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Transgranular environment-induced cracking of 7050 aluminium alloy under cyclic loading conditions at low frequencies
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International Journalof Fatigue
International Journal of Fatigue 30 (2008) 1827–1837
www.elsevier.com/locate/ijfatigue
Transgranular environment-induced cracking of 7050 aluminium alloy under cyclic loading conditions at low frequencies Reinhold Braun * DLR – German Aerospace Center, Institute of Materials Research, D-51170 Ko¨ln, Germany Received 25 October 2007; accepted 11 February 2008 Available online 19 February 2008
Abstract The environment sensitive cracking behaviour of alloy 7050 plate in different tempers was studied under cyclic loading conditions using short transverse oriented notched tension specimens immersed in substitute ocean water. The cyclic loading tests were carried out at displacement rates in the range from 2 106 to 2 104 mm s1 applying a sawtooth wave form. The maximum load corresponded to an initial stress of 600 MPa, except for alloy 7050-T7651 being loaded up to 500 MPa. Number of cycles to failure decreased with decreasing displacement rate. A linear relationship was found between log values of the number of cycles to failure and the loading frequency. Fractographic examinations of failed specimens revealed brittle transgranular fracture with crack arrest markings on flat facets. The sensitivity to transgranular environment-induced cracking proposed to be associated with hydrogen embrittlement did not depend on heat treatment conditions. Ó 2008 Elsevier Ltd. All rights reserved. Keywords: AA7050; Cyclic loading; Transgranular environment-induced cracking; Striations
1. Introduction Age-hardenable 7000-series aluminium alloys are sensitive to environmentally assisted cracking including stress corrosion cracking (SCC), hydrogen embrittlement and corrosion fatigue [1–8]. The SCC susceptibility of high strength aluminium alloys decreases from underaged to overaged microstructures [1–3]. Whereas Al–Zn–Mg–Cu alloys are quite sensitive to stress corrosion cracking in the peak-aged T6 temper, two-stage overaging improves the SCC resistance, resulting in virtual immunity for the T73 temper [4]. Stress corrosion cracking causing failure in commercial aluminium alloys is generally intergranular [1,2,4]. A similar behaviour was observed for hydrogenenvironment assisted cracking [9]. Using double-cantilever beam fracture mechanics specimens exposed to 90% relative humidity air, the stage II (plateau) crack growth rate *
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of alloy 7050 was found to decrease from the underaged to overaged conditions. In all tempers, the fracture path was intergranular, predominantly along high-angle grain boundaries. Water vapour was also found to promote corrosion fatigue. The fatigue resistance of a high purity Al–Zn– Mg alloy was significantly reduced when exposed to a humid nitrogen environment [10]. The degradation of fatigue resistance in water vapour was concluded to be associated with hydrogen embrittlement. Transgranular or mixed intergranular and transgranular cracking was usually observed with corrosion fatigue in aluminium alloys [7,10]. Whereas alloy 7010-T76 exhibited entirely ductile fatigue fracture in dry environments, failure in moist air occurred by cleavage with well defined surface striations [11]. Again, crack growth by cleavage was suggested to be the result of hydrogen embrittlement involving dislocation transport of hydrogen. The effect of water vapour accelerating the fatigue crack growth rate depended upon the frequency of load fluctuations and the water
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vapour pressure [12]. The kinetics of fatigue crack growth was controlled by the rate of transport of water vapour to the crack tip and, at higher pressures, by surface reactions of water vapour. A similar behaviour was found for corrosion fatigue of 7000-series aluminium alloys exposed to aggressive aqueous environments. Crack propagation rates for 7017T651 plate in seawater were dependent on loading frequency and stress intensity factor range DK [13]. Increasing these parameters resulted in changes of the fracture mode from intergranular to flat transgranular to striated ductile transgranular. The dependence of the fracture mode transitions upon the square root of the reciprocal of frequency indicated that hydrogen diffusion ahead of the crack tip was involved in crack growth. Enhanced fatigue crack growth rates were also found for 7150-T651 plate material immersed in an acidified aqueous salt solution, compared to those determined in dry laboratory air [14]. Corrosion fatigue crack growth rates increased with decreasing frequency. Intergranular and transgranular crack propagations predominated at low and high DK, respectively. An analysis of the frequency dependence of the intergranular–transgranular failure mode transition supported grain boundary diffusion of hydrogen being the controlling step in intergranular corrosion fatigue. Besides intergranular environmentally assisted cracking, transgranular SCC was also observed in aluminium alloys, particularly in Al–Zn–Mg–(Cu) alloys [2,4,15,16]. Transgranular SCC occurred when severe loading conditions were applied, being observed in static loading and slow strain rate tests. As supported by fractography, cracks initiated very often from critical defects produced by localized anodic dissolution, which promote localization of plasticity as well as a localized hydrogen discharge, entry and subsequent embrittlement [17]. This environment-induced cleavage-like fracture revealed crystallographic parallel facets separated by steps frequently exhibiting fan-like pattern [18,19]. Crack arrest markings on the fracture surfaces indicated discontinuous crack propagation. The occurrence of transgranular environment-induced cracking was very pronounced when cyclic loading was applied using displacement rates being in the range from 106 to 104 mm s1 [20]. These cyclic loading tests are in the interfacial region between stress corrosion cracking testing under static loading conditions and corrosion fatigue tests which are typically carried out at frequencies in the order of magnitude of 1–10 Hz, and, therefore, they could elucidate degradation mechanisms being operative under different loading conditions. The aim of the present work was to study the environment-induced cracking behaviour of alloy 7050 in different tempers under cyclic loading conditions at very low frequencies. Notched specimens were used being well established to study hydrogen-assisted fracture and considered to be valuable in identifying the micromechanisms involved [21].
2. Experimental The material used was a 73 mm thick plate of the aluminium alloy AA7050 received in the solution heat-treated and quenched condition W51. Cylindrical threaded bars machined from the plate were artificially aged to the peak-aged temper T651 (24 h at 120 °C + 12 h at 154 °C) and to the overaged tempers T7X51 (5 h at 120 °C + 12 h at 165 °C), T7651 (5 h at 120 °C + 15 h at 165 °C) and T7351 (5 h at 120 °C + 18 h at 165 °C). Heating-up rates of 30 °C/h and 15 °C/h were used in the first and second stages, respectively. After heat treatment, the bars were machined to notched, short transverse oriented round tension specimens with a gauge length of 23 mm. The 60° circumferential V-notch had a depth of 0.61 mm and a notch tip radius of 0.14 mm. The root diameter was 3.50 mm. Short transverse tensile properties of the alloy 7050 in the different tempers are given in Table 1. Notch strengths related to the specified sample geometry are also included, determined by tensile testing in an inert environment at a displacement rate of 2.5 106 mm s1. Fig. 1 shows transmission electron micrographs of alloy 7050 in the different overaged tempers. The matrix precipitates in the less overaged T7X heat treatment condition were finer than those observed in the T76 and T73 conditions. The microstructure of the former temper might approximately resemble that of the peak-aged condition which was found to be characterized predominantly by the presence of a fine dispersion of the g0 (MgZn2) transition phase being most probably coherent [22]. The matrix precipitates coarsened with increasing aging with a simultaneous loss of coherency [23]. Along the grain boundaries particles (equilibrium phase g) precipitated with concomitant precipitate-free zones (PFZs). The size of the grain boundary particles appeared to be smaller in the T7X temper than those in the more overaged T76 and T73 tempers, whereas the width of the PFZs seemed to be independent on aging, in agreement with other studies on grain boundary microstructure of Al–Zn–Mg–Cu alloys reporting a larger particles size in overaged material than that for the peak-aged condition and no significant difference in the PFZ width with aging treatment [24–26]. Static loading tests were carried out under constant load conditions using dead weight loaded tensile testing machines. Duplicate specimens were permanently immersed in aerated substitute ocean water without heavy Table 1 Short transverse tensile properties of alloy 7050 plate in different tempers Temper
0.2% Proof stress (MPa)
Ultimate tensile strength (MPa)
Fracture elongation (%)
Notch strength (MPa)
7050-T651 7050-T7X51 7050-T7651 7050-T7351
501 ± 4
571 ± 8
4.3 ± 0.8
454 ± 7 444 ± 9
522 ± 9 514 ± 13
4.8 ± 0.4 5.8 ± 0.9
735 ± 29 709 ± 18 695 ± 16 685 ± 6
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Fig. 1. TEM micrographs of alloy 7050 in the tempers (a) T7X51, (b) T7651 and (c) T7351, showing matrix precipitates, grain boundary particles and precipitate-free zones.
metals (ASTM D 1141) at free corrosion potential. The electrolyte was replaced every two weeks. Failure criterion was fracture. The maximum exposure length was 1000 h. Cyclic loading tests were performed using a slow strain rate tensile testing machine. The load applied exhibited a sawtooth shaped waveform, increasing from a lower to an upper limit at constant cross-head speed. The displacement rates were in the range from 2 106 to 2 104 mm s1. If the upper limit was attained, the stress was reduced very fast to the lower limit and was raised again. The maximum load corresponded to a stress of 600 MPa with regard to the original cross-sectional area at the notch, except with alloy 7050-T7651 which was mainly tested applying a maximum load of 500 MPa. The synthetic environment used was again aerated substitute ocean water under permanent immersion condition at free corrosion potential.
3. Results Time-to-failure data obtained from constant load tests for alloy 7050 in the different tempers are listed in Table 2. Alloy 7050-T651 was quite sensitive to stress corrosion cracking when notched tension specimens were stressed at 500 MPa and above in short transverse direction. Intergranular cracking was observed on the fracture surfaces (Fig. 2a). At an applied load of 400 MPa, one sample failed in the smooth part of the gauge length outside the notch, related to the low intergranular SCC resistance of 7050T651. Short transverse threshold stresses are typically below 100 MPa for peak-aged 7000-series aluminium alloys using smooth SCC specimens [27,28]. The other 7050-T651 sample passed the maximum exposure length of 1000 h at an applied stress of 400 MPa. The latter specimen was not deteriorated by immersion in the aggressive environ-
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Table 2 Time-to-failure data of alloy 7050 plate in different tempers. Notched tension specimens were permanently immersed in substitute ocean water under constant load conditions Alloy
600 MPa
500 MPa
400 MPa
7050-T651 7050-T7X51 7050-T7651 7050-T7351
9 h, 44 h 208 h, 241 h 234 h, 358 h 248 h, 302 h, 350 h
12 h, 244 h 2 >1000 h 2 >1000 h 2 >1000 h
435 h*, 1 >1000 h 2 >1000 h 2 >1000 h 2 >1000 h
* Specimen failed in the smooth part of the gauge length outside the notch.
and below, no failure occurred during the maximum exposure length of 1000 h. After exposure the specimens were tensile tested under inert environmental conditions at a displacement rate of 2.5 106 mm s1. Immersion in substitute ocean water did not deteriorate the notch strength of these specimens. Fractographic examinations revealed a dimple-like ductile overload fracture. No cleavage-like fracture was found. Results of cyclic loading tests for alloy 7050-T7651 are shown in Fig. 3. Specimens were loaded at different displacement rates, applying a stress range Dr = rmax– rmin = 250 MPa with rmax = 500 MPa, corresponding to an R-ratio of 0.5. For samples immersed in substitute ocean water, the number of cycles to failure decreased with decreasing frequency, e.g. decreasing displacement rate (Fig. 3a). A linear dependence was found between the log values (log Y = 0.507 log X + 4.152). Under inert environmental conditions (embedding the specimens in magnesium perchlorate hydrate), fatigue failure occurred after 1700– 2100 cycles. Due to long testing time periods, fatigue tests conducted in an inert environment were limited and
Fig. 2. Scanning electron fractographs of alloy 7050 in the tempers (a) T651 and (b) T7651, showing intergranular and transgranular stress corrosion cracking, respectively. Notched specimens were immersed in substitute ocean water under constant load conditions at applied stresses of (a) 500 and (b) 600 MPa and failed after 244 and 234 h of exposure, respectively.
ment as indicated by subsequent tensile testing revealing a notch strength of 716 MPa which was similar to that of the peak-aged material tensile tested under inert environmental conditions (Table 1). The fracture surface exhibited only ductile dimple-like overload fracture. Notched specimens of alloy 7050 in the overaged tempers failed after exposure time periods between 200 and 360 h when loaded at 600 MPa. Transgranular SCC was found on the fracture surfaces of failed specimens (Fig. 2b). Near the transition to overload fracture, striations consisting of deep slots were often observed on the flat transgranular facets. At applied stresses of 500 MPa
Fig. 3. Curves of (a) cycles to failure and (b) time to failure vs. frequency for alloy 7050-T7651. Notched specimens were cyclically loaded in substitute ocean water and in an inert environment applying a stress range Dr = 250 MPa at a maximum stress rmax = 500 MPa.
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performed at higher frequencies. Time to failure data as function of the frequency are plotted in Fig. 3b. The logarithm of time to failure decreased linearly with increasing logarithm of frequency. The slopes of the straight lines were 0.50 and 1.06 for cyclic loading in corrosive and inert environments, respectively. According to the intersection of the lines, the specimens would fail by (pure) mechanical fatigue at frequencies of 1.5 102 cycles/s and above without crack growth enhancement by the corrosive environment. However, it should be mentioned that the slope of the line representing time to failure data obtained under inert environmental condition was rather uncertain due to the small frequency range used. Fractographic examinations of specimens failed by pure mechanical fatigue revealed a mainly ductile dimple fracture with small areas of cleavage-like fracture at the rim of the fracture surfaces (Fig. 4). No striations were found. The fracture surfaces of specimens cyclically loaded in substitute ocean water exhibited a brittle transgranular fracture, being more flat than those of samples tested in an inert environment (Fig. 5). On the areas of flat transgranular crack growth, striations were observed traversing several facets. Striation spacing increased as the crack propagated in the interior of the material due to the higher stress resulting from reduced cross-sectional area. With increasing loading frequency regularly spaced striations were difficult to find. Irregularly shaped rough markings were observed instead, resembling secondary cracks (Fig. 5d). Results of the cyclic loading tests for alloy 7050-T7X51 are plotted in Fig. 6. A stress range Dr = 200 MPa with rmax = 600 MPa was applied corresponding to an R-ratio of 0.67. Again, for specimens immersed in substitute ocean water, the number of cycles to failure decreased with decreasing displacement rate, depending linearly upon the frequency on a log–log basis (log Y = 0.592 log X + 4.357). Under inert environmental conditions, failure occurred after number of cycles in the range from 2700 to 3500. The number of cycles to failure caused by pure mechanical fatigue were higher for 7050-T7X51 specimens compared to those obtained for the alloy in the T7651 temper. This might be associated with the higher R-ratio and the higher notch strength of alloy 7050-T7X51. Plotting time to failure data as function of the frequency slopes of 0.42 and 0.90 were obtained for cyclic loading tests in corrosive and inert environments, respectively (Fig. 6b). Both lines intersect at a frequency of 8.4 102 cycles/s. Thus, the corrosive environment had no effect on crack growth at frequencies above 8 102 cycles/s. Fractographic examinations of the 7050-T7X51 specimens cyclically loaded under inert environmental conditions revealed again cleavage facets at the rim of the fracture surface (Fig. 7). No striations were observed. Fig. 8 shows the fracture surface of a 7050-T7X51 sample cyclically tested in substitute ocean water. Pronounced brittle transgranular crack propagation was observed. On the flat transgranular facets, coarse striations consisting of
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Fig. 4. Scanning electron fractographs of alloy 7050-T7651, showing cleavage-like fracture. The notched specimens were cyclically loaded under inert environmental conditions at frequencies of (a) 5.1 104 and (b) 1.3 103 cycles/s applying a stress range Dr = 250 MPa at a maximum stress rmax = 500 MPa.
deep slots were found. The distance between the crack arrest markings increased with decreasing frequency. Fig. 9 displays results of cyclic loading tests for alloy 7050-T7351 at R-ratios of 0.83 and 0.67 with rmax = 600 MPa. In the frequency range used, number of cycles to failure were higher for the higher R-ratio, but the slope of a linear plot on a log–log basis was slightly steeper compared to that for R = 0.67. However, cyclic loading tests carried out at different R-ratios did not indicate a marked dependence of the slope upon the R-ratio [20]. Fractographic examinations revealed cleavage facets with coarse striations (Fig. 10). The striation spacing increased with increasing stress range; it decreased with higher frequency. At frequencies above 103 s1, striations were difficult to find. Results of cyclic loading tests carried out with notched specimens of alloy 7050-T651 are presented in Fig. 11, applying stress ranges Dr of 100 and 200 MPa with rmax = 600 MPa. At higher frequencies (corresponding to a displacement rate of 3.2 105 mm s1), the number of cycles to failure were similar to those observed for alloy 7050T7351. Applying a displacement rate of 2.5 106 mm s1 (resulting in frequencies below 5 105 s1), the 7050-
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Fig. 5. Scanning electron fractographs of alloy 7050-T7651, showing brittle transgranular fracture. Notched specimens were cyclically loaded in substitute ocean water at frequencies of (a) 1.9 105, (b) 4.6 105, (c) 2.4 104, and (d) 8.0 104 cycles/s applying a stress range Dr = 250 MPa at a maximum stress rmax = 500 MPa.
T651 specimens failed after lower number of cycles. This was associated with the occurrence of intergranular stress corrosion cracking, as corroborated by fractography. On the fracture surfaces of samples cyclically tested at a displacement rate 2.5 106 mm s1, both intergranular and transgranular environment-induced cracking were observed, whereas brittle transgranular fracture was more pronounced when 7050-T651 specimens were loaded at higher frequencies (Fig. 12). Both coarse slot-like and fine striations were observed (Fig. 13). The number of cycles to failure as function of the frequency is plotted in Fig. 14 for all cyclic loading tests conducted with alloy 7050 applying a stress range Dr = 200 MPa with rmax = 600 MPa. At an applied R-ratio of 0.67, the frequency dependence of failure caused by transgranular environment-induced cracking was found to be similar for alloy 7050 in all tempers studied. On a log–log basis, the dependence was linear described by log Y (number of cycles to failure) = 0.596 log X (frequency) + 4.363. Therefore, a unique mechanism causing brittle transgranular environment-induced cracking seems to operate in alloy 7050, being not influenced by the heat treatment conditions. 4. Discussion Stress corrosion cracking of high strength aluminium alloys is a well known degradation process in loaded struc-
tural components described in several review papers [1–4]. The crack propagates predominantly intergranular, and the SCC sensibility can be reduced by artificial aging. Two main mechanisms have been proposed for stress corrosion cracking in aluminium alloys: highly localized anodic dissolution of the grain boundaries and embrittlement due to hydrogen ingress into the alloy. When severe loading conditions are applied, such as dynamic straining to fracture (slow strain rate testing) or stress concentration effects associated with pitting in statically loaded smooth tensile specimens, transgranular stress corrosion cracking can also occur in aluminium alloys, revealing cleavage-like fracture and crack arrest markings [2,4,15–19,29]. The mechanisms proposed for the latter type of environmentinduced cracking include a combined action of anodic dissolution and extremely localized ductile fracture, an adsorption-induced localized-slip process, or the corrosion enhanced plasticity model [19,30,31]. Fatigue crack growth in aluminium alloys is generally transgranular and, under inert environmental conditions, related to alternating shear at the crack tip and controlled by the crack tip opening displacement [32]. In alloy 7150T651, the fatigue crack growth rate in air was found to be independent of frequency in the range between 0.1 and 20 Hz [14]. Ductile transgranular crack propagation was predominantly observed on the fracture surfaces. When fatigue testing was performed in an aggressive aqueous
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Fig. 6. Curves of (a) cycles to failure and (b) time to failure vs. frequency for alloy 7050-T7X51. Notched specimens were cyclically loaded in substitute ocean water and in an inert environment applying a stress range Dr = 200 MPa at a maximum stress rmax = 600 MPa.
environment, the crack growth rates increased by up to an order of magnitude compared to those measured in laboratory air. Fractography revealed both brittle intergranular and trangranular fracture modes, being dominant at low and high DK, respectively. The extent of intergranular fracture mode and the enhancement of fatigue crack growth rates depended upon cyclic frequency, supporting a mechanism of hydrogen grain boundary diffusion. The crystallographic and flat transgranular crack growth being operative at high DK exhibited little ductility and was characterized by the presence of brittle striations. Transgranular corrosion fatigue was proposed to be dependent on strain-controlled hydrogen diffusion ahead of the crack [14]. Similar results were found by Holroyd and Hardie conducting fatigue crack growth tests of alloy 7017-T651 in seawater [13]. At all frequencies tested, the crack growth rates significantly increased in seawater, compared to dry air, and the enhancement was more marked at the lower frequencies. Crack growth rates in seawater depended upon both loading frequency and DK. An increase of these
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Fig. 7. Scanning electron fractographs of alloy 7050-T7X51, showing cleavage-like fracture. Notched specimens were cyclically loaded under inert environmental conditions at frequencies of (a) 1.7 103 and (b) 1.0 103 cycles/s applying a stress range Dr = 200 MPa at a maximum stress rmax = 600 MPa.
parameter resulted in a change of fracture mode from intergranular to flat transgranular to striated ductile transgranular. A linear fit was obtained plotting logarithmically the crack velocities associated with the fracture mode transitions vs. frequency of loading. The dependence of the velocity involved in each transition on the square root of the cycle period confirmed again that the enhanced crack growth in seawater was controlled by diffusion of hydrogen ahead of the crack tip. For alloy 7050-T7651 cyclically loaded under inert environmental conditions, the time to failure tF decreased linearly with increasing frequency m on a log–log basis with a slope of 1 (Fig. 3b). Therefore, the number of cycles to failure nF, given by nF = tF m, was constant being independent upon frequency. When the cyclic loading tests were carried out in substitute ocean water, a slope of 0.5 was obtained fitting the logarithmic plot of time to failure vs. frequency. The number of cycles to failure was dependent on p the square root of frequency, i.e. proportional to 1/ s with s being the cycle period. This square root dependence indicated an embrittlement mechanism controlled by diffusion. Thus, the enhanced crack growth in substitute ocean water was related to hydrogen diffusion ahead of the crack tip, in accordance with the suggestion made by Holroyd and Hardie [13].
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Fig. 8. Scanning electron fractographs of alloy 7050-T7X51, showing (a) brittle transgranular fracture and (b) crack arrest markings. The notched specimen was cyclically loaded in substitute ocean water at a frequency of 2.3 105 cycles/s, applying a stress range Dr = 200 MPa at a maximum stress rmax = 600 MPa.
Fig. 9. Number of cycles to failure vs. frequency for alloy 7050-T7351. Notched specimens were cyclically loaded in substitute ocean water applying stress ranges Dr of 100 and 200 MPa with rmax = 600 MPa.
When the maximum stress was increased to 600 MPa, the slope of the time-to-failure vs. frequency curve plotted logarithmically was 0.90 for notched specimens cyclically loaded in an inert environment (Fig. 6b). The maximum load exceeded significantly the yield strength of the alloy 7050-T7X51, resulting in a slight dependence of the num-
Fig. 10. Scanning electron fractographs of alloy 7050-T7351, showing (a) brittle transgranular fracture and (b) crack arrest markings. The notched sample was cyclically loaded in substitute ocean water at a frequency of 3.1 104 cycles/s applying a stress range Dr = 200 MPa with rmax = 600 MPa.
Fig. 11. Number of cycles to failure vs. frequency for alloy 7050-T651. Notched specimens were cyclically loaded in substitute ocean water applying stress ranges Dr of 100 and 200 MPa with rmax = 600 MPa.
ber of cycles to failure upon the frequency. Besides hydrogen embrittlement this mechanical effect also contributed to the failure of samples immersed in an aggressive environment, leading to a slope of 0.42. Fatigue crack growth in alloy 7075-T651 was found to be intermittent, not proceeding cycle-by-cycle both in moist
R. Braun / International Journal of Fatigue 30 (2008) 1827–1837
Fig. 12. Scanning electron fractographs of alloy 7050-T651, showing (a) intergranular and transgranular and (b) transgranular environmentinduced cracking. The notched samples were cyclically loaded in substitute ocean water at frequencies of (a) 4.8 105 and (b) 6.4 104 cycles/s applying a stress range Dr of 100 MPa with rmax = 600 MPa.
air and vacuum [33]. Before the crack propagates, it blunts and accumulates local strain damage near the tip during a number of cycles which decreases with higher cyclic stress intensity. The plasticity required for an increment of crack extension was reduced in moist air compared to that in vacuum. Whereas fracture surfaces resulting from purely plastic processes in inert environments were non-crystallographic and relatively featureless, crystallographic fatigue crack growth on {1 0 0} and {1 1 0} planes was observed in moist environments [11,12,14]. In the present study, cleavage facets were found at the rim of the fracture surfaces of notched specimens cyclically loaded under inert environmental conditions (Figs. 4 and 7). This might be caused by a residual humidity of the laboratory air not entirely removed by embedding the specimens in magnesium perchlorate hydrate. However, slow strain rate tests carried out to evaluate the SCC behaviour of aluminium alloys indicated that this wrapping procedure can be used for reference tests in an inert environment [28,34]. Therefore, the crystallographic facets were rather attributed to slip processes at the crack tip activated particularly by the high maximum load. When specimens were immersed in substitute ocean water hydrogen ahead of the crack tip promoted slip, thus enhancing crack advance. The failure of the specimens cyclically loaded under inert environmen-
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Fig. 13. Scanning electron fractographs of alloy 7050-T651, showing (a) coarse and (b) fine striations. The notched sample was cyclically loaded in substitute ocean water at a frequency of 6.4 104 cycles/s applying a stress range Dr of 100 MPa with rmax = 600 MPa.
Fig. 14. Number of cycles to failure vs. frequency for alloy 7050 in different tempers. Notched specimens were cyclically loaded in substitute ocean water applying a stress range Dr = 200 MPa at a maximum stress rmax = 600 MPa. The log values were linearly fitted.
tal conditions might be also associated with accumulation of plastic strain caused by cyclic creep/ratcheting when stress levels close to or exceeding the ultimate tensile strength were applied [35,36]. However, as found with SAE 1045 steel, fractographic examinations of samples
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failed due to cyclic creep/ratcheting revealed ductile dimple fracture resulting from microvoid coalescence [37]. Thus, the cleavage fracture observed at the rim of alloy 7050 specimens was related to fatigue cracking. Cleavage facets were not observed with notched specimens monotonically tensile tested in an inert environment. Striations are typically found on corrosion fatigue fracture surfaces [11,12,14,33]. Several types of striations were observed varying from a sawtooth shaped profile (type A) to another one consisting of deep slots traversing a flat surface (type B) [11]. It was concluded that the striation profile depended upon the orientation of the cleavage plane and the slip systems to the tensile and shear stresses. In the present study, the flat transgranular facets of the samples cyclically loaded in substitute ocean water were covered with striations. These crack arrest markings indicated an intermittent crack growth in the corrosive environment. The striations were predominantly regularly arranged and consisted of slots (Figs. 8, 10 and 13), being similar to the type B and brittle striations observed by Nix and Flower and Gao et al., respectively [11,12]. This type of striation profile occurred when the cleavage plane was oriented orthogonal to the tensile stress axis [11]. Occasionally, fine regularly spaced striations were also found related to crack advance during a cycle (Figs. 5a and 13b), resembling the ductile striations observed by Gao et al. on fracture surfaces of 7075-T651 specimens tested in water vapour [12]. The latter type of striations was probably created by dissolution of the exposed bare fracture surface at crack arrest during the period of reduced load. With increasing frequency, the regularly arranged striations disappeared, and irregularly shaped markings were observed near the transition region to overload fracture (Figs. 5d and 12b). These markings might be secondary cracks related to the high final stress before failure. Because these cracks did not occur on the cleavage fracture of specimens cyclically loaded in an inert environment, they were also associated with the corrosive environment, as confirmed by their presence on the fracture surfaces of specimens failed under static loading conditions (Fig. 2b). Whereas the intergranular SCC susceptibility of 7000series aluminium alloys depends upon heat treatment [1], sensitivity to transgranular environment-induced cracking was found to be independent upon temper condition, at least when a maximum stress of 600 MPa was applied which was close to the notch strength of the specimens used. This was confirmed by the occurrence of transgranular cleavage-like fracture in alloy 7050 in all tempers as well as the single straight line fitting the time to failure data of all notched specimens cyclically loaded at a stress range Dr = 200 MPa with rmax = 600 MPa (Fig. 14). Therefore, the diffusion process of hydrogen causing embrittlement was not significantly influenced by changes of the microstructure due to artificial aging. Because fractographic features of transgranular environmentally assisted cracking were similar when specimens were stressed under static, dynamic and cyclic loading conditions [20,29], the same
transgranular embrittlement mechanism might be operative in both stress corrosion cracking and corrosion fatigue at low frequencies. 5. Conclusions 1. Alloy 7050 was found to be sensitive to transgranular environment-induced cracking when cyclically loaded in substitute ocean water at frequencies between 2 105 and 4 103 cycles/s. 2. The number of cycles to failure decreased with decreasing frequency. On a log–log basis, a linear dependence was determined with slopes of 0.5 and 0.6 applying maximum stresses of 500 and 600 MPa, respectively. 3. Striations on the transgranular facets of the fracture surfaces indicated discontinuous crack propagation. Crack arrest markings were more pronounced when cyclic loading was carried out at lower frequencies. 4. No influence of the heat treatment on transgranular environment-induced cracking of alloy 7050 was observed. 5. Environment-induced failure under cyclic loading conditions at low frequencies was probably associated with hydrogen embrittlement.
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R. Braun / International Journal of Fatigue 30 (2008) 1827–1837 [10] Ricker RE, Duquette DJ. The role of hydrogen in corrosion fatigue of high purity Al–Zn–Mg exposed to water vapor. Metall Trans A 1988;19A(7):1775–83. [11] Nix KJ, Flower HM. The micromechanisms of fatigue crack growth in a commercial Al–Zn–Mg–Cu alloy. Acta Metall 1982;30(8): 1549–59. [12] Gao M, Pao PS, Wei RP. Chemical and metallurgical aspects of environmentally assisted fatigue crack growth in 7075-T651 aluminum alloy. Metall Trans A 1988;19A(7):1739–50. [13] Holroyd NJH, Hardie D. Factors controlling crack velocity in 7000 series aluminium alloys during fatigue in an aggressive environment. Corros Sci 1983;23(6):527–46. [14] Gingell ADB, King JE. The effect of frequency and microstructure on corrosion fatigue crack propagation in high strength aluminium alloys. Acta Mater 1997;45(9):3855–70. [15] Magnin T. Recent advances in the environment sensitive fracture mechanisms of aluminium alloys. Mater Sci Forum 1996;217– 222:83–94. [16] Braun R. Environmentally assisted cracking of aluminium alloys in chloride solution. In: Sato T, Kumai S, Kobayashi T, Murakami Y, editors. Aluminum alloys – their physical and mechanical properties (ICAA-6). Tokyo: The Japan Institute of Light Metals; 1998. p. 153–64. [17] Najjar D, Magnin T, Warner TJ. Influence of critical surface defects and localized competition between anodic dissolution and hydrogen effects during stress corrosion cracking of a 7050 aluminium alloy. Mater Sci Eng 1997;A238(2):293–302. [18] Pugh EN. Progress toward understanding the stress corrosion problem. Corrosion 1985;41(9):517–26. [19] Kaufman MJ, Fink JL. Evidence for localized ductile fracture in the ‘‘brittle” transgranular stress corrosion cracking of ductile F.C.C. alloys. Acta Metall 1988;36(8):2213–28. [20] Braun R. Transgranular environmentally induced cracking in the Al– Zn–Mg–Cu alloy 7050. In: Moody NR, Thompson AW, Ricker RE, Was GW, Jones RH, editors. Hydrogen effects on material behavior and corrosion deformation interactions. Warrendale: The Minerals, Metals & Materials Society; 2003. p. 883–92. [21] Thompson AW. Hydrogen-assisted fracture at notches. Mater Sci Technol 1985;1(9):711–8. [22] Park JK, Ardell AJ. Microstructure of the commercial 7075 Al alloy in the T651 and T7 tempers. Metall Trans A 1983;14A(10):1957–65. [23] Dorward RC. Precipitate coarsening during overaging of Al–Zn–Mg– Cu alloy. Mater Sci Technol 1999;15(10):1133–8.
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[24] Hepples W, Jarrett MR, Crompton JS, Holroyd NJH. The influence of microstructure on the stress corrosion cracking and exfoliation of commercial Al–Zn–Mg–Cu alloys. In: Gangloff RP, Ives MB, editors. Environment-induced cracking of metals. Houston: National Association of Corrosion Engineers; 1989. p. 383–7. [25] Park JK, Ardell AJ. Effect of retrogression and reaging treatments on the microstructure of Al-7075-T651. Metall Trans A 1984;15A(8): 1531–43. [26] Rao BVN. Influence of aging treatment on the grain boundary precipitation and chemistry in 7075 Al alloy. Metall Trans A 1981;12A(7):1356–9. [27] Hollingsworth EH, Hunsicker HY. Corrosion of aluminum and aluminum alloys. Metals handbook. Corrosion, 9th ed., vol. 13. Metals Park, Ohio: ASM International; 1987. p. 583–609. [28] Braun R. Slow strain rate testing of aluminum alloy 7050 in different tempers using various synthetic environments. Corrosion 1997;53(6): 467–74. [29] Braun R. Environmentally assisted cracking of aluminium alloys. Mat-wiss u Werkstofftech 2007;38(9):674–89. [30] Lynch SP. Environmentally assisted cracking: overview of evidence for an adsorption-induced localised-slip process. Acta Metall 1988;36(10):2639–61. [31] Magnin T, Chambreuil A, Chateau JP. Stress corrosion cracking mechanisms in ductile FCC materials. Int J Fract 1996;79(2):147–63. [32] Speidel MO. Stress corrosion and corrosion fatigue crack growth in aluminum alloys. In: Arup H, Parkins RN, editors. Stress corrosion research. The Netherlands: Sijthoff & Noordhoff, Alphen aan den Rijn; 1979. p. 117–76. [33] Lankford J, Davidson DL. Fatigue crack micromechanisms in ingot and powder metallurgy 7XXX aluminum alloys in air and vacuum. Acta Metall 1983;31(8):1273–84. [34] Braun R, Buhl H. Stress corrosion testing. In: Buhl H, editor. Advanced aerospace materials. Berlin: Springer Verlag; 1992. p. 296–307. [35] Satyadevi A, Sivakumar SM, Bhattacharya SS. A new failure criterion for materials exhibiting ratcheting during very low cycle fatigue. Mater Sci Eng A 2007;452–453:380–5. [36] Pals TG, Stephens RI. The influence of high R ratio on mild and sharp notched and unnotched fatigue behaviour of 1045 steel with three different heat treatments. Int J Fatigue 2004;26(6):651–61. [37] Vantiger TR, Stephens RI, Karadag M. The influence of high R ratio on notched fatigue behaviour of 1045 steel with three different heat treatments. Int J Fatigue 2002;24(12):1275–84.
International Journalof Fatigue
International Journal of Fatigue 30 (2008) 1827–1837
www.elsevier.com/locate/ijfatigue
Transgranular environment-induced cracking of 7050 aluminium alloy under cyclic loading conditions at low frequencies Reinhold Braun * DLR – German Aerospace Center, Institute of Materials Research, D-51170 Ko¨ln, Germany Received 25 October 2007; accepted 11 February 2008 Available online 19 February 2008
Abstract The environment sensitive cracking behaviour of alloy 7050 plate in different tempers was studied under cyclic loading conditions using short transverse oriented notched tension specimens immersed in substitute ocean water. The cyclic loading tests were carried out at displacement rates in the range from 2 106 to 2 104 mm s1 applying a sawtooth wave form. The maximum load corresponded to an initial stress of 600 MPa, except for alloy 7050-T7651 being loaded up to 500 MPa. Number of cycles to failure decreased with decreasing displacement rate. A linear relationship was found between log values of the number of cycles to failure and the loading frequency. Fractographic examinations of failed specimens revealed brittle transgranular fracture with crack arrest markings on flat facets. The sensitivity to transgranular environment-induced cracking proposed to be associated with hydrogen embrittlement did not depend on heat treatment conditions. Ó 2008 Elsevier Ltd. All rights reserved. Keywords: AA7050; Cyclic loading; Transgranular environment-induced cracking; Striations
1. Introduction Age-hardenable 7000-series aluminium alloys are sensitive to environmentally assisted cracking including stress corrosion cracking (SCC), hydrogen embrittlement and corrosion fatigue [1–8]. The SCC susceptibility of high strength aluminium alloys decreases from underaged to overaged microstructures [1–3]. Whereas Al–Zn–Mg–Cu alloys are quite sensitive to stress corrosion cracking in the peak-aged T6 temper, two-stage overaging improves the SCC resistance, resulting in virtual immunity for the T73 temper [4]. Stress corrosion cracking causing failure in commercial aluminium alloys is generally intergranular [1,2,4]. A similar behaviour was observed for hydrogenenvironment assisted cracking [9]. Using double-cantilever beam fracture mechanics specimens exposed to 90% relative humidity air, the stage II (plateau) crack growth rate *
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of alloy 7050 was found to decrease from the underaged to overaged conditions. In all tempers, the fracture path was intergranular, predominantly along high-angle grain boundaries. Water vapour was also found to promote corrosion fatigue. The fatigue resistance of a high purity Al–Zn– Mg alloy was significantly reduced when exposed to a humid nitrogen environment [10]. The degradation of fatigue resistance in water vapour was concluded to be associated with hydrogen embrittlement. Transgranular or mixed intergranular and transgranular cracking was usually observed with corrosion fatigue in aluminium alloys [7,10]. Whereas alloy 7010-T76 exhibited entirely ductile fatigue fracture in dry environments, failure in moist air occurred by cleavage with well defined surface striations [11]. Again, crack growth by cleavage was suggested to be the result of hydrogen embrittlement involving dislocation transport of hydrogen. The effect of water vapour accelerating the fatigue crack growth rate depended upon the frequency of load fluctuations and the water
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vapour pressure [12]. The kinetics of fatigue crack growth was controlled by the rate of transport of water vapour to the crack tip and, at higher pressures, by surface reactions of water vapour. A similar behaviour was found for corrosion fatigue of 7000-series aluminium alloys exposed to aggressive aqueous environments. Crack propagation rates for 7017T651 plate in seawater were dependent on loading frequency and stress intensity factor range DK [13]. Increasing these parameters resulted in changes of the fracture mode from intergranular to flat transgranular to striated ductile transgranular. The dependence of the fracture mode transitions upon the square root of the reciprocal of frequency indicated that hydrogen diffusion ahead of the crack tip was involved in crack growth. Enhanced fatigue crack growth rates were also found for 7150-T651 plate material immersed in an acidified aqueous salt solution, compared to those determined in dry laboratory air [14]. Corrosion fatigue crack growth rates increased with decreasing frequency. Intergranular and transgranular crack propagations predominated at low and high DK, respectively. An analysis of the frequency dependence of the intergranular–transgranular failure mode transition supported grain boundary diffusion of hydrogen being the controlling step in intergranular corrosion fatigue. Besides intergranular environmentally assisted cracking, transgranular SCC was also observed in aluminium alloys, particularly in Al–Zn–Mg–(Cu) alloys [2,4,15,16]. Transgranular SCC occurred when severe loading conditions were applied, being observed in static loading and slow strain rate tests. As supported by fractography, cracks initiated very often from critical defects produced by localized anodic dissolution, which promote localization of plasticity as well as a localized hydrogen discharge, entry and subsequent embrittlement [17]. This environment-induced cleavage-like fracture revealed crystallographic parallel facets separated by steps frequently exhibiting fan-like pattern [18,19]. Crack arrest markings on the fracture surfaces indicated discontinuous crack propagation. The occurrence of transgranular environment-induced cracking was very pronounced when cyclic loading was applied using displacement rates being in the range from 106 to 104 mm s1 [20]. These cyclic loading tests are in the interfacial region between stress corrosion cracking testing under static loading conditions and corrosion fatigue tests which are typically carried out at frequencies in the order of magnitude of 1–10 Hz, and, therefore, they could elucidate degradation mechanisms being operative under different loading conditions. The aim of the present work was to study the environment-induced cracking behaviour of alloy 7050 in different tempers under cyclic loading conditions at very low frequencies. Notched specimens were used being well established to study hydrogen-assisted fracture and considered to be valuable in identifying the micromechanisms involved [21].
2. Experimental The material used was a 73 mm thick plate of the aluminium alloy AA7050 received in the solution heat-treated and quenched condition W51. Cylindrical threaded bars machined from the plate were artificially aged to the peak-aged temper T651 (24 h at 120 °C + 12 h at 154 °C) and to the overaged tempers T7X51 (5 h at 120 °C + 12 h at 165 °C), T7651 (5 h at 120 °C + 15 h at 165 °C) and T7351 (5 h at 120 °C + 18 h at 165 °C). Heating-up rates of 30 °C/h and 15 °C/h were used in the first and second stages, respectively. After heat treatment, the bars were machined to notched, short transverse oriented round tension specimens with a gauge length of 23 mm. The 60° circumferential V-notch had a depth of 0.61 mm and a notch tip radius of 0.14 mm. The root diameter was 3.50 mm. Short transverse tensile properties of the alloy 7050 in the different tempers are given in Table 1. Notch strengths related to the specified sample geometry are also included, determined by tensile testing in an inert environment at a displacement rate of 2.5 106 mm s1. Fig. 1 shows transmission electron micrographs of alloy 7050 in the different overaged tempers. The matrix precipitates in the less overaged T7X heat treatment condition were finer than those observed in the T76 and T73 conditions. The microstructure of the former temper might approximately resemble that of the peak-aged condition which was found to be characterized predominantly by the presence of a fine dispersion of the g0 (MgZn2) transition phase being most probably coherent [22]. The matrix precipitates coarsened with increasing aging with a simultaneous loss of coherency [23]. Along the grain boundaries particles (equilibrium phase g) precipitated with concomitant precipitate-free zones (PFZs). The size of the grain boundary particles appeared to be smaller in the T7X temper than those in the more overaged T76 and T73 tempers, whereas the width of the PFZs seemed to be independent on aging, in agreement with other studies on grain boundary microstructure of Al–Zn–Mg–Cu alloys reporting a larger particles size in overaged material than that for the peak-aged condition and no significant difference in the PFZ width with aging treatment [24–26]. Static loading tests were carried out under constant load conditions using dead weight loaded tensile testing machines. Duplicate specimens were permanently immersed in aerated substitute ocean water without heavy Table 1 Short transverse tensile properties of alloy 7050 plate in different tempers Temper
0.2% Proof stress (MPa)
Ultimate tensile strength (MPa)
Fracture elongation (%)
Notch strength (MPa)
7050-T651 7050-T7X51 7050-T7651 7050-T7351
501 ± 4
571 ± 8
4.3 ± 0.8
454 ± 7 444 ± 9
522 ± 9 514 ± 13
4.8 ± 0.4 5.8 ± 0.9
735 ± 29 709 ± 18 695 ± 16 685 ± 6
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Fig. 1. TEM micrographs of alloy 7050 in the tempers (a) T7X51, (b) T7651 and (c) T7351, showing matrix precipitates, grain boundary particles and precipitate-free zones.
metals (ASTM D 1141) at free corrosion potential. The electrolyte was replaced every two weeks. Failure criterion was fracture. The maximum exposure length was 1000 h. Cyclic loading tests were performed using a slow strain rate tensile testing machine. The load applied exhibited a sawtooth shaped waveform, increasing from a lower to an upper limit at constant cross-head speed. The displacement rates were in the range from 2 106 to 2 104 mm s1. If the upper limit was attained, the stress was reduced very fast to the lower limit and was raised again. The maximum load corresponded to a stress of 600 MPa with regard to the original cross-sectional area at the notch, except with alloy 7050-T7651 which was mainly tested applying a maximum load of 500 MPa. The synthetic environment used was again aerated substitute ocean water under permanent immersion condition at free corrosion potential.
3. Results Time-to-failure data obtained from constant load tests for alloy 7050 in the different tempers are listed in Table 2. Alloy 7050-T651 was quite sensitive to stress corrosion cracking when notched tension specimens were stressed at 500 MPa and above in short transverse direction. Intergranular cracking was observed on the fracture surfaces (Fig. 2a). At an applied load of 400 MPa, one sample failed in the smooth part of the gauge length outside the notch, related to the low intergranular SCC resistance of 7050T651. Short transverse threshold stresses are typically below 100 MPa for peak-aged 7000-series aluminium alloys using smooth SCC specimens [27,28]. The other 7050-T651 sample passed the maximum exposure length of 1000 h at an applied stress of 400 MPa. The latter specimen was not deteriorated by immersion in the aggressive environ-
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Table 2 Time-to-failure data of alloy 7050 plate in different tempers. Notched tension specimens were permanently immersed in substitute ocean water under constant load conditions Alloy
600 MPa
500 MPa
400 MPa
7050-T651 7050-T7X51 7050-T7651 7050-T7351
9 h, 44 h 208 h, 241 h 234 h, 358 h 248 h, 302 h, 350 h
12 h, 244 h 2 >1000 h 2 >1000 h 2 >1000 h
435 h*, 1 >1000 h 2 >1000 h 2 >1000 h 2 >1000 h
* Specimen failed in the smooth part of the gauge length outside the notch.
and below, no failure occurred during the maximum exposure length of 1000 h. After exposure the specimens were tensile tested under inert environmental conditions at a displacement rate of 2.5 106 mm s1. Immersion in substitute ocean water did not deteriorate the notch strength of these specimens. Fractographic examinations revealed a dimple-like ductile overload fracture. No cleavage-like fracture was found. Results of cyclic loading tests for alloy 7050-T7651 are shown in Fig. 3. Specimens were loaded at different displacement rates, applying a stress range Dr = rmax– rmin = 250 MPa with rmax = 500 MPa, corresponding to an R-ratio of 0.5. For samples immersed in substitute ocean water, the number of cycles to failure decreased with decreasing frequency, e.g. decreasing displacement rate (Fig. 3a). A linear dependence was found between the log values (log Y = 0.507 log X + 4.152). Under inert environmental conditions (embedding the specimens in magnesium perchlorate hydrate), fatigue failure occurred after 1700– 2100 cycles. Due to long testing time periods, fatigue tests conducted in an inert environment were limited and
Fig. 2. Scanning electron fractographs of alloy 7050 in the tempers (a) T651 and (b) T7651, showing intergranular and transgranular stress corrosion cracking, respectively. Notched specimens were immersed in substitute ocean water under constant load conditions at applied stresses of (a) 500 and (b) 600 MPa and failed after 244 and 234 h of exposure, respectively.
ment as indicated by subsequent tensile testing revealing a notch strength of 716 MPa which was similar to that of the peak-aged material tensile tested under inert environmental conditions (Table 1). The fracture surface exhibited only ductile dimple-like overload fracture. Notched specimens of alloy 7050 in the overaged tempers failed after exposure time periods between 200 and 360 h when loaded at 600 MPa. Transgranular SCC was found on the fracture surfaces of failed specimens (Fig. 2b). Near the transition to overload fracture, striations consisting of deep slots were often observed on the flat transgranular facets. At applied stresses of 500 MPa
Fig. 3. Curves of (a) cycles to failure and (b) time to failure vs. frequency for alloy 7050-T7651. Notched specimens were cyclically loaded in substitute ocean water and in an inert environment applying a stress range Dr = 250 MPa at a maximum stress rmax = 500 MPa.
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performed at higher frequencies. Time to failure data as function of the frequency are plotted in Fig. 3b. The logarithm of time to failure decreased linearly with increasing logarithm of frequency. The slopes of the straight lines were 0.50 and 1.06 for cyclic loading in corrosive and inert environments, respectively. According to the intersection of the lines, the specimens would fail by (pure) mechanical fatigue at frequencies of 1.5 102 cycles/s and above without crack growth enhancement by the corrosive environment. However, it should be mentioned that the slope of the line representing time to failure data obtained under inert environmental condition was rather uncertain due to the small frequency range used. Fractographic examinations of specimens failed by pure mechanical fatigue revealed a mainly ductile dimple fracture with small areas of cleavage-like fracture at the rim of the fracture surfaces (Fig. 4). No striations were found. The fracture surfaces of specimens cyclically loaded in substitute ocean water exhibited a brittle transgranular fracture, being more flat than those of samples tested in an inert environment (Fig. 5). On the areas of flat transgranular crack growth, striations were observed traversing several facets. Striation spacing increased as the crack propagated in the interior of the material due to the higher stress resulting from reduced cross-sectional area. With increasing loading frequency regularly spaced striations were difficult to find. Irregularly shaped rough markings were observed instead, resembling secondary cracks (Fig. 5d). Results of the cyclic loading tests for alloy 7050-T7X51 are plotted in Fig. 6. A stress range Dr = 200 MPa with rmax = 600 MPa was applied corresponding to an R-ratio of 0.67. Again, for specimens immersed in substitute ocean water, the number of cycles to failure decreased with decreasing displacement rate, depending linearly upon the frequency on a log–log basis (log Y = 0.592 log X + 4.357). Under inert environmental conditions, failure occurred after number of cycles in the range from 2700 to 3500. The number of cycles to failure caused by pure mechanical fatigue were higher for 7050-T7X51 specimens compared to those obtained for the alloy in the T7651 temper. This might be associated with the higher R-ratio and the higher notch strength of alloy 7050-T7X51. Plotting time to failure data as function of the frequency slopes of 0.42 and 0.90 were obtained for cyclic loading tests in corrosive and inert environments, respectively (Fig. 6b). Both lines intersect at a frequency of 8.4 102 cycles/s. Thus, the corrosive environment had no effect on crack growth at frequencies above 8 102 cycles/s. Fractographic examinations of the 7050-T7X51 specimens cyclically loaded under inert environmental conditions revealed again cleavage facets at the rim of the fracture surface (Fig. 7). No striations were observed. Fig. 8 shows the fracture surface of a 7050-T7X51 sample cyclically tested in substitute ocean water. Pronounced brittle transgranular crack propagation was observed. On the flat transgranular facets, coarse striations consisting of
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Fig. 4. Scanning electron fractographs of alloy 7050-T7651, showing cleavage-like fracture. The notched specimens were cyclically loaded under inert environmental conditions at frequencies of (a) 5.1 104 and (b) 1.3 103 cycles/s applying a stress range Dr = 250 MPa at a maximum stress rmax = 500 MPa.
deep slots were found. The distance between the crack arrest markings increased with decreasing frequency. Fig. 9 displays results of cyclic loading tests for alloy 7050-T7351 at R-ratios of 0.83 and 0.67 with rmax = 600 MPa. In the frequency range used, number of cycles to failure were higher for the higher R-ratio, but the slope of a linear plot on a log–log basis was slightly steeper compared to that for R = 0.67. However, cyclic loading tests carried out at different R-ratios did not indicate a marked dependence of the slope upon the R-ratio [20]. Fractographic examinations revealed cleavage facets with coarse striations (Fig. 10). The striation spacing increased with increasing stress range; it decreased with higher frequency. At frequencies above 103 s1, striations were difficult to find. Results of cyclic loading tests carried out with notched specimens of alloy 7050-T651 are presented in Fig. 11, applying stress ranges Dr of 100 and 200 MPa with rmax = 600 MPa. At higher frequencies (corresponding to a displacement rate of 3.2 105 mm s1), the number of cycles to failure were similar to those observed for alloy 7050T7351. Applying a displacement rate of 2.5 106 mm s1 (resulting in frequencies below 5 105 s1), the 7050-
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Fig. 5. Scanning electron fractographs of alloy 7050-T7651, showing brittle transgranular fracture. Notched specimens were cyclically loaded in substitute ocean water at frequencies of (a) 1.9 105, (b) 4.6 105, (c) 2.4 104, and (d) 8.0 104 cycles/s applying a stress range Dr = 250 MPa at a maximum stress rmax = 500 MPa.
T651 specimens failed after lower number of cycles. This was associated with the occurrence of intergranular stress corrosion cracking, as corroborated by fractography. On the fracture surfaces of samples cyclically tested at a displacement rate 2.5 106 mm s1, both intergranular and transgranular environment-induced cracking were observed, whereas brittle transgranular fracture was more pronounced when 7050-T651 specimens were loaded at higher frequencies (Fig. 12). Both coarse slot-like and fine striations were observed (Fig. 13). The number of cycles to failure as function of the frequency is plotted in Fig. 14 for all cyclic loading tests conducted with alloy 7050 applying a stress range Dr = 200 MPa with rmax = 600 MPa. At an applied R-ratio of 0.67, the frequency dependence of failure caused by transgranular environment-induced cracking was found to be similar for alloy 7050 in all tempers studied. On a log–log basis, the dependence was linear described by log Y (number of cycles to failure) = 0.596 log X (frequency) + 4.363. Therefore, a unique mechanism causing brittle transgranular environment-induced cracking seems to operate in alloy 7050, being not influenced by the heat treatment conditions. 4. Discussion Stress corrosion cracking of high strength aluminium alloys is a well known degradation process in loaded struc-
tural components described in several review papers [1–4]. The crack propagates predominantly intergranular, and the SCC sensibility can be reduced by artificial aging. Two main mechanisms have been proposed for stress corrosion cracking in aluminium alloys: highly localized anodic dissolution of the grain boundaries and embrittlement due to hydrogen ingress into the alloy. When severe loading conditions are applied, such as dynamic straining to fracture (slow strain rate testing) or stress concentration effects associated with pitting in statically loaded smooth tensile specimens, transgranular stress corrosion cracking can also occur in aluminium alloys, revealing cleavage-like fracture and crack arrest markings [2,4,15–19,29]. The mechanisms proposed for the latter type of environmentinduced cracking include a combined action of anodic dissolution and extremely localized ductile fracture, an adsorption-induced localized-slip process, or the corrosion enhanced plasticity model [19,30,31]. Fatigue crack growth in aluminium alloys is generally transgranular and, under inert environmental conditions, related to alternating shear at the crack tip and controlled by the crack tip opening displacement [32]. In alloy 7150T651, the fatigue crack growth rate in air was found to be independent of frequency in the range between 0.1 and 20 Hz [14]. Ductile transgranular crack propagation was predominantly observed on the fracture surfaces. When fatigue testing was performed in an aggressive aqueous
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Fig. 6. Curves of (a) cycles to failure and (b) time to failure vs. frequency for alloy 7050-T7X51. Notched specimens were cyclically loaded in substitute ocean water and in an inert environment applying a stress range Dr = 200 MPa at a maximum stress rmax = 600 MPa.
environment, the crack growth rates increased by up to an order of magnitude compared to those measured in laboratory air. Fractography revealed both brittle intergranular and trangranular fracture modes, being dominant at low and high DK, respectively. The extent of intergranular fracture mode and the enhancement of fatigue crack growth rates depended upon cyclic frequency, supporting a mechanism of hydrogen grain boundary diffusion. The crystallographic and flat transgranular crack growth being operative at high DK exhibited little ductility and was characterized by the presence of brittle striations. Transgranular corrosion fatigue was proposed to be dependent on strain-controlled hydrogen diffusion ahead of the crack [14]. Similar results were found by Holroyd and Hardie conducting fatigue crack growth tests of alloy 7017-T651 in seawater [13]. At all frequencies tested, the crack growth rates significantly increased in seawater, compared to dry air, and the enhancement was more marked at the lower frequencies. Crack growth rates in seawater depended upon both loading frequency and DK. An increase of these
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Fig. 7. Scanning electron fractographs of alloy 7050-T7X51, showing cleavage-like fracture. Notched specimens were cyclically loaded under inert environmental conditions at frequencies of (a) 1.7 103 and (b) 1.0 103 cycles/s applying a stress range Dr = 200 MPa at a maximum stress rmax = 600 MPa.
parameter resulted in a change of fracture mode from intergranular to flat transgranular to striated ductile transgranular. A linear fit was obtained plotting logarithmically the crack velocities associated with the fracture mode transitions vs. frequency of loading. The dependence of the velocity involved in each transition on the square root of the cycle period confirmed again that the enhanced crack growth in seawater was controlled by diffusion of hydrogen ahead of the crack tip. For alloy 7050-T7651 cyclically loaded under inert environmental conditions, the time to failure tF decreased linearly with increasing frequency m on a log–log basis with a slope of 1 (Fig. 3b). Therefore, the number of cycles to failure nF, given by nF = tF m, was constant being independent upon frequency. When the cyclic loading tests were carried out in substitute ocean water, a slope of 0.5 was obtained fitting the logarithmic plot of time to failure vs. frequency. The number of cycles to failure was dependent on p the square root of frequency, i.e. proportional to 1/ s with s being the cycle period. This square root dependence indicated an embrittlement mechanism controlled by diffusion. Thus, the enhanced crack growth in substitute ocean water was related to hydrogen diffusion ahead of the crack tip, in accordance with the suggestion made by Holroyd and Hardie [13].
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Fig. 8. Scanning electron fractographs of alloy 7050-T7X51, showing (a) brittle transgranular fracture and (b) crack arrest markings. The notched specimen was cyclically loaded in substitute ocean water at a frequency of 2.3 105 cycles/s, applying a stress range Dr = 200 MPa at a maximum stress rmax = 600 MPa.
Fig. 9. Number of cycles to failure vs. frequency for alloy 7050-T7351. Notched specimens were cyclically loaded in substitute ocean water applying stress ranges Dr of 100 and 200 MPa with rmax = 600 MPa.
When the maximum stress was increased to 600 MPa, the slope of the time-to-failure vs. frequency curve plotted logarithmically was 0.90 for notched specimens cyclically loaded in an inert environment (Fig. 6b). The maximum load exceeded significantly the yield strength of the alloy 7050-T7X51, resulting in a slight dependence of the num-
Fig. 10. Scanning electron fractographs of alloy 7050-T7351, showing (a) brittle transgranular fracture and (b) crack arrest markings. The notched sample was cyclically loaded in substitute ocean water at a frequency of 3.1 104 cycles/s applying a stress range Dr = 200 MPa with rmax = 600 MPa.
Fig. 11. Number of cycles to failure vs. frequency for alloy 7050-T651. Notched specimens were cyclically loaded in substitute ocean water applying stress ranges Dr of 100 and 200 MPa with rmax = 600 MPa.
ber of cycles to failure upon the frequency. Besides hydrogen embrittlement this mechanical effect also contributed to the failure of samples immersed in an aggressive environment, leading to a slope of 0.42. Fatigue crack growth in alloy 7075-T651 was found to be intermittent, not proceeding cycle-by-cycle both in moist
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Fig. 12. Scanning electron fractographs of alloy 7050-T651, showing (a) intergranular and transgranular and (b) transgranular environmentinduced cracking. The notched samples were cyclically loaded in substitute ocean water at frequencies of (a) 4.8 105 and (b) 6.4 104 cycles/s applying a stress range Dr of 100 MPa with rmax = 600 MPa.
air and vacuum [33]. Before the crack propagates, it blunts and accumulates local strain damage near the tip during a number of cycles which decreases with higher cyclic stress intensity. The plasticity required for an increment of crack extension was reduced in moist air compared to that in vacuum. Whereas fracture surfaces resulting from purely plastic processes in inert environments were non-crystallographic and relatively featureless, crystallographic fatigue crack growth on {1 0 0} and {1 1 0} planes was observed in moist environments [11,12,14]. In the present study, cleavage facets were found at the rim of the fracture surfaces of notched specimens cyclically loaded under inert environmental conditions (Figs. 4 and 7). This might be caused by a residual humidity of the laboratory air not entirely removed by embedding the specimens in magnesium perchlorate hydrate. However, slow strain rate tests carried out to evaluate the SCC behaviour of aluminium alloys indicated that this wrapping procedure can be used for reference tests in an inert environment [28,34]. Therefore, the crystallographic facets were rather attributed to slip processes at the crack tip activated particularly by the high maximum load. When specimens were immersed in substitute ocean water hydrogen ahead of the crack tip promoted slip, thus enhancing crack advance. The failure of the specimens cyclically loaded under inert environmen-
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Fig. 13. Scanning electron fractographs of alloy 7050-T651, showing (a) coarse and (b) fine striations. The notched sample was cyclically loaded in substitute ocean water at a frequency of 6.4 104 cycles/s applying a stress range Dr of 100 MPa with rmax = 600 MPa.
Fig. 14. Number of cycles to failure vs. frequency for alloy 7050 in different tempers. Notched specimens were cyclically loaded in substitute ocean water applying a stress range Dr = 200 MPa at a maximum stress rmax = 600 MPa. The log values were linearly fitted.
tal conditions might be also associated with accumulation of plastic strain caused by cyclic creep/ratcheting when stress levels close to or exceeding the ultimate tensile strength were applied [35,36]. However, as found with SAE 1045 steel, fractographic examinations of samples
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failed due to cyclic creep/ratcheting revealed ductile dimple fracture resulting from microvoid coalescence [37]. Thus, the cleavage fracture observed at the rim of alloy 7050 specimens was related to fatigue cracking. Cleavage facets were not observed with notched specimens monotonically tensile tested in an inert environment. Striations are typically found on corrosion fatigue fracture surfaces [11,12,14,33]. Several types of striations were observed varying from a sawtooth shaped profile (type A) to another one consisting of deep slots traversing a flat surface (type B) [11]. It was concluded that the striation profile depended upon the orientation of the cleavage plane and the slip systems to the tensile and shear stresses. In the present study, the flat transgranular facets of the samples cyclically loaded in substitute ocean water were covered with striations. These crack arrest markings indicated an intermittent crack growth in the corrosive environment. The striations were predominantly regularly arranged and consisted of slots (Figs. 8, 10 and 13), being similar to the type B and brittle striations observed by Nix and Flower and Gao et al., respectively [11,12]. This type of striation profile occurred when the cleavage plane was oriented orthogonal to the tensile stress axis [11]. Occasionally, fine regularly spaced striations were also found related to crack advance during a cycle (Figs. 5a and 13b), resembling the ductile striations observed by Gao et al. on fracture surfaces of 7075-T651 specimens tested in water vapour [12]. The latter type of striations was probably created by dissolution of the exposed bare fracture surface at crack arrest during the period of reduced load. With increasing frequency, the regularly arranged striations disappeared, and irregularly shaped markings were observed near the transition region to overload fracture (Figs. 5d and 12b). These markings might be secondary cracks related to the high final stress before failure. Because these cracks did not occur on the cleavage fracture of specimens cyclically loaded in an inert environment, they were also associated with the corrosive environment, as confirmed by their presence on the fracture surfaces of specimens failed under static loading conditions (Fig. 2b). Whereas the intergranular SCC susceptibility of 7000series aluminium alloys depends upon heat treatment [1], sensitivity to transgranular environment-induced cracking was found to be independent upon temper condition, at least when a maximum stress of 600 MPa was applied which was close to the notch strength of the specimens used. This was confirmed by the occurrence of transgranular cleavage-like fracture in alloy 7050 in all tempers as well as the single straight line fitting the time to failure data of all notched specimens cyclically loaded at a stress range Dr = 200 MPa with rmax = 600 MPa (Fig. 14). Therefore, the diffusion process of hydrogen causing embrittlement was not significantly influenced by changes of the microstructure due to artificial aging. Because fractographic features of transgranular environmentally assisted cracking were similar when specimens were stressed under static, dynamic and cyclic loading conditions [20,29], the same
transgranular embrittlement mechanism might be operative in both stress corrosion cracking and corrosion fatigue at low frequencies. 5. Conclusions 1. Alloy 7050 was found to be sensitive to transgranular environment-induced cracking when cyclically loaded in substitute ocean water at frequencies between 2 105 and 4 103 cycles/s. 2. The number of cycles to failure decreased with decreasing frequency. On a log–log basis, a linear dependence was determined with slopes of 0.5 and 0.6 applying maximum stresses of 500 and 600 MPa, respectively. 3. Striations on the transgranular facets of the fracture surfaces indicated discontinuous crack propagation. Crack arrest markings were more pronounced when cyclic loading was carried out at lower frequencies. 4. No influence of the heat treatment on transgranular environment-induced cracking of alloy 7050 was observed. 5. Environment-induced failure under cyclic loading conditions at low frequencies was probably associated with hydrogen embrittlement.
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