Study on double peaks aging strengthening and stress corrosion cracking behaviour of 7075 aluminium alloy

Abstract

The aging strengthening effect, stress corrosion cracking (SCC) behaviour and microstructure of 7075 aluminium alloys under double peaks aging treatment have been investigated in this paper. The results show that the hardness and strength of alloys under the second aging peak are equal or slightly higher than those under the first peak. The microstructure analysis demonstrates that the aging strengthening effect of the first peak and second peak mainly depend on fine and intensive GP zones (2-6 nm in size with an inter-particle distance of 4-10 nm) and η’ phase (4-8 nm in size), respectively. Meanwhile, the SCC resistance of specimens under the second aging peak is better than the first peak aging treated ones. Lower Mg and higher Cu segregation on grainboundary induced by second peak aging treatment are considered to be the reason for better SCC behaviour according to high-resolution transmission electron microscopy and energy dispersive spectrometer results.

Keywords

Al–Zn–Mg–Cu aluminium alloys double peaks aging treatment stress corrosion cracking microstructure

Introduction

Al–Zn–Mg–Cu (7000 series) aluminium alloys have been widely used in aeronautical and industrial fields for several decades due to their high strength, low density and excellent workability [1 3]. However, 7000 series aluminium alloys have shown significant susceptibility to stress corrosion cracking (SCC) in aqueous environments (particularly in chloride environments), which is the cause of premature failures in various accidents [4,5]. The heat treatment technology of 7000 series Al alloys and its relationship with SCC have been extensively investigated by researchers [6 10]. So far, there has not yet concluded a specific heat treatment that can both satisfy strength and SCC. For the highest strength temper (T6), 7000 series Al alloys are susceptible to intergranular stress attack; alloys with overaged temper (T73) show lower susceptibility to SCC but sacrifice strength a lot (about 10%–15% in comparison with T6) [11]; in addition, numerous investigators [12,13] have provided experimental evidence indicating that the Retrogression and Re-aging (RRA) process can ensure relatively high strength without affecting the stress corrosion resistance. However, from the perspective of industrial application, RRA treatment is much more complex and energy consumption in comparison with T6 and T73. Moreover, RRA treatment is only applicable for thin sections due to the short period of heating at retrogression. Recently, Song [14] found double hardness peaks during long-term aging process of a 7175 Al alloy and the strengths of two peaks were almost equal, but his work did not verify the universality of double peaks phenomenon among all 7000 series Al alloys.

The SCC of Al–Zn–Mg–Cu alloys has been the subject of extensive studies for many years, some consensuses have been reached for the operative mechanism(s), for example, the hydrogen embrittlement mechanism [4,15,16]. One primary agreement by a majority of researchers is that the SCC is closely related to the microstructure evolution during aging treatment. Previous studies [17,18] had reported the precipitation sequence of 7000 series Al alloy i.e. main α-supersaturated solid solution, GP zones and then to form η’ phase. The majority of η’ phase is a semi-coherent metastable MgZn₂, which will be transformed into incoherent stable η phase. Some researchers suggested that the SCC probably be linked to η phase and alloying elements precipitated on grain boundary. Wang's [19] research suggests that the second peak of a 7N01 alloy during long-term aging showing not only equivalent strength with T6 but also excellent SCC resistance, but the phase transformation and SCC mechanism are not clear.

The aim of this paper is to further research the aging strengthening behaviour of double peaks aging process by mechanical properties measurement coupling with energy spectroscopy analysis on a 7075 Al alloy. Meanwhile, the microstructure evolution and its relationship with strengthening effect and SCC also has been investigated through scanning electron microscope (SEM), transmission electron microscope (TEM) and high-resolution transmission electron microscopy (HRTEM).

Experimental procedure

Materials and heat treatment method

The 7075 high-strength aluminium alloys used were 40 mm thick rolled plates, which were supplied by U.S. ALCOA. The chemical compositions are shown in Table 1. According to previous works, the best solid solution temperature for 7000 series Al alloys is about 743 K. 7000 series alloys under this temperature can obtain sufficiently dissolved intermetallic particles and avoid coarsened grain and sub-grain under lower temperature, meanwhile can also avoid recrystallised-grain or even over burning under higher temperature. Detailed solid solution and aging methods see Table 2.

Table 1

Chemical compositions of 7075 Al alloy, wt-%.

Alloy	Zn	Mg	Cu	Zr	Ti	Mn	Cr	Fe	Si	Al
7075	6.02	2.20	1.55	–	0.20	0.30	0.25	0.50	0.40	bal.

Table 2

Heat treatment methods for 7075 Al alloy.

Alloys	Heat treatment*
7075	743 K/70 min + water quench+383 K/∼h
	743 K/70 min + water quench+393 K/∼h
	743 K/70 min + water quench+413 K/∼h
	743 K/70 min + water quench+433 K/∼h
	743 K/70 min + water quench+463 K/∼h

*Symbol∼ means a variable of aging hours.

Experimental methods

Rockwell hardness of Al alloys under various heat treatments was measured using a hardness tester (model name: HR-150) along the middle region of the cross section and an interval between two points of 1 mm, the load applied was 98 N. Tensile test was carried out on an AG-10TA model test machine.

The double cantilever beam (DCB) test was conducted in accordance with GB 12 445.1 [20] in a neutral 3.5 wt-% chloride solution, The sampling direction of DCB specimens was parallel to the S-L direction on rolled plates with a size of 26 mm × 26 mm × 127 mm (height × width × length). Before testing, the specimens were polished using emery paper of increasing grain size from 400 up to 1200 grit. A 2.5-4.0 mm pre-crack was made by a wrench for each specimen, then the specimen was sealed by transparent polyester tapes except the head and the screw which were sealed by wax (mixed by 50% paraffin wax and 50% rosin). A reading microscope was used to periodically measure the length of crack. The dimensions of tensile test and DCB test specimens were shown in Figure 1.

Phase analysis was performed by a D8 Discover X-ray diffractometer (produced by Bruker Company) with the scanning range changing from 30° to 80°. The specimens for XRD analysis were 10 × 10 × 4 mm³ in size. Micro-fractography, microstructure and phase structure were characterised using SEM (Hitachi S-4700), TEM (JEM-2010) equipped with HRTEM (JEM-2010HR). The thin foils for TEM observation were prepared by twin jet-polishing in 30% HNO₃, 70% ethanol solution cooled to −35°C with liquid nitrogen at 19 V. An Energy Dispersive X-ray Spectrometer (JEM-2010) was used to analysis the composition on and near grain boundaries.

It should be noted that the tests were duplicated for reproducibility. The same features observed in SEM, TEM and HRTEM for the first batch were found in duplicate. Differences in the amount of each feature on the surface were attributed to small differences.

Results

Hardness and strength during long-term aging

Figure 2 illustrates the aging hardening behaviour of 7075 Al alloy during long-term aging under various aging temperatures. From Figure 2, two obvious aging hardening peaks are observed, although with the aging temperature changing, there are some differences in the hardness values and the aging time when the second peak hardness appears. This in combination with literatures [14] and [19], demonstrate the universality of double-peaking aging phenomenon of 7000 series high-strength aluminium alloy. It needs to be noted that when the aging temperature reaches a critical point, the second peak will disappear and the hardness at the only first peak is much lower than those at other aging temperatures. For 7075 Al alloy, the second peak disappears when the aging temperature reaches about 463 K. Besides, the time both at first and second peak will be advanced when the aging temperature increases (not higher than critical temperature), which means the double peak aging is sensitive to the variation of aging temperature. On the other hand, the aging hardening effect of the second peak aging is generally better than that of the first peak aging, as shown in Figure 2, the hardness of 7075 Aluminium alloy under second peak aging increases approximately 5% compared to its first peak (T6) during 743 K/70 min+393 K/∼h treatment. In addition, 743 K/70 min+393 K/∼h (named ‘Aging 1’) exhibits a highest hardening effect.

The result of tensile tests of aluminium alloys during ‘Aging 1’ treatment in Table 3 shows a similar changing law of strength when comparing with Figure 2, which indicates once again the effect of double peaks strengthening of 7000 series aluminium alloy. The second peak of 7075 alloy is observed at 128 h, during which both yield strength (YS) and ultimate tensile strength (UTS) reach their maximum values.

Table 3

Tensile properties of 7075 Al alloy during ‘Aging 1’.

Al alloy	Heat treatment methods	YS/MPa	UTS/MPa	Elongation/%
7075	743 K × 70 min+393 K × 16 h	470	547	15.1
	743 K × 70 min+393 K × 36 h	504	579	14.4
	743 K × 70 min+393 K × 78 h	491	551	13
	743 K × 70 min+393 K × 128 h	515	580	14.8
	743 K × 70 min+393 K × 144 h	476	537	15.3

XRD Analysis and lattice constant calculation

Since the 7075 Al alloy with ‘Aging 1’ treatment (four aging states: under aging (16 h), first peak aging (36 h), second peak aging (128 h) and over aging (144 h), respectively) shows the best comprehensive mechanical properties, this aging treatment method was chosen for following tests. XRD patterns of the processed specimens are reported in Figure 3. As shown in Figure 3, all XRD patterns display strong peaks of α-Al matrix, and several weaker peaks of precipitates. Very few η’ (MgZn₂) nucleates after under aging and first peak aging treatment (Figure 3(a), (b)), then the intensity of η’ (MgZn₂) increases with increasing the aging time, thereby indicating the accelerated production of precipitate particles. During second peak aging treatment (Figure 3(c)), partial η’ (MgZn₂) begins to transform into equilibrium η (MgZn₂) and this process will be fully completed after 144 h aging treatment (Figure 3(d)) as there almost no signals of η’ phase. According to previous work [21], η (MgZn₂) is incoherent with matrix and has no strengthening effect.

XRD is not sensitive enough to detect phases less than 1 wt-%, the change of lattice constants during aging treatment is, therefore employed to identify other precipitated particles. The lattice constant results of α-Al matrix of 7075 Aluminium alloy under various aging hours are summarised in Table 4. According to Pearson's research [22], the precipitation of Mg and Zn atoms can be reflected by the change of lattice constants as this precipitation process can result in lattice distortion of crystals. The lattice constant of α-Al matrix will increase during an Mg atom-dominated precipitation process. On the contrary, it will decrease during a Zn atom-dominated precipitation process. As shown in Table 4, the lattice constant increases when the aging time increases from 16 to 36 h, suggesting that the first peak is strengthened by Zn enriched GP zones. Furthermore, more Mg atoms precipitated when the aging time increases to 78 h, which creates a good environment for the formation of MgZn₂.

Table 4

Lattice constants of α-Al matrix of 7075 Aluminium alloy during ‘aging 1’ treatment.

condition	16 h	36 h	78 h	128 h	144 h
a/nm	0.404900	0.405725	0.404933	0.405203	0.405305
Δa*/nm
0.000825	−0.000792	0.000270	0.000102

*Δa is the difference value of two adjacent lattice constants.

Fracture surface of tested tensile specimens

Figure 4 shows the fracture surface of ‘Aging 1’ treated 7075 Al specimens after tensile tests. Some dimples mixed with cleavage facets can be found on fracture surfaces of specimens at both first and second aging peaks (Figure 4(b,c)), however for the second peak, the dimples are deeper. Additionally, for underaged treatment (Figure 4(a)), fracture surfaces produced by tensile test exhibit typical intergranular fracture. In contrast, the fracture surface of overaged specimens (144 h, see Figure 4(d)) is dominated by dimples which indicating a better plasticity however the strength decreases a lot in comparison with the two aging peaks.

Stress corrosion testing

The relationship between crack propagation rate (da/dt) and stress intensity factor (K) of DCB specimens during immersion in 3.5 wt-% NaCl solution was evaluated in accordance with following equation [20]: (1) where δ represents the displacement on loading line, H is the half height of the specimen, E denotes Young's modulus, which is 70 560 MPa for 7075 aluminium alloy and will hardly be affected by heat treatment. The greater the crack propagation rate (da/dt) or the less the SCC critical stress intensity factor (K _ISCC). the higher the SCC susceptibility is. Figure 5 shows the da/dt versus K _I curves of 7075 Al alloy under ‘Aging 1’ treatment. From Figure 5, the crack propagation can be divided into two stages, one is described as fast rising stage (stage I), and another one is plateau stage (stage II), the crack propagation rate keeps unchanged during this stage (named as plateau rate). The plateau rates of stress corrosion cracks for specimens with various aging states after immersed in 3.5 wt-% NaCl solution are summarised in Table 5, in which the plateau rate of DCB specimens decreases with increasing the aging hours, however the K _ISCC is on the contrary. According to Figure 5 and Table 5, the SCC resistance of DCB specimens with the second peak aging is much better than that of DCB specimens with the first peak aging.

Table 5

SCC behaviour of 7075 Al alloy when immersed in 3.5 wt-% NaCl solution.

Aging treatments	Plateau rates (mm s⁻¹)	KI_SCC (MPam^1/2)
393 K × 16 h	2.34 × 10⁻⁵	5.6
393 K × 36 h	6.24 × 10⁻⁶	12.5
393 K × 78 h	5.23 × 10⁻⁷	14.3
393 K × 128 h	9.15 × 10⁻⁸	17.5
393 K × 144 h	7.14 × 10⁻⁸	20.1

When da/dt reached 10⁻⁹ m s⁻¹ after 5760 h immersion in 3.5 wt-% NaCl solution, the test was stopped, then the specimens were loaded to fracture to analyse the fracture surface. The macroscopic fracture surfaces of specimens with different aging states are shown in Figure 6. The length of propagated crack decreases with increasing the aging hours which is in compliance with the results in Figure 5. In addition, the SEM micrographs of specimens with two aging peaks are presented in Figure 7, which are significantly different from those of specimens after tensile tests (see Figure 4). As shown in Figure 7, corrosion products can be seen on almost all the fracture surfaces. For the first peak aged specimen, the local direction of crack growth at crack tip region (Figure 7(a)) is very tortuous, indicating severe stress corrosion. Meanwhile, many secondary cracks and intergranular cracks are produced at crack propagation region (Figure 7(b)), this probably because of the preferential deformation of the precipitate-free zone (PFZ) around the grain boundary area [22]. In terms of the second peak aged specimen, the fracture surface produced by DCB test in 3.5 wt-% NaCl solution exhibits predominant cleavage river pattern coupled with a small quantity of intergranular facets in crack tip region (Figure 7(c)). Moreover, there are very few secondary cracks in the crack propagation region (Figure 7(d)). As a result, the second peak aging treated 7075 Al alloy can achieve higher aging strengthening effect without sacrificing too much performance of SCC resistance in comparison with first peak aging treated one. This means the double peak aging treatment of 7000 series aluminium alloys has the potential for industrial application.

Microstructure comparison between the double peaks

The microstructure of 7075 Al alloy in grain and on grain boundary is characterised by TEM, corresponding results are shown in Figure 8. A large volume fraction of finely distributed matrix precipitates (MPt) can be seen in specimens after first peak aging treatment which mainly are GP zones (Figure 8(a)), suggesting that the hardening effect at first peak is mainly attributed to deformable GP zones. Similar results have been reported before [19]. On the other hand, GP zones are coherent with Al matrix and existing without fixed structure, this is why they cannot be clearly detected by XRD. The precipitates on the grain boundaries of the first peak aged specimens are small and distributed continuously (Figure 8(b)). Concerning the specimens after second peak aging, fine and dispersive η’ (MgZn₂) combined with small amount of GP zones are distributed in grain and sub-grain (Figure 8(c)), but the size of η’ particle is slightly greater than GP zones and in round or short bar shape. Furthermore, η’ particles are always easy to distinguish from α-Al matrix, and are hexagonal crystallographic microstructure. Meanwhile, as shown in Figure 8(d), PFZ along the grain boundary is wider than those in Figure 8(b). Since grain boundaries are preferential paths for SCC which has been proved by the intergranular corrosion under stress in chloride environment, PFZ can act as a barrier to hinder the forming of SCC path. Therefore, wider PFZ is considered to be beneficial to SCC resistance [23]. In addition, Coarsened η (MgZn₂) is mainly precipitated on grain boundary which is incoherent with α-Al matrix.

HRTEM is employed to further understand the microstructure and phase transformation of 7075 Al alloy under double peaks aging treatment. Fine and round shape GP zones are precipitated in grain after first peak aging with sizes between 2 and 6 nm, as shown in Figure 9(a) and (b). Moreover, in Figure 9(b), the GP zones present an alternate layer structure which probably due to the opposite-size effect between Mg and Zn atom layers. In addition, all the GP zones are coherent with the [100] and [111] zone axes near α-Al matrix and their inter-particle distances are about 4-10 nm which are shorter than the distance between each two [100] zone axis. Therefore, there should exist a lattice distortion band transiting from Al matrix to GP zones. For the second peak, the incoherent relationship between precipitates and [100] zone axis near α-Al matrix proves the presence of η’ (MgZn₂) as some η’ (MgZn₂) is semi-coherent with Al matrix (round shape, Figure 9(c)). The short bar shaped η’ (MgZn₂) is also a nanoscale particle (about 4-10 nm) and coherent with [111] zone axis (see Figure 9(d)).

Figure 10 shows the HRTEM images of grain boundary precipitates (GBPs) of 7075 Al alloy under double peaks aging treatment. As shown in Figure 10, the GBPs of specimens under first and second peak aging treatment are both η (MgZn₂). They are about 6-10 times coarser than the MPts in Figure 9 as there is more energy on grain boundary that will induce the preferential nucleation of particles.

Discussion

Strengthening mechanism of Al alloys under double peaks aging treatment

It is generally accepted that the hardening and strengthening effect of Al–Zn–Mg–Cu alloy are associated with volumetric fraction and composition of precipitates produced in grain during aging treatment [24,25]. In present work, the precipitates and microstructure of Al–Zn–Mg–Cu alloy under double peaks aging treatment are significantly different. For ease of understanding, the microstructure evolution process is summarised in Figure 11 according to the results of Figures 2, 3 and 8. As a result, the precipitates of specimens with first peak aging are almost fine and intensive GP zones. They are homogeneously distributed in grain and fully coherent with Al matrix. Previous experimental data and literatures [26] reported that a large number of vacancies would be generated in grain after solid solution which is the beneficial condition for the nucleation of GP zones. Meanwhile, for GP zones, the activation energy of nucleation was higher than that of nuclear growth. This probably the reason why the GP zones precipitated during first peak aging are in fine size and distributed such densely. However, with aging process continuing, GP zones will gradually grow up and transform into η’ (MgZn₂). Short inter-particle distance of GP zones will induce lattice distortion bands between Al matrix and GP zones so that more energy will be consumed when dislocations shearing the GP zones. In other words, dislocation glide is hindered which reflects the strengthening effect at the macro level.

The main precipitate in grain for the second peak is metastable η’ which is harder than GP zone, so dislocation cannot shear but bypass it through the Orowan mechanism [19]. When a dislocation bypassing the η’ particles, a line tension must be overcome which depends on the distance between two η’ particles. The shorter the distance is, the higher the line tension must be overcome. Meanwhile, the theoretical distance between two η’ particles of 7075 Al alloy under the second peak aging can be calculated by the equation of critical shear stress: (2) where τ is the critical shear stress which is equal to half the UTS of specimens under second peak aging (see Table 3); b is the Burgers vector which is equal to for a fcc Al alloy; L is the theoretical inter-particle distance of η’; T is the line tensile; G is the shear modulus of 7075 Al alloy. According to our experimental data, τ = 310 MPa; b = 2.81485 Å; G = 23.8 × 10³ MPa; substituting them into Equation (2) and yield L = 21.6 nm. As shown in Figure 9(d), the distance between coherent η’ is about 10-40 nm which is similar to the value of L. This suggests that η’ and its distribution play a critical role in aging strengthening at the second peak.

SCC mechanism of Al alloys under double peaks aging treatment

Numerous researches [3,4,27 29] have been conducted on the SCC of 7000 series Al alloys, some conclusions have been made on a micro level, including the effect of microstructure, alloy element and grain boundary on SCC behaviour. Song et al. [30] reported that the type and composition of grain boundaries are closely related to the SCC behaviour of 7000 series Al alloys as SCC often starts from grain boundary, meanwhile, the Mg segregation on grain boundaries could increase the amount of hydrogen absorbed and consequently accelerated the hydrogen embrittlement. Birbilis and Buchheit [31] suggested that increasing the copper content of GBP would result in reducing both anodic-reaction rates and cathodic generation of hydrogen at crack-tips. Both Song and Birbilis's results are consistent with our EDS results in Table 6 which indicating both Mg and Cu segregated towards grain boundary. The amount of Mg segregation at first peak is higher than that of second peak, but for Cu segregation, it shows a significantly reverse result. The sampling positions of SEM/EDS analysis for the two aging peaks are shown in Figure 12. It should be noted that the SEM/EDS was used as a tool in this paper to qualitatively analyse the trend of element segregation on or near grain boundary. If assuming that almost all Zn atoms on grain boundary are present in the form of η phase (MgZn₂), then the excess magnesium is free on grain boundary. Figure 13 shows the relationship between excess Mg and aging time, in which the content of excess Mg decreases with increasing the aging time. Therefore, the close correlation of Mg segregation on gain boundary with SCC resistance is further proved, although this is not entirely true it serves as an assumption to calculate excess Mg. On the other hand, as shown in Figure 8, the width of PFZ seems also play a role in SCC behaviour. According to Deng et al. [32], the corrosion potential of MgZn₂ on grain boundary is more negative than Al matrix which will act as the anode when forming a micro battery with Al matrix. The existence of wider PFZ can decrease the forming velocity of micro battery, however, this explanation is only suitable for anodic dissolution dominated SCC.

Table 6

Composition of grain boundaries of 7075 Al alloy analysed by SEM/EDS.

Heat treatment	Element	Analysis points* (at.-%)
		1	2	3	4	5
First peak aging (36 h)	Mg	2.95	3.25	3.21	3.11	2.91
	Zn	2.79	3.62	3.68	3.64	2.89
	Cu	0.08	0.18	0.30	0.23	0.09
Second peak aging (128 h)	Mg	2.52	2.90	2.95	2.95	2.49
	Zn	3.67	3.80	3.94	3.87	2.95
	Cu	0.10	0.24	0.35	0.28	0.11

*Point 3 is mostly near the grain boundary, other four points gradually leave away the grain boundary.

Conclusions

The strengthening effect of 7075 Al alloy under the second aging peak is equal or even greater than that under the first aging peak.

The strengthening mechanism of Al–Zn–Mg–Cu Aluminium alloys under first peak depends on GP zones as more energy will be consumed when dislocations shearing the GP zones. In terms of second peak, the strengthening effect relies on dispersive η’ (MgZn₂).

The Mg and Cu segregation may contribute to the SCC of Al–Zn–Mg–Cu Aluminium alloys. The SCC resistance will increase with the decreasing of the segregation amount of Mg and the increasing of the segregation amount of Cu. The width of PFZ also probably plays a role in SCC.

Footnotes

Acknowledgements

The financial aids of the National Natural Science Foundation of China under grant number 51871031 and the Priority Academic Program Development of Jiangsu Higher Education Institutions (PAPD) are gratefully acknowledged.

Disclosure statement

No potential conflict of interest was reported by the author(s).

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