Stress corrosion cracking susceptibility of Al–Mg alloy sheet with high Mg content

Journal of Materials Processing Technology 125–126 (2002) 275–280

Stress corrosion cracking susceptibility of Al–Mg alloy sheet with high Mg content Miljana Popovic´*, Endre Romhanji Department of Metal Forming, Faculty of Technology and Metallurgy, University of Belgrade, Karnegijeva 4, POB 35-03, 11120 Belgrade, Yugoslavia Received 4 December 2001; accepted 22 January 2002

Abstract Slow strain rate testing (SSRT) was used to study the effect of the microstructure on the stress corrosion cracking (SCC) susceptibility of Al– Mg alloy sheet containing 6.8 wt.% Mg. In the cold-rolled and fully annealed conditions, high SCC susceptibility was experienced. In those cases the ductility was strongly affected by the presence of corrosive environment (for hard temper: Elair ¼ 13:6%, ElSCC ¼ 0:6%; for annealed condition: Elair ¼ 24:1–25.3%, ElSCC ¼ 3:2–4.2%) and the elongation loss was great, Elloss ¼ 81:7–95.6%. It is supposed that the high SCC susceptibility results from a continuous network of the b-phase (Mg5Al8) precipitate at grain boundaries for the annealed temper, and heavy precipitation of b-phase along the planes of localized deformation for the hard temper. High SCC resistance attained after thermal exposure at the temperature range 225–285 8C (stabilized condition). The ductility was almost unaffected by the presence of corrosive environment (Elair ¼ 12:8–23.2%, ElSCC ¼ 12:8–22%) and the elongation loss was small, Elloss < 7%. High SCC resistance was related to the stabilized structure, which causes discontinuous b-phase (Mg5Al8) precipitation in a globular form, uniformly distributed throughout the structure. # 2002 Published by Elsevier Science B.V. Keywords: Al–Mg alloy; Stress corrosion cracking (SCC); Slow strain rate testing (SSRT)

1. Introduction Wrought alloys of the 5XXX series have been considered in the past for use in a wide variety applications as in automotive body structures, trains, shipbuilding or cryogenic vessels, due to their excellent properties such as high strength, good formability, corrosion resistance and weldability [1–3]. In general, they offer the best combination of strength and corrosion resistance of all Al-alloys. Al–Mg alloys are not heat treatable and they achieve high strength through solution hardening by the Mg atoms retained in solid solution, dispersion hardening by precipitates (Mn, Cr, Zr—bearing ones) and strain hardening effects [1,2]. The addition of Mg markedly increases the strength of these alloys without a significant decrease of ductility [1]. On the other hand, Al–Mg alloys containing more than 3 wt.% Mg become susceptible to stress corrosion cracking (SCC) due to supersaturation of solid solution and increased tendency of Mg atoms to precipitate at grain boundaries as highly anodic b-phase (Mg5Al8) [3–6]. The precipitation is accelerated by heating at elevated temperature (80–160 8C) or after very long aging at room temperature. *

Corresponding author. E-mail address: [email protected] (M. Popovic´). 0924-0136/02/$ – see front matter # 2002 Published by Elsevier Science B.V. PII: S 0 9 2 4 - 0 1 3 6 ( 0 2 ) 0 0 3 9 8 - 9

Many of the recent studies [3,4,7,8] are related to the question of how to minimize or avoid the SCC susceptibility of highly alloyed Al–Mg alloys. It was reported [7,9,10] that Zn addition of 1–2 wt.% improve the SCC resistance of Al–Mg alloys, due to the formation of stable ternary Al–Zn–Mg phase (t-phase) at grain boundaries. The application of protective coatings, such as 7072 alloy for aluminium alloys, is another very useful and important method in avoiding SCC susceptibility [3,4]. Some works [4,5,11,12] have indicated a correlation between the microstructure and SCC susceptibility, i.e. the possible control of the SCC by applying specific TMTs. The aim of the present paper was to examine and discuss various TMTs or the influence of appropriate microstructures on the SCC behaviour of highly alloyed Al–Mg type alloy sheet. 2. Experimental 2.1. Material This study is performed using a 3 mm thick Al–Mg alloy sheet in a fully annealed condition (O temper) with the chemical composition given in Table 1.

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Table 1 Chemical composition of the tested alloy (wt.%) Mg Mn Fe Si Cu Zn Ti Cr

6.8 0.51 0.2 0.1 0.01 0.03 0.05 0.01

The as-received material was further cold-rolled to a thickness of 1.8 mm, with intermediate annealing at 320 8C for 3 h, and cold-rolled to a final thickness of 0.9 mm. After rolling, the material was subjected to various thermal exposures at 225, 245, 265 and 285 8C for 3, 6 and 12 h, at 320 8C for 30 min and 3 h, and finally at 450 8C for 10 min. Both the as-cold-rolled and heat-treated specimens have undergone sensitizing treatment for 7 days at 100 8C. 2.2. Tensile testing Uniaxial tensile testing was performed using ‘‘Zwick 1484’’ tensile testing machine, on a specimens taken transverse to the rolling direction (L–T orientation), with a gauge length of 25 mm and the width of 12.5 mm. The tensile properties of as-cold-rolled and various heat-treated specimens were tested at room temperature with an initial strain rate of 6:7  103 /s.

(Elloss) calculated according to the equation: Elloss ¼ ½1  ðElSCC =Elair Þ  100 ð%Þ.

3. Results 3.1. Tensile properties The effects of cold working and different thermal exposures on transverse tensile properties of the tested alloy are shown in Fig. 1(a) and (b). Annealing of the cold-rolled samples in the temperature range 225–450 8C is followed by decrease of YS and UTS, and by increase of total elongation, as it is shown in Fig. 1(a) and (b). Fig. 1(a) shows that the basic softening process is limited to the temperature interval of 225–285 8C. At higher temperatures the YS and UTS values are rather unaffected by the temperature increase. Increasing the exposure time from 3 to 6 h or 12 h is followed with the drop of strength parameters, which is the highest at 265 8C: 35% for the YS and 10% for UTS (Fig. 1(a)). The total elongation appeared to increase continuously up to 320 8C independently of the exposure time (Fig. 1(b)) and after that is rather unaffected with the temperature change.

2.3. Optical microscopy Samples for optically metallographic observation, of cold-rolled and heat-treated specimens, were prepared by electrochemical polishing and etching in Barker’s reagent. To identify the b-phase precipitates, sensitized specimens were mechanically polished and etched in 10% orthophosphoric acid at 50 8C for 2.5 min. 2.4. Slow strain rate testing (SSRT) The SCC susceptibility was evaluated by SSRT on the specimens (both as-cold-rolled and heat-treated) that had undergone sensitization treatment for 7 days at 100 8C. Tension specimens with a gauge length of 25 mm and the width of 6.25 mm were taken transverse to the rolling direction (L–T orientation). In all cases the SSRT was conducted in triplicate in dry air (at an initial strain rate of 6:7  104 /s) and in 2% NaCl þ 0:5% Na2CrO4 solution as a corrosive environment (at an initial strain rate of 3:3  105 /s). The elongation of each fractured specimen was measured as the length between gauge marks and it was expressed as a percentage of the original gauge length. Relative elongation in the corrosive environment (ElSCC) was compared with that in dry air (Elair), and the SCC susceptibility was determined by the elongation loss

Fig. 1. Effect of annealing conditions: (a) yield stress (YS), ultimate tensile strength (UTS); (b) total elongation (El).

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Table 2 Transverse tensile properties of the specimens before (YS, UTS, El) and after sensitization (YSS, UTSS, ElS) t (8C)

t (h)

YS (MPa)

UTS (MPa)

El (%)

YSS (MPa)

UTSS (MPa)

ElS (%)

225

3 6 12

265.5 247.8 253.1

416.3 396.7 389.3

17.8 16.4 17.5

300.2 282.3 288.6

423.7 402.4 395.2

13.4 16 16.2

245

3 6 12

273.6 248.6 246.1

399.5 386.5 379

19.5 18 18.7

277.4 266.6 268.4

407.9 391.6 386.4

17.6 16.2 18.4

265

3 6 12

265.3 164 156.8

393.5 359.5 346.5

20.1 21.1 21.3

262.8 175.2 151

398 361 345

22.4 21.8 24.4

285

3 6 12

170.5 161.1 160.4

356.3 343.1 341.1

23.9 28.1 26.3

168.7 159.8 156

363 345 341.3

27.4 29.6 28.6

320

0.5 3

166 167.8

343.9 343.5

29.5 30.8

165.2 166.1

344.1 345.9

32 30.4

450

10min

172.2

345.2

28.8

170.5

346.2

31.6

381.8

464.2

8

289.4

420.2

12.2

As-cold-rolled

The effect of sensitizing treatment on the tensile properties (YS, UTS and El) of cold-rolled and annealed specimens is given in Table 2. Sensitization of cold-rolled specimen was followed by decrease in tensile strength and increase in ductility. The strength parameters were slightly higher after sensitization of specimens annealed from 225 to 265 8C, while those were slightly lower after sensitizing treatment of specimens annealed from 285 to 450 8C. 3.2. Optical microscopy 3.2.1. Microstructure after cold rolling and heat treatment Microstructural examination has shown markedly elongated grains in the rolling direction for the as-cold-rolled sample (Fig. 2(a)). After annealing at 225 8C for 3 h, the grain structure was still highly elongated in the rolling direction (Fig. 2(b)). Such structures were also observed after annealing at 225 8C (for 6, 12 h), 245 8C (for 3, 6 and 12 h) and 265 8C for 3 h. More or less homogenous structures were obtained after annealing at 265 8C for 6 and 12 h, at 285 8C for 3, 6 and 12 h, at 320 8C for 30 min and 3 h, and at 450 8C for 10 min. Such a structure shown in Fig. 2(c) was obtained from a sample annealed at 320 8C for 3 h. 3.2.2. Microstructure after sensitization Microstructural examinations were also carried out to determine the effect of the microstructure on the b-phase precipitation during sensitization at 100 8C for 7 days. The effect of sensitization on b-phase precipitation in as-cold-rolled samples is shown in Fig. 3(a). The micrograph shows heavy precipitation inside the elongated grains and also at grain boundaries. It is obvious that there is a pronounced precipitation along the planes of localized

Fig. 2. Microstructures of the tested alloy after: (a) cold rolling with 50% reduction; (b) heat treatment at 225 8C for 3 h; (c) annealing at 320 8C for 3 h.

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for 3 h before sensitization (Fig. 3(c)). Precipitation of b-phase inside the grains is less evident. Similar b-phase distributions observed in the samples heat-treated at 285 8C for 6 and 12 h, at 320 8C for 30 min, and at 450 8C for 10 min, before sensitizing treatment. 3.3. SSRT The SSRT of SCC susceptibility was carried out on specimens sensitized for 7 days at 100 8C. The effect of cold working (50% reduction) and different thermal exposures (from 225 to 450 8C) before sensitization on the SCC susceptibility, was evaluated by the elongation loss (Elloss) calculated using the equation:    ElSCC Elloss ¼ 1   100 ð%Þ Elair Total elongation obtained by SSRT in dry air, Elair, and in 2% NaCl þ 0:5% Na2CrO4 as corrosive environment, ElSCC, are given in Table 3. It can be seen that the ductility was strongly affected by the corrosive environment in the case of cold-rolled or specimens annealed at 320 and 450 8C before sensitizing treatment. Total elongation decreased from Elair ¼ 13:6% in dry air to ElSCC ¼ 0:6% in the corrosive environment, for deformed specimen, and from Elair ¼ 24:1–25.3% to ElSCC ¼ 3:2–4.4% for the annealed specimens (Table 3). For the specimens heat-treated at 225, 245, 265 and 285 8C for 3, 6 and 12 h, before sensitization, the ductility was almost unchanged by corrosive environment. Total elongations in dry air Elair ¼ 12:8–23.2% were approximately the same as total elongations in corrosive environment ElSCC ¼ 12:8–22% (Table 3). Table 3 Slow strain rate tested total elongation in dry air (Elair), in 2% NaCl (ElSCC), and elongation loss (Elloss) for various specimens after sensitization

Fig. 3. Microstructures of tested Al–Mg alloy after sensitization treatment for 7 days at 100 8C with varying degrees of susceptibility to SCC: (a) cold-rolled with 50% reduction—highly SCC susceptible; (b) cold-rolled 50%, heat-treated at 245 8C for 12 h—highly SCC resistant; (c) cold-rolled 50%, heat-treated at 320 8C for 3 h—highly SCC susceptible.

deformation (slip planes and shear bands). The micrograph on Fig. 3(b), corresponding to the sample annealed at 245 8C for 12 h before sensitization, indicates the discontinuous bphase precipitates in a globular form, throughout the structure (inside the grains and at grain boundaries). The same distribution of b-phase precipitates was observed in the samples heat-treated at 225, 245 and 2658C for 3, 6 and 12 h, and also at 285 8C for 3 h, before sensitizing treatment. A continuous network of b-phase precipitates at grain boundaries was revealed in the sample annealed at 320 8C

t (8C)

t (h)

Elair (%)

ElSCC (%)

225

3 6 12

13.3 14.4 12.8

13.1 14 12.8

1.8 2.8 0

245

3 6 12

16.4 17.7 15.3

16 16.5 15.3

2.4 6.8 0

265

3 6 12

17.1 20.3 19.6

17.1 19 20

0 6.3 0

285

3 6 12

20.4 23.2 23.1

20.3 21.7 22

0.6 6.9 4.7

Elloss (%)

320

0.5 3

25.3 24.1

4.2 4.4

83.4 81.7

450

10min

24.8

3.2

87.1

13.6

0.6

95.6

As-cold-rolled

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which is non-coherent with the matrix [2,5]. Such changes are in a good agreement with some previously published results about recrystallization in a highly alloyed Al–Mg6 type sheet [1] and the decomposition of a-solid solution [5]. 4.2. Microstructure and tensile properties after sensitizing treatment

Fig. 4. The elongation percentage loss, Elloss, for the various specimens after sensitization.

Table 3 and Fig. 4 clearly indicate that the elongation loss (Elloss) for the deformed and at 320 or 450 8C annealed specimens was much larger (Elloss ¼ 81:7–95.6%) than in the case of specimens heat-treated in the temperature range 225–285 8C (Elloss ¼ 0–6.9%).

4. Discussion 4.1. Tensile properties with different TMTs Variations of the tensile properties (YS, UTS, El) shown in Fig. 1(a) and (b) were affected by microstructure development during different thermal exposure. A decrease in tensile strength and an increase in ductility, with raising the temperature from 225 to 450 8C, were accomplished by recovery and recrystallization processes, and also by a reduction of the concentration of Mg-atoms in solid solution. On the other hand, Al–Mg alloy with 6.8 wt.% Mg would be dual phase (a-Al þ b-phase) below 285 8C (from 225 to 285 8C) and single phase (a-Al) above 285 8C (at 320 and 450 8C) [2]. The high tensile strengths after annealing at 225, 245 8C (for 3, 6 and 12 h), or at 265 8C for 3 h (Fig. 1(a)), were related to the recovered structures with more or less elongated grains in the rolling direction (Fig. 2(b)), and also to the precipitation of b0 -phase, which is partly coherent with the matrix, as it was observed in other Al–Mg alloys [2,13,14]. Lower strength approached after annealing for 6 and 12 h at 225 and 245 8C can be attributed to more recovered structure, the transition of b0 -phase to b-phase, and the coarsening of b-phase precipitates, as it was published earlier [2,5,13]. The decrease in tensile strength observed after annealing at 265 8C for 6 and 12 h, and also in the temperature interval from 285 to 450 8C (Fig. 1(a)), was related to the completely recrystallized grain structure (Fig. 2(c)), and also to the formation and coarsening of b-equilibrium phase precipitate

During sensitization of cold-rolled or heat-treated specimens, for 7 days at 100 8C, the transformation of supersaturated a-solid solution has occurred. As it was reported earlier [2,5,15,16], the decomposition of solid solution in Al–Mg alloys can be followed by direct formation of equilibrium b-phase precipitates, as the a ! b process, or through the formation of partly coherent b0 -phase, as the a ! b0 ! b process. Since different structural heterogeneities facilitate the formation of b0 -phase precipitates [14–16], it was supposed that the sensitization of samples annealed at t ¼ 225–285 8C produced b0 -phase precipitates. In the case of cold-rolled or samples annealed at 320 and 450 8C, sensitizing treatment resulted in equilibrium b-phase precipitation. Fig. 3(a)–(c) have shown that the form and distribution of b- or b0 -phase precipitates were affected by TMTs before sensitization. During sensitization of samples annealed at t ¼ 225–285 8C, a precipitation of globular b0 -phase took place throughout the structure (Fig. 3(b)). It is supposed that random distribution of discontinuous b0 -phase precipitates was provided by the presence of b- or b0 -phase precipitates in dual phase structure before sensitization, and also by a certain amount of dislocations in recovered or non-fully recrystallized structure, which could make a favourable place for b0 -phase precipitation during sensitization. Such a behaviour was also observed in other Al–Mg alloys [3,5,13,15]. During sensitization of samples annealed at 320 8C or at 450 8C, with fully recrystallized structure, a continuously distributed b-phase precipitation layer formed at grain boundaries (Fig. 3(c)). Sensitization of cold-rolled specimens induced a pronounced precipitation of b-phase on the grain boundaries and also within the grains, in the areas of localized deformation, which corresponded to deformation bands or shear bands (Fig. 3(a)). The presented results have shown that tensile properties after sensitizing treatment were affected by previous TMTs (Table 2). Lower tensile strength and higher ductility of coldrolled specimens after sensitization (Table 2) were attributed to a dislocation rearrangement due to recovery processes and to the formation of b-equilibrium phase precipitates. An improved tensile strength of specimens annealed at t ¼ 225–265 8C and sensitized at 100 8C, in comparison with those before sensitization (Table 2), can be related to the formation of b0 -phase precipitate, partly coherent with a-Al matrix. The tensile strength of a specimens annealed at higher temperatures (from 285 to 450 8C), almost unaffected or slightly lower after sensitizing treatment (Table 2), was

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related to the precipitation of b-equilibrium phase, noncoherent with a-Al matrix. 4.3. SCC susceptibility Experimental results of the SSRT have shown a serious degradation of ductility for cold-rolled sample in corrosive environment, ElSCC ¼ 0:6%, in comparison with ductility in dry air, Elair ¼ 13:6% (Table 3). The elongation loss was great Elloss ¼ 95:6% and a high SCC susceptibility was observed (Table 3, Fig. 4). According to the theories [17] about SCC in Al-alloys, it was supposed that high SCC susceptibility of hard temper was due to anodic dissolution of a very concentrated b-phase precipitates (Fig. 3(a)), and accelerated cracking along the planes of localized deformation. Results of the SSRT for the samples annealed at 225, 245, 265 and 285 8C for 3, 6 and 12 h, before sensitization, indicated a slight difference between total elongation in dry air, Elair ¼ 12:8–23.2%, and in a corrosive environment, ElSCC ¼ 12:8–22% (Table 3). The ductility, almost unchanged by the presence of a corrosive environment, and low values of the elongation loss, Elloss ¼ 0–6.8% (Table 3, Fig. 4), indicated high SCC resistance for stabilizing temper. This can be attributed to random distribution of b0 -phase precipitates throughout the structure (Fig. 3(b)) which retarded the corrosion cracking and failure propagation. The SSRT results for specimens annealed at 320 8C for 30 min and 3 h, or at 450 8C for 10 min, has shown very pronounced changes of ductility in a corrosive environment, ElSCC ¼ 3:2–4.4%, in comparison with ductility in dry air, Elair ¼ 24:1–25.3% (Table 3). The elongation loss was very high, Elloss ¼ 81:7–87.1% (Table 3, Fig. 4), and severe SCC susceptibility occurred. This can be attributed to anodic dissolution of a continuously distributed b-phase precipitation layer along grain boundaries (Fig. 3(c)), which induced and accelerated the failure propagation by intergranular mechanism.

5. Conclusions Cold rolling with 50% reduction results in elongated grain structure, high strength level (YS 380 MPa, UTS 465 MPa), low ductility (El 8%), and high SCC susceptibility (Elloss 96%) due to the dissolution of b-phase precipitates along the planes of localized deformation. Thermal exposures at t ¼ 225–285 8C for 3, 6 and 12 h produce good mechanical properties (YS 160–275 MPa, UTS 345–420 MPa, El 16–28%) and an improved SCC resistance ðElloss 7%Þ due to discontinuous b-phase precipitates in a globular form throughout the structure.

Thermal exposures at t ¼ 320 8C for 30 min and 3 h, and at 450 8C for 10 min result in a recrystallized grain structure, moderate strength level (YS 165–175 MPa, UTS 345 MPa), high ductility (El 30–32%) and high SCC susceptibility (Elloss 82–87%) due to continuous network of b-phase precipitate at grain boundaries. In order of increasing SCC resistance the tempers of the tested Al–Mg alloy with 6.8 wt.% Mg are as follows: (1) as-cold-rolled (hard temper); (2) annealed condition (t ¼ 320 and 450 8C); (3) stabilized temper (from t ¼ 225 to 285 8C).

Acknowledgements The authors are grateful to Aluminium Rolling Mill, Sevojno, for supplying the material for testing; Laboratory for Metallography and Mechanical Testing in Copper Rolling Mill, Sevojno, for help with microstructural examinations and tensile testing; Military Institute, Belgrade, for help with slow strain rate testing of stress corrosion cracking susceptibility.

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