poten-tial changes from −350 to 90 mV. For specimen subjected to T7 = 200 ◦ C, a continuous sharp anodic dissolution is observed.funcited It can be concluded, that retrogression at low temperature (160 ◦ C) is responsible for the low corrosion rate and consequently a good corrosion resistance. A high corrosion rate accompanied with more active corrosion potential was recorded under the due to retrogression at T7 200 ◦ C. Under no stress conditions, the difference between the two breakdown potentials observed in some cases, increases with increasing the retrogression temperature. 4. Discussion High strength Al alloys such as AA 7075 are commonly used in aircraft structure applications. The addition of Cu greatly improves the mechanical strength of this alloy by precipitation hardening. The peak strength is developed through the T6 condition, while the T7 produces superior corrosion resistance. In general, these temper specific properties are controlled by the precipitation sequence: Solid solution → Guinier Preston (GP) zones → η → η equilibrium phase G-P zones are metastable, coherent solute clusters of Zn, Mg and Cu. The metastable η , MgZn2 or more correctly Mg(ZnCuAl)2 appear as discrete platelet particles that are semi coherent with the matrix, which is known to populate within the grains, and η is pseudostable, non-coherent of the same phase appearing as rods or plates, which is known to populate the grain boundary [8]. During retrogression, the hardness and yield strength of alloy 7075 decreases rapidly as the Guinier-Preston (GP) zones dissolve. Stage 2 is a transient recovery period where the hardness increases as the remaining η grows to a near optimum size and distribution. The hardness reaches a maximum and then begins to decrease as this η coarsen excessively and starts transforming to η [8]. According to Park [1], the high strength of the aluminum is believed to arise from the high concentration of the fine η particles within the matrix, while the η phase hardly contributes to it. With X-ray photoelectron spectroscopy (XPS) analysis, it was found that Al and Cu were rich in the grain boundary of AA 7055 Li, and inferred that it was θ (Al2 Cu) phase [9]. From Table 2, the specimens with T7 temper at temperature <200 ◦ C had lower hardness and strength compared with the T6 temper. But theses values exceeded that of T6 temper upon retrogression at 200 ◦ C. The maximum tensile strength (698 MPa) accompanied by maximum hardness (130 HB5) due to preaging at 120 ◦ C and retrogression at 200 ◦ C. The microstructure
was characterized by high intermetallic particles intensity both at grain boundaries and inside the grains. The stress corrosion resistance of the alloy is generally thought to be controlled by the microstructure near to grain boundaries (the size and spacing of precipitates, free precipitation zones, and solute concentration gradients). The effect of these parameters is however not fully understood, but it is well known that the susceptibility of these alloys to stress corrosion cracking decreases with ageing time [7]. Stressed specimens subjected to preaging at 100 ◦ C and retrogression at 160 ◦ C, proved to have the highest corrosion resistance. The microstructure under this condition was composed of ␣ matrix and fine intermetallic particles with low intensity. On the other hand, the specimens under these conditions exhibited low hardness and tensile values. The hardness and ultimate tensile strength values were 102 HB5 and 469.3 Mpa, respectively. According to Zielinski et al. [11], T6 temperature demonstrated the best mechanical properties and the highest susceptibility to stress corrosion cracking. The increase in retrogression temperature and decrease in retrogression time, positively influenced the resistance to SCC. According to Ferrer et al. [6] various retrogression and reaging tempers <200 ◦ C produced strength similar to that of T6 with improved SCC. The RRA temper with retrogression at 160 ◦ C for 660 min produced the greatest improvement in SCC, with only 4% reduction in strength below T6 . Specimens subjected to preaging at 120 ◦ C followed by retrogression at 200 ◦ C, showed an active corrosion behavior. The corrosion potential and corrosion rate were −800 mV and 1.32 mm/y, respectively. Under this condition, alloy 7075 proved to have the highest tensile strength (698 Mpa) and the highest hardness value (130 HB5). The reason for the enhanced SCC in RRA treated 7xxx series alloys, is that RRA promotes coarse precipitation of the equilibrium phase η in the grains and subgrain boundaries, while maintaining a fine distribution of η in the grain interiors. These coarse intermetallic particles then act as hydrogen trapping sites, locally reducing the hydrogen concentration in the matrix around the grain boundary [2]. The most resistant microstructure possess a great number of very small particles. The relative resistance to SCC is related to the hydrogen trapping by numerous interphase boundaries [10]. By studying the electrochemical polarization test results illustrated in Fig. 5, phenomena of two breakdown potential were only detected for T6 = 120 ◦ C. Meanwhile, only one break down potential was detected for all other T7 temper conditions. Specimens retrogressed at 200 ◦ C, showed continuous dissolution behavior. Maitra and English [12] have conducted electrochemical measurements on alloy AA 7075 and AA 7150 in the T6 temper in deareated NaCl solution. Two breakdown potentials have been found in potentiodynamic polarization scans. It was suggested that for AA 7075, the active break down potential was associated with IGC, and the noble breakdown potential with pitting in the matrix. They attributed the IGC susceptibility of the T6 temper to Mg and Zn solute segregation or enrichment in the grin bound-
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ary regions. The breakdown potential in AA 7075 is the result of preferential dissolution or dealloying of fine Mg(ZnCuAl)2 hardening precipitates in the matrix. Since these fine hardening precipitates are highly dispersed coherently or semicoherently in the Al solid solution containing Zn, Mg, and Cu, the preferential dissolution of theses fine precipitates may also cause the reaction of the Al solid solution in the matrix. This combined dissolution of fine hardening precipitates and surrounding Al solid solution results in the formation of an Al(OH)3 product layer, which limit the depth of attack, making it a transient process. When the applied potential is above the second breakdown potential, stable pits form on the matrix by breakdown of the passive film formed between E1 and E2 . Pitting corrosion usually occurs in the Al matrix near Cu or Fe- containing intermetallic particles owing to galvanic interaction with the Al matrix. IGC is generally believed along grain boundaries [13]. Only one breakdown potential was found for AA 7075 in the T7 temper. It is possible that the T7 temper decreases the critical current density for stable pitting [12]. 5. Conclusions As a result of the work took place, it was found that: 1. Preaging at T6 120 ◦ C and retrogressing at T7 200 ◦ C resulted in the highest hardness and tensile properties. 2. Preaging at T6 120 ◦ C and retrogressing at T7 200 ◦ C lead to deterioration of corrosion properties of alloy 7075. The
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coarse precipitates formed during retrogression and reaging act as hydrogen trapping sites, locally reducing the hydrogen concentration in the matrix around the grain boundary. 3. The optimum condition for SCC was aging at T6 100 ◦ C and retrogression at T7 160 ◦ C. The relative resistance to SCC is related to the hydrogen trapping by numerous interphase boundaries. References [1] J.K. Park, Mater. Sci. Eng. A103 (1988) 223–231. [2] F. Viana, A.M.P. Pinto, H.M.C. Santos, A.B. Lopes, J. Mater. Prop. Technol. 92–93 (1999) 54–59. [3] M. Gao, C.R. Feng, et al., Met. Mater. Eng. Trans. A 29A (1998) 1145–1155. [4] Yasser Reda, Randa Abdel-Karim, Iman Elmahallawi, Cina, Ranish. Reducing stress corrosion cracking in aluminum Alloys, US Patent 3,856,584. [5] ASTM B597 preparation and use of bent beam stress corrosion test specimens, ASTM G39. [6] C.P. Ferrer, M.G. Koul, B.J. Connolly, A.L. Moran, Corrosion 59 (6) (2003) 520–528. [7] N. Birbilis, M.K. Cavanaugh, R.G. Buchheit, et al., Proceedings Symposium Applications of Materials Science to Military System, Materials Science and Technology “05, Pisttsburgh, PA, 2005. [8] J.M. Papazian, Mater. Sci. Eng. 79 (1986) 97–104. [9] G.-W. Zhu, D. Li, P.-Y. Liu, J.-H. Liu, Mater. Forum 28 (2004) 805–810. [10] J.S. Robinson, S.D. Whelan, R.L. Cudd, J. Mater. Sci. Technol. 15 (1999) 717–724. [11] A. Zielinski, M. Warmuzek, et al., Adv. Mater. Sci. 2 (1 (2)) (2002) 33–42. [12] S. Maitra, G.C. English, Metall. Mater. Trans A 12 (1981) 535. [13] Q. Meng, G.S. Frankel, JES 151 (5) (2004) B271–B283.