Materials Science and Engineering, A 125 ( 1990) 1 - 14 l

Effects of Environment and Microstructure on Stress Corrosion Crack Propagation in an AI-Li-Cu-Mg Alloy

M. AHMAD

Department of Mechanical Engineering, Linkiiping Institute of Technology, S-581 83 Linkrping (Sweden)

(Received September 1, 1988; in revised form August 28, 1989)

Abstract

Stress corrosion cracking behaviour of the 8090 (Al-Li-Cu-Mg-Zr) alloy has been studied using fatigue-precracked double-cantilever-beam (DCB) specimens. The corrosion environment was either 3.5% NaCI solution or synthetic seawater. Total stress corrosion crack growth in the DCB speci- mens exposed to the synthetic seawater environ- ment was generally greater than that for the 3.5% NaCI solution, Grain boundary microstructures of the 8090 alloy in a variety of aging conditions have been studied using transmission electron micros- copy. Precipitation of l phase and 6 (AILi) phase on high angle grain boundaries resulted in 6' (AI~Li) precipitate-free zones along the bound- aries. Stress corrosion cracking resistance of the alloy is discussed in terms of grain boundary microstructure and corrosion environment.

I. Introduction

Lithium-based aluminium alloys are well known for their low density and high stiffness as compared to the conventional aerospace alloys. However, the resistance to stress corrosion crack- ing (SCC) in the short transverse direction, par- ticularly in thick sections, may be a problem. There are numerous investigations dealing with the problem of stress corrosion cracking of these alloys [1-6]. Early work by Rinker et al. [1] on alloy 2020 have shown that aging the alloy from the underaged to the peak-aged condition results in significant reductions in the plateau velocity. This change in SCC behaviour was attributed to the potential difference between grain boundary T 1 precipitates and the grain interior. Christo- doulou et al. [2] studied the stress corrosion cracking of A1-Li binary alloys. It was suggested that hydrogen embrittlement may play a role in the SCC mechanisms. Holroyd et al. [4] investi-

gated the SCC behaviour of binary A1-Li, ternary AI~Li-Zr and quaternary A1-Li-Cu-Mg alloys. Various mechanisms possibly involved in the stress corrosion cracking of aluminium-lithium alloys were discussed. Work by Vasudevan et al. [5] on stress corrosion cracking of A1-Cu-Li-Zr alloys has shown that aging had little effect on plateau velocities, and reduction in Kiscc values with aging was related to a concomitant reduction in toughness. The SCC behaviour of the 8090 alloy [7-12] and the 2090 alloy [13, 14] have also been the subject of more recent publications.

The present investigation was undertaken to study the effects of complex aging heat treatments on the fracture toughness and stress corrosion resistance of the 8090 alloy. Both smooth and precracked specimens were used to study the stress corrosion behaviour of the alloy. The results from the tuning-fork-type smooth test specimens, fracture toughness and early double- cantilever-beam (DCB) test results have been published elsewhere [7, 8]. Synthetic seawater was chosen as a corrosion medium since it is known that chloride, bicarbonate and sulphate ions can b e severe SCC environments for the AI-Li-Cu-Mg alloys [15, 16].

2. Experimental procedures

A plate of the 8090 alloy 44 mm thick was obtained from British Alcan Aluminium. The as- received plate was in the peak-aged condition, i.e. aged for 16 h at 190 °C. The chemical composi- tion of the alloy plate is shown in Table 1.

The plate was sawn along the rolling direction into bars with cross-sectional dimensions of 44 mm × 28 mm. These bars were solution heat treated in an air furnace for 1 h at 530°C, quenched in room temperature water, stretched 2.5% and aged. After quenching but before the,

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2

TABLE 1 Chemical composition of 8090 alloy plate

Element Li Cu Mg Zr Fe Si AI Amount (wt.%) 2.36 1.08 0.67 0.11 0.05 0.03 Balance

beginning of aging heat treatments, the alloy bars were kept in a refrigerator in order to limit natural aging. Aging heat treatments were nor- mally carried out in an oil bath. A 4 min, 250 °C heat treatment was performed in a salt bath.

The microstructure of the alloy in a variety of aging heat-treated conditions was studied using transmission electron microscopy (TEM) (Phitips EM400). Thin foils for TEM work were pre- pared using a twin-jet polishing apparatus with a solution of nitric acid and methanol (1:2) at -15 °C.

The stress corrosion resistance of the alloy in the various aging tempers was evaluated using precracked DCB specimens. The DCB speci- mens, 25 m m x 25 mm x 130 mm, were ma- chined from the centre of the plate thickness and oriented so that cracking would occur in the short transverse plane and propagate in the rolling direction. Further details about the geometry of the DCB specimens have been published else- where [7]. Specimens were fatigue precracked at 10 Hz and the maximum stress intensity was less than 80% of K~c. Fatigue cycling continued until a crack had grown 2-2.5 mm beyond the end of the 90 ° chevron notch. A clip gauge was attached and the specimens were bolt loaded to predetermined starting stress intensity levels. The DCB tests were performed in a climatized room at a temper- ature of 20 °C.

As a corrosion medium, 3.5% NaCl solution and synthetic seawater were chosen. The syn- thetic seawater was prepared according to the ASTM specification D 1141-75. Specimens were exposed to 3.5% NaCl solution by applying drops at two different frequencies, once an hour and three times daily. Specimens were also exposed to synthetic seawater by either alternate immersion or moistening by drops three times daily. Alter- nate immersion cycling was such that the lower portions of the DCB specimens were soaked in the solution for 10 min followed by 50 rain in air. The loading steel bolts were always above the water surface level. The alternate immersion cycle and dropwise addition of the solutions to the notch were accomplished with the aid of electronic timers. Crack lengths were measured optically (x 500 magnification) with the help of a

TABLE 2 Results of fracture toughness tests [71

Specimen Aging treatment of Vickers Ksa a number 8090alloy hardness (MPa

(10 kP) m 112 )

1 4 h, 190 °C 130.0 33.0 2 16 h, 190 °C 142.9 16.0 3 80 h, 190 °C 139.3 8.9 5 190 h, 130 °C 133.1 34.8 6 50 h, 130 *C+ 16 h, 190 °C 137.1 19.2 7 4 min, 250 °C+ 16 h, 190"C 141.7 13.7 8 As received (16 h, 190 °C) 142.5 12.9 8a As received (16 h, 190 °C) -- 23.7 9 7075-T651 -- --

aFracture toughness test orientation was S-L, except for no. 8a (T-L). S-L and T-L are defined according to ASTM E399.

travelling microscope on both sides of the speci- mens and the measurements averaged. The fol- lowing equation was used to calculate the stress intensity K~ [17]:

31/2E(2V) KI 4Hl/Z(a/H+O.6728)z (1)

where E is the modulus of elasticity (79 GPa for the 8090 alloy and 71.1 GPa for 7075-T651 alloy), 2 V is the deflection of the specimen arms at the load line, H is half the specimen height and a is the crack length measured from the load line.

At the end of testing, the DCB specimens were unloaded and the remaining displacement was measured. Finally, specimens were reloaded in a tensile testing machine until complete fracture occurred. Load-displacement plots were recorded during reloading to fracture.

3. Resul ts

The various aging tempers included in the present investigation are listed in Table 2. Vickers hardness and short-bar fracture toughness (KsB) data of the alloy tempers are also presented in Table 2. This is a continuation work, the first part of it having been published elsewhere [7].

3.1. Microstructure The transmission electron micrographs in Figs.

l(a)-l(l) show the development of 6' (AlsLi) precipitate-free zones (PFZs) along grain bound-

aries with aging. Both the bright field and corre- sponding dark field electron micrographs taken using a 6' superlattice reflection are shown in Fig. 1. The alloy sample in the low temperature underaged condition (190 h, 130°C) showed no 6' PFZ (Fig. l(b)). The precipitation density of 6' was greater in the alloy aged at the low tempera- ture than in those aged at the high temperature. In

addition, the size of 6' particles and 6' PFZs increases with aging. The other phases present in the alloy are S' (AI2CuMg), T 1 (A12CuLi) and t3' (A13Zr) [18, 19, 8]. The growth mechanisms of the 6', T 1 and S' phases have been analysed previ- ously [20, 19, 8]. For the aging conditions studied, the 6' PFZs were present only along high angle grain boundaries. Electron micrographs (g) and

~, 4-- P F Z

Fig. 1 (continued).

(h) in Fig. 1 show a high angle grain boundary and a couple of low angle grain boundaries; no 6' PFZs are present along the low angle boundaries.

The typical transmission electron micrographs in Figs. 2-4 show precipitation on high angle grain boundaries. Figures 2(a) and 2(b) represent high angle grain boundary precipitation in the alloy samples aged at the low temperature (190 h, 130 °C). Apart from the I phase, few T 1 precipi- tates nucleated on the high angle grain boundary. Historically the icosahedral (I) phase has been

called T 2 (Al6CuLi3). This phase has received attention from a number of researchers [21-24]. In Fig. 2 numerous spherical fl' (AI3Zr) disper- soid particles are also visible within the matrix and at grain boundaries. Figure 3 shows precipi- tation of I phase on a high angle grain boundary of high temperature underaged temper (4 h, 190 °C). The number density of high angle grain boundary precipitates in the high temperature underaged condition is much greater than that observed in the low temperature underaged

PF Z --b// t

Fig. 1. TEM micrographs illustrating the development of 6' PFZs along the high angle grain boundaries, bright field (left) and corresponding 6' dark field micrograph (right): (a), (b) low temperature underaged alloy (190 h, 130 *C)--no 6' PFZ; (c), (d) high temperature underaged alloy (4 h, 190 °C); (e), (f) peak-aged alloy ( 16 h, 190 °C); (g), (h) overaged alloy (80 h, 190 °C)--note no 6' PFZ along the low angle boundaries; (i), (j) aged (4 min, 250 *C + 16 h, 190 °C); (k), (I) as-received ( 16 h, 190 °C).

material. Figures 4(a) and 4(b) show the high angle grain boundary microstructure of the alloy material aged to peak (16 h, 190°C) and over- aged (80 h, 190 °C) tempers respectively. In addi- tion to the I phase, large equilibrium 6 (A1Li) phase particles are also present on high angle

boundaries. A large particle (Fig. 4(b) (A)), pos- sibly c phase, is also seen. Massive precipitation of S' phase within the matrix is also obvious from Fig. 4. Owing to the reactive nature of 6, this phase is preferentially attacked during electro- polishing of thin foils, which makes it difficult to

Fig. 2. TEMmicrographs showing precipitation along high angle grmn boundaries oflow temperature underaged alloy(190 h, 130°C).

~i !ii!~!ill iili~ !i ~!ii !i! i%111~i iii!iii!! !!~i

iiii!iiii ii iiiiiiiiiiiiiiii !!iiii!!ii!i!i!!ii!ilili

Fig. 3. TEM micrograph showing precipitation along high angle grain boundaries of high temperature underaged alloy (4 h, 190 *C).

obtain diffraction information from the particles. These panicles are normally surrounded by dislocations owing to the large difference in the lattice parameters of 6 and the matrix [25]. The strain field surrounding the 6 panicles can be seen e.g. in Fig. 4(c). The electron micrographs in Figs. 4(c) and 4(d) are representative of high angle grain boundary precipitates of the alloy

samples aged (4 min, 250 °C + 16 h, 190 °C) and in the as-received condition respectively. Grain boundaries of the alloy with the two-step aging condition (no. 7) were severely attacked during electropolishing of thin foils. In addition, precipi- tation on grain boundaries was more pronounced in the two-step aged condition as compared to the other aging heat treatments.

The size of equilibrium grain boundary par- ticles 6 and I phase increased with aging. How- ever, it seems that the 6 phase coarsens more rapidly as compared to the I phase. This is prob- ably due to the fact that the diffusion of copper in the alloy is slow as compared to that of lithium [20], and the growth of I phase is controlled by copper diffusion while the growth of 6 (A1Li) phase is controlled by lithium diffusion.

A high angle grain boundary of the alloy in the overaged condition is shown in Fig. 5(a). Conver- gent beam electron diffraction (CBED) patterns of a panicle (Fig. 5(a) (A)) are shown in Fig. 5(b) (matrix zone axis [110]) and Fig. 5(c) (matrix zone axis [111]). This particle may be the c phase reported by Lapasset and Loiseau [26]. These types of particles were seen only rarely. However, the more common grain boundary particle I phase displays icosahedral symmetry. Figure 6 is a CBED pattern of a single I phase particle dis- playing fivefold symmetry. A selected area dif- fraction (SAD) pattern with the matrix zone axis [111] is shown in Fig. 5(d).

Fig. 4. TEM micrographs showing precipitation along high angle boundaries in the 8090 alloy: (a) peak aged (16 h, 190 °C); (b) overaged (80 h, 190 °C); (c) aged (4 rain, 250 °C + 16 h, 190 °C); (d) as-received condition ( 16 h, 190 °C).

Figure 7 shows S' phase precipitating on a low angle grain boundary and within the matrix of the low temperature underaged alloy. Precipitations of the S' and T~ phases within the matrix of the alloy in the underaged tempers are compared in Fig. 8. It is clear that the T~ and S' precipitation is less advanced in the alloy aged at the low temper- ature (190 h, 130°C, Fig. 8(a)) as compared to the high temperature (4 h, 190 °C, Fig. 8(c)). The corresponding [112] SAD patterns are shown in

Figs. 8(b) and 8(d). The SAD micrographs show streaking along (210) aluminium directions caused by S' precipitates. It is seen that the (210) streaks are weaker in the alloy aged at the low temperature (Fig. 8(b)) as compared to the high temperature (Fig. 8(d)).

3.2. Stress corrosion crack growth Table 3 summarizes the stress corrosion

crack growth data of fatigue precracked DCB

i i i ! i ¸ iiii! iiii!iiii i ¸

Fig. 5. (a) High angle grain boundary of overaged alloy (80 h, 190 °C). (b) CBED pattern from grain boundary precipitate "A" in (a), matrix zone axis [110]. (c) CBED pattern from precipitate '~" in (a), matrix zone axis [111]. (d) SAD pattern, matrix zone axis[Ill].

specimens exposed to 3.5% NaC1 solution and synthetic seawater environments. It contains the stress intensity factor to which the specimens were initially loaded, the final stress intensity factor calculated from eqn. (1), the total exposure time and the total crack growth. The total crack

growth values within parentheses are based on the average of crack surface measurements. The crack lengths were measured at the following three positions: at the centre of the crack front and at midpoints between the centre and the ends

of the crack front. The crack growth values based on the fracture surface measurements were usually greater than those measured at the speci- men edges because of the irregular nature of the crack fronts. Owing to pitting corrosion attack on the surface, the crack growth measurements in the 7075-T651 alloy exposed to the 3.5% NaC1 solution were terminated after 83 days testing since it was difficult to measure crack lengths accurately. Total crack growth is generally lower in the specimens exposed to 3.5% NaCI solution

Fig. 6. CBED pattern from I phase precipitate showing five- fold symmetry axis.

Fig. 7. TEM micrograph showing S' precipitates on a low angle boundary and within matrix of low temperature under- aged alloy (190 h, 130 °C).

than in those exposed to synthetic seawater. The DCB specimen numbers with index R had been previously SCC tested and resultant crack growth was up to 1 mm. They were re-fatigued and loaded to the stress intensity levels shown in Table 3.

Stress corrosion crack growth vs. exposure time data are also plotted in Figs. 9-11. The aging heat treatment, corrosive environment and starting stress intensity factor applied to the respective DCB specimens are listed in the figures. By comparing the equivalent alloy tem- pers and starting stress intensity levels, it can be seen from Figs. 9(a) and 9(b) that crack growth in the alloy specimens alternately immersed in syn- thetic seawater is faster and greater than in those specimens exposed to 3.5% NaCI solution.

The crack growth vs. exposure time data for the underaged alloy specimens alternately immersed in synthetic seawater are shown in Fig. 10. The specimen aged at the low temperature (190 h, 130 °C) shows far lower crack growth as compared to the one aged at the high temperature (4 h, 190 °C). Both of the underaged specimens had almost the same hardness and fracture tough- ness and were loaded to the same starting stress intensity. Comparison of the alloy specimens either alternately immersed in synthetic seawater (no. 8.1, Fig. 9(b)) or periodically moistened three times per day with synthetic seawater (no. 8.2, Fig. 11) shows that crack growth was faster for the alloy specimen moistened periodically. One of the specimens two-step aged (50 h, 130°C+ 16 h, 190°C) and alternately immersed in syn- thetic seawater showed an incubation period of 50 days (Fig. 9(b)). In a previous study on stress corrosion cracking of alumim'um-lithium alloys, Pizzo and Bell [27] also observed an incubation period of 10-60 days before the subcritical crack growth starts. A more detailed examination of the stress corrosion crack growth curves in Figs. 9-11 reveals a more or less stepped behaviour for some of them. The extent of self-loading caused by the wedging action of corrosion prod- ucts may be estimated from (1) the remaining displacement after unloading the specimen [28] and (2) the load-deflection plot obtained during reloading to fracture [29]. The load-displacement plot was non-linear for specimens 2.1, 3R, 6.1 and 8R. The corresponding remaining displace- ment values were also greater for the above specimens as compared to the others. This indi- cates that pronounced stepwise crack growth in specimens 6.1 and 2.1 (Fig. 9(b)) and 8R (Fig. 11) was due to the corrosion product wedging effects. The crack growth data (Figs. 9-11) and eqn. (1) were combined to obtain crack growth rate vs. stress intensity plots. It is apparent from Fig. 12 that the crack growth rate at a given stress

10

Fig. 8. TEM micrographs showing precipitation of S' and T I phases within matrix of alloy: (a) low temperature underaged ( 190 h, 130 °C); (b) corresponding SAD pattern of (a), [112] beam direction; (c) high temperature underaged (4 h, 190 °C); (d) corre- sponding SAD pattern of (c), [ 112] beam direction.

intensity increases and the threshold value decreases when the 8090 alloy is aged from the underaged (no. 1.1) to the peak-aged (no. 8R) condition.

The high temperature underaged alloy speci- men (no. 1.1) alternately immersed in synthetic seawater showed a rather strange crack mor- phology. Instead of growing continuously from the fatigue precrack, the stress corrosion crack started almost 2 mm away from the mid-thick- ness of the DCB specimen. However, with contin-

ued exposure the crack propagated towards the mid-thickness. Multiple parallel stress corrosion cracks were quite common on the DCB specimen surfaces.

4. Discussion

It is interesting to compare the stress corrosion crack growth behaviour of the underaged alloy tempers. The alloy aged at the low temperature (190 h, 130°C) had almost the same (actually

TABLE 3 Stress corrosion crack propagation data of fatigue-precracked DCB specimens

l l

Specimen Aging treatment of Environment Kt ~ Kf b Total Total number 8090 alloy (MPa m I/2) (MPa m 1/2) exposure time crack growth c

(days) (mm)

1 4 h, 190 °C Wet once per hour, 14.0 13.1 139 1.548 1R 4 h, 190 °C 3.5% NaCI 25.0 20.2 351 4.938 (5.868) 2 16 h, 190 °C 14.0 12.2 469 1.795 (3.660) 3 80 h, 190 °C 7.5 7.4 139 0.250 5 190 h, 130 °C 14.0 13.9 139 0.120 5R 190h, 130°C 25.0 22.2 351 0.229 (3.310) 6 50 h, 130 °C+ 16 h, 190 °C 14.0 12.4 469 1.283 (3.133) 7 4 min, 250 °C + 16 h, 190 °C 12.5 10.4 434 4.308 (4.690) 9 7075-T651 12.6 8.8 83 9.565

1.1 4h, 190°C Alternate immersion 25.0 13.5 351 16.637(17.453) 2.1 16 h, 190°C in synthetic seawater: 13.6 10.1 351 7.443 (7.798) 3R 80 h, 190 °C wet 10 min, 7.5 6.4 351 2.372 (3.995) 5.1 190 h, 130 °C dry 50 min 25.0 22.5 351 3.164 (2.637) 6.1 50 h, 130 °C + 16 h, 190 °C 16.3 11.5 351 8.765 (9.462) 7.1 4 rain, 250 °C + 16 h, 190 °C 11.6 10.3 351 3.547 (3.039) 8.1 As received (16 h, 190 °C) 11.0 9.4 351 3.612 (3.972)

8.2 As received (16 h, 190°C) Wet three times per day, 11.0 8.2 351 5.558 (7.452) 9.1 7075-T651 synthetic seawater 20.0 7.4 351 28.252(31.373)

8R As received (16 h, 190 °C) Wet three times per day, 11.0 7.2 351 9.365(11.965) 3.5% NaCI

~Crack tip stress intensity factor to which the specimens were initially loaded. "Final crack tip stress intensity factor calculated from eqn. ( 1 ). In such calculations, whenever possible, crack length data based on fracture surface measurements were used. CTotal crack growth based on average of measurements on both sides of the specimens at the end of the exposure period (numbers within parentheses are based on fracture surface measurements).

slightly higher) hardness and fracture toughness (Table 2). On the other hand, the difference in SCC behaviour was quite large. Microstructural studies showed that the precipitation of S' within the matrix and at low angle grain boundaries was slightly more advanced in the high temperature underaged condition. More significantly, the density of precipitates at high angle grain bound- aries, i.e. of I phase, was high and they were closely spaced, and a small 6' PFZ was also present, for the alloy aged at the high tempera- ture. In addition, low angle grain boundaries were decorated by S' for both aging treatments. Since the stress corrosion cracks grow intergranularly in the alloy [7], it is logical to think that the increased precipitation on high angle grain boundaries was responsible for accelerating the stress corrosion crack growth in the high temper- ature underaged alloy. In the high temperature underaged condition there are a large number of small, discrete I phase particles. These will tend to deplete copper and lithium at the grain bound- ary PFZ and generate greater potential difference between the grain boundary and the matrix. This apparent difference in the SCC resistance of the low temperature and high temperature underaged

alloy tempers demonstrates that grain boundary precipitation plays a very important role in deter- mining the stress corrosion resistance of the alloy. The present findings are in good agreement with those of Vasudevan et al. [30]. These research workers showed that the grain boundary precipi- tates play a major role in the SCC resistance of A1-Li-Cu-type alloys. The present results are also consistent with those of Lewis et al. [9]. They studied heavy section forgings of 8090 alloy and reported that those aging treatments which reduce the precipitation of I phase can improve the fracture toughness and SCC resistance. The reduction in threshold value (KIscc) upon aging from underaged (4 h, 190 °C) to peak-aged (16 h, 190 °C) tempers (Fig. 12) is probably related to the reduced fracture toughness for longer aging .times (Table 2). A similar trend in the reduction of apparent Kiscc on aging has been observed by Gray [10]. A possible reason for the alloy speci- men moistened three times per day (no. 8.2, Fig. 11 ) exhibiting faster crack growth as compared to the specimen alternately immersed in synthetic seawater (no. 8.1, Fig. 9(b)) might be that for the alternate immersion the crack surface is washed more frequently and corrosion products are

12

B -

6 -

/ ~_e:~vow~..~j .......... ] ~ Ki.(]~8 ) I%~.S. fiOK/laO*C+lab/120"c /

.=?~ _ , ~ o _ .c_+ u ~ _~-~ l " ~.~,_.,~,,a...e.~_ . . . . . . J /

/

i ~.o) ^ . - -

] ~ (la@ / a /

2 41i (1¢o) / (,~,o) / 1 ~ .,,-..'**:+_+-_* ~ : ' = o,~, = = ~ " = ~ "

0 r I l i I I I I l

0 50 100 150 2 0 0 2 5 0 3 0 0 3 5 0 4 0 0 4 5 0

Exposure t i m e (days)

(a)

io

.=.~."-~'~-~ao~O~:c+,~c' x~'£'~.c, ~.~s n) ' 1 ~ < 8 - ~ _z.~,.v_..x,~"c~ewLao'c_, ~-(kz.ej /

41 / / _ o -.,,_ ~__~.,

~ t ~ ~ -~1- = -.--+ = - - .

0 ~ i l I I I I I I 0 50 I00 150 2 0 0 250 300 350 400

Exposure t i m e (days)

(b)

Fig. 9. Stress corrosion crack growth vs . exposure time curves for the 8090 alloy in a variety of aging conditions and the 7075-T651 alloy. Stress intensity levels to which the DCB specimens were initially loaded are shown in parentheses.

30-

25-

15-

~0

10-

5-

wet lOmbv'dl 7 50rain . . . . . - - - -

/ * ~ (~.o1 ,::::::: : : : : ; ; :~ ~ ~ ~ o D o o o o

' 5'0 ~o ~o ~o ~ ~o ~o Exposure t i m e (days)

400

Fig. 10. Comparison of stress corrosiou crack growth in the underaged alloy tempers. The DCB specimens were initially loaded to K~=25 MPa m j/2 and alternately immersed in synthetic seawater.

30- . . . . . . . . . A

gl,-~O.O ) Z( ~"

15- ~ ~..~a~'~

, wlnthet~ ~mv~,er (t l.O)

0 50 100 150 200 250 300 350 400

Exposure t i m e (days)

Fig. 11. Stress corrosion crack growth as a function of exposure time for the 8090 alloy (nos. 8.2 and 8R) and the 7075-T651 alloy. Stress intensity factors (K~) to which the DCB specimens were initially loaded are shown in paren- theses.

1 0 " 8

E v

m l -

.o

0

o

10"9

t 10-1

10"1 i 5

f " 04 ) ' 0 "0 ' ' 4 ) . . . . . . . -B

B . . . . -m- .... No.1.1 I

I N0.SR

10 15 20 25

Stress intensity (MPa4m)

Fig. 12. Stress corrosion crack growth rate as a function of stress intensity for the underaged (no. 1.1) and peak-aged (no. 8R) 8090 alloy tempers and the 7075-T651 (no. 9.1) alloy.

removed, which reduces the corrosion product wedging effect, or it might be difficult to develop a critical environment at the crack tip.

The total crack growth in the alloy specimens exposed to synthetic seawater was generally greater than that for the 3.5% NaC1 exposure. This suggests that the synthetic seawater environ- ment is more severe than 3.5% NaCI solution for stress corrosion crack growth in the alloy. Craig et al. [15, 16], in a recent study on stress corrosion crack initiation in 8090 alloy, have shown that small concentrations of HCO3- and 8042- are able to induce stress corrosion cracking in dilute

chloride solutions. Since all the above ions are present in artificial seawater solution, it is expected that synthetic seawater will be a severe SCC environment for the alloy. The stress corro- sion test results reported by Dorward and Hasse [11] also showed that the total crack growth in the underaged material exposed to a coastal marine environment was much greater than that for the 3.5% NaCI environment. This was explained in terms of corrosion product wedging effects in the marine environment. However, in the present experimental conditions the alternate immersion in synthetic seawater should limit the action of corrosion product wedging; the higher total crack growth observed for the synthetic seawater is due to the severe nature of the SCC environment.

The SCC performance of the alloy specimen pre-aged at the high temperature ( 4 m in, 250 °C + 16 h, 190 °C) was rather poor (Fig. 9(b) and ref. 7). This is consistent with the TEM microstructural study, which showed enhanced precipitation of I phase and 6 phase on grain boundaries, although Gray [10] claimed that overaging via a duplex high-low aging heat treat- ment improved the SCC resistance of 8090 alloy. This improvement in SCC resistance has been attributed to a more uniform precipitation of 6 phase. The exact aging time and temperature employed by Gray [10] as a first step in the duplex aging process are unknown. In the present study the alloy was two-step aged (no. 7) to a nearly peak-aged condition, the first step being 4 min at 250 °C. As a first step in the duplex aging treatment, a wide range of temperatures and times may be used. Obviously, more work is needed to establish the impact of the two-step aging heat treatments on the stress corrosion resistance of the alloy.

5. Summary and conclusions

From the results of the present investigations the following conclusions can be made.

(1) The number density of precipitates nucle- ated along the high angle grain boundaries was low for the low temperature underaged material (190 h, 130 °C).

(2) Precipitation of I phase and & phase on high angle grain boundaries resulted in &' precipi- tate-free zones along the boundaries.

(3) Total stress corrosion crack growth in the alloy specimens alternately immersed in synthetic seawater was generally greater than that for the

13

3.5% NaC1 environment. This suggests that syn- thetic seawater is a more severe SCC environ- ment for the alloy.

(4) Stress corrosion crack growth in the low temperature underaged alloy (190 h, 130 °C) was smaller than in the high temperature underaged alloy (4h, 190°C). This difference has been explained in terms of reduced grain boundary precipitation of I phase at the low temperature.

(5) Aging the 8090 alloy from underaged (4 h, 190°C) to peak-aged (16h, 190°C) tempers results in increased stress corrosion crack growth rates and decreased threshold value.

(6) Total SCC growth in the 7075-T651 alloy was much greater than for the various aging tempers of the 8090 alloy.

Acknowledgments

The author is grateful to Linktping University and SAAB--Scania AB Aircraft Division for the provision of financial support. He is also very grateful to Professor T. Ericsson and Dr. S. Johansson for many stimulating discussions. The 8090 alloy material was kindly supplied by Alcan International Ltd.

References

1 J. G. Rinker, M. Marek and T. H. Sanders, Mater Sci. Eng., 64 (1984) 203.

2 L. Christodoulou, L. Struble and J. R. Pickens, in T. H. Sanders and E. A. Starke (eds.), Aluminiurn-Lithium Alloys II, Metallurgical Society of AIME, Warrendale, PA, 1984, p. 561.

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