Short transverse fatigue crack propagation in 8090 AlLi alloy: re-ageing and environmental effects

MIIlUlI~ EININIEImG ELSEVIER

Materials Science and Engineering A206 (1996) 163 175

A

Short transverse fatigue crack propagation in 8090 A1-Li alloy: re-ageing and environmental effects X. Chen, D. Tromans The UniversiO' (~["British Columbia, Department of Metals and Materials Engineering, 309-6_750 Stores Road, Vancouver, B.C. V6T IZ4, Canada Received 10 May 1995

Abstract

Rising AK fatigue tests were conducted at 30 Hz on S-L oriented 8090 A1-Li alloy plate in the as-received T8771 temper and after a toughness-enhancing re-ageing treatment at 230 °C. Tests were conducted at an R-ratio of 0.5 in desiccated air, distilled water and 1 M A1C13. Attention was directed to the near threshold and low AK (Stage I) region of crack propagation near and below ~ 3 MPa m u2. The fatigue behaviour of the as-received temper was compared to a previous study on the T - L orientation. The investigation included some intergranular corrosion tests in the 1 M A1CI3 environment, plus fractographic, microstructural and analytical electron microscopical analyses. It was found that S-L fatigue followed an intergranular path under all test conditions. Crack rates in the aqueous environments were demonstrably higher in the S-L orientation than in the T - L orientation, but were similar when tested in desiccated air. The re-ageing treatment had no beneficial effect on S-L fatigue behaviour in the aqueous environments and produced a slightly deleterious effect in desiccated air. The grain boundary (GB) regions exhibited Cu solute concentration gradients that influenced intergranular corrosion behaviour and localized corrosion processes during intergranular fatigue. Discussion of fatigue crack propagation processes included the influence of slip reversibility, localized corrosion, hydrogen embrittlement, and solute distribution at the GB.

Keywords: Fatigue behaviour

1. Introduction

A recent study by the authors [1] investigated the effect of several environments on high-frequency (80 Hz) corrosion fatigue cracking of 8090 A1-Li alloy plate in the T8771 temper. Crack propagation was studied in the T - L (long transverse) orientation and it was found that the near-threshold and Stage I behaviours in aqueous environments were affected by S - L orientation (short transverse) grain boundary splitting (delamination), causing the crack tip region to be placed under plane stress conditions. The splitting was attributed to a combination of solute segregation effects associated with the anisotropic grain structure, including localized dissolution, hydrogen embrittlement processes, and a low S - L orientation fracture toughness (Kic). * Corresponding author. 0921-5093/96/$15.00 © 1996 S S D I 0921-5093(95)10008-3

Elsevier Science S.A. All rights reserved

The poor S - L toughness of commercial A1-Li alloys is well recognized. Lynch [2] reasoned that it was related to grain boundary (GB) segregation of lithium, although he was unable to obtain any analytical confirmation, whereas Webster [3] attributed the low toughness behaviour to GB segregation of trace impurities of alkali metals (e.g. Na and K). Based on the premise that low toughness in the 8090-T8771 alloy was associated with segregation, Lynch [2] devised a double-ageing treatment whereby the as-received '178771 temper was subjected to a short-term re-ageing treatment that doubled or almost tripled the S - L fracture toughness with only a small loss in tensile strength. Blankenship and Starke [4] also improved the short transverse fracture toughness by re-ageing treatments similar to those of Lynch [2]. However, they attributed the improvement to a decrease in strain localization arising from a change in the slip deformation characteristics [4]. Lynch [2] improved the S - L toughness of the T8771 temper by re-ageing for 300 s at temperatures that were

X. Chen, D. Tromans / Materials Science and Engineering A206 (1996) 163-175

164

Table 1 Composition and properties of AI-Li 8090-T8771 plate [12] Chemical composition (wt.%) Li

Cu

Mg

Fe

Si

Zr

Na

Other

AI

2.35

1.2

0.64

0.05

0.03

0.11

0.0005

0.07

bal.

Mechanical properties [12] O'y (MPa)

GUTS (MPa)

e (%)

Kic (MPa m I/2)

L-direction, 480 T-direction, 450

L-direction, 540 T-direction, 525

L-direction, 9 T-direction, 12

L - T = 36, T L = 28 S - T = 19, S - L = 14

30-60 °C higher than the normal T8771 ageing temperature of 170 °C, followed by quenching in cold water. These procedures were very similar in principle to those used to control the problem of reversible temper embrittlement (RTE) of low alloy steels, whereby steels embrittled by tempering or slow cooling in the range 375-575 °C have their toughness restored by heating to a higher temperature (e.g. >~600 °C) followed by rapid cooling to below ~ 300 °C [5]. As noted by Lynch [2], RTE in steels is frequently associated with segregation of impurity elements (e.g. Sn, P, Sb and S) to prior austenite grain boundaries during tempering [6-11]. The improvement in toughness after the de-embrittling treatment is then attributed to the redistribution (removal) of the segregant away from the GB. Studies on the S - L fatigue crack propagation behaviour of A1-Li alloys are scarce and even less is known about the effects of re-ageing treatments. Initial considerations suggest that the beneficial effects of reageing on S - L toughness may be accompanied by enhanced resistance to S - L fatigue, but supposed correlations between toughness and fatigue ignore any microstructural changes due to cyclic deformation processes at the propagating fatigue crack tip. Some preliminary studies by Blankenship and Starke [4] have shown that improvements in S - L toughness of the 8090 alloy after re-ageing at 230 °C did not lead to improved fatigue behaviour in moist air. The present study was designed to provide a more complete understanding of the factors controlling S - L fatigue cracking in 8090-T8771 alloy plate in the near threshold and low AK region (Stage I), particularly the influence of environment and the effect of a toughnessenhancing re-ageing treatment. Fatigue tests were conducted in desiccated air and two different aqueous environments. The study was supplemented by fractographic and analytical electron microscopical examinations, together with the effects of re-ageing on aqueous intergranular corrosion behaviour.

2. Experimental procedures 2.1. Materials and heat treatment

The A1-Li 8090 alloy was received as 12.7 mm thick plate in the T8771 condition. It was part of the same plate used for previous T - L fatigue tests by the authors [1] and for mechanical property tests by Morrison and Zhai [12]. The reported chemical composition of the ingot from which the plate was manufactured is shown in Table 1, together with the mechanical properties of the plate [12]. A separate determination of the alkaline trace elements was conducted by International Plasma Lab Ltd, using atomic absorption techniques, and their analyses are listed in Table 2. The grain structure of the plate was anisotropic with grains that were flattened, in the rolling plane (S-plane) and elongated along the rolling direction (L-direction) [1]. The average grain thickness normal to the rolling plane (S-direction) was .~ 7 pm. Average grain diameters in the rolling plane were g 20 pm in the long transverse direction (T-direction) and ~ 70/zm in the L-direction. The as-received T-8771 temper involved a solution treatment at 545 °C, a 7% stretch, and a final ageing treatment at 170 °C for 32 h [2]. Re-ageing of sectioned coupons of the as-received plate was conducted in a molten salt bath composed of 50% NaNO 3 + 50% KNO 3 at 230 °C for 600 s, followed by rapid quenching into cold water at ~ 20 °C. Vickers microhardness (Hv) measurements were conducted across the plate thickness of both the as-received and re-aged materials. Re-ageing had no effect on the grain anisotropy. Table 2 Atomic absorption analyses of as-received 8090-T8771 plate Element

Analysis, wt.% (p.p.m.)

Detectable lower limit

Na K

0.0011 (11 p.p.m.) Not detected

~>2 p.p.m. ~>25 p.p.m.

X. Chen, D. Tromans

Materials Science and Engineering .4206 (1996) 163. 175

S

i

T

Z_____ B = 12.7mm 2 H = 12.7 m m W=49mm

W

--q

P = Load

V P

Fig. 1. DCB specimen geometry and S L crack orientation.

Double cantilever beam (DCB) specimens were used to determine the fatigue and fracture toughness behaviour of as-received and re-aged specimens in the S - L orientation. The crack plane was parallel to the rolling plane (S-plane) and the macroscopic crack propagation direction was parallel to the L-direction [13]. The DCB specimen dimensions are shown in Fig. 1, together with the orientation of the specimen with respect to the S-, T- and L-directions. Planes normal to these directions are referred to in the text as the S-, Tand L-planes respectively. A chevron-shaped starter notch, ~ 6 . 4 mm at the centre and 13 mm at the surface, was introduced parallel to the S-plane along the midplane of the specimen. A ~ 17 mm pre-crack, with a straight crack front, was then produced from the notch by fatigue cracking in air at 30 Hz. The S - L fracture toughness (K~c) was determined by increasing the load on the DCB specimen in a tensile machine until fracture occurred. The load (P) and crack length (a) at fracture were used to calculate the toughness via Eq. (1) [14,151: Pa / K,- ~3.45

+ 2"415Ha

)

mum loads (Pmin) was produced by a rotating, unbalanced mass. This machine was chosen because it had a relatively high loading frequency of 30 Hz and was able to test compliant specimens, such as those with the DCB geometry. The higher frequency ( ~ 80 Hz) electromechanical resonant fatigue machine used previously for T - L testing [1] was unsuitable because it required very stiff specimen geometries. Cracking was monitored and controlled at 22 °C using procedures described in the earlier study [1]. These involved careful cyclic load shedding until no cracking could be detected ( ~< 1 × 10 s ram/cycle), followed by a rising (AK) fatigue test in the chosen environment at an R-ratio (P,,,i,,P ....... ) of 0.5. The (AK) values were calculated from Eqs. (1) and (12) by substituting P,~,~ and P ..... for P: A K = K, .... - K,~,m = (1 - R)Km~x

2.2. Test specimen geometry

(1)

where B is the specimen thickness (12.7 ram), H is the beam height (6.35 ram) and W is the specimen length (49 mm) measured from the loading line. 2.3. Fatigue testing

Fatigue testing was conducted on pre-cracked DCB specimens with a Sonntag machine (Model SF-1-U). Sinusoidal loading between maximum (Pn~) and mini-

165

(2)

Crack length (a) was measured on a metallographically polished external surface with a travelling microscope and was recorded as a function of the number of cycles (N). The average crack growth rate per cycle, d a / d N , was obtained by the secant method [16]. Fatigue tests were conducted under freely corroding conditions in distilled water, 1 M A1CI3 solution (pH 2.0) and desiccated air. The environments were contained in a transparent acrylic cell mounted around the specimen. The top of the cell was left open when testing in the aqueous environments and the fi'ee corrosion potential (Eco,.,.) in the AICI 3 solution was measured with respect to a saturated calomel electrode (SCE). Prior to a fatigue test in desiccated air, the specimen was placed in vacuum ( ~ 1 × 10 5 Tort) for over 60 h to eliminate any possibility of reversible hydrogen embrittlement effects caused by interaction with water moisture in the laboratory air during load shedding [1]. It was then placed in a test cell containing freshly baked silica gel desiccant and the cell was dosed with a rubber sheet. Following fatigue testing, the specimens were broken open by tensile overload. The fracture surfaces were cleaned ultrasonically in ethyl alcohol, dried in an air stream, and examined in a conventional scanning electron microscope (SEM) under secondary electron imaging conditions at 5 keV excitation. 2.4. Corrosion tests

Corrosion tests were conducted on metallographically polished coupons immersed in 1 M A1C13. Metallographic sections were examined by reflected light microscopy to determine the extent of intergranular corrosion. Relatively large ( > 1 ,urn) undissolved grain boundary particles were analysed in an SEM equipped with an energy-dispersive X-ray (EDX) system.

166

X. Chen, D. Tromans / Materials Science and Engineering A206 (1996) 163-175

2.5. Microstructure and microanalysis

10-3

The microstructures of the 8090 alloy in the as-received and re-aged tempers were studied with a conventional transmission electron microscope (TEM) and scanning transmission electron microscope (STEM). A standard tungsten filament electron gun, operating at 200 keV, was used in both microscopy modes. Microanalyses were conducted on electron transparent specimens with EDX and electron energy loss spectroscopy (EELS) systems that were interfaced to the STEM.

3. Results

Desiccated Air 10-4

•-o

(R=0.5, -22 °C, 30 Hz, S-L)

¢

I0 -5

10 -6

Io I I

10_7

3. I. S - L toughness and hardness The short transverse (S-L) fracture toughness of the as-received material was 14.7 MPa m 1/2, close to that reported in Table 1, whereas the re-aged S - L toughness was increased twofold to 29.8 MPa m ~/2. At the same time, the microhardness of the re-aged material was decreased by ~ 8% relative to the as-received plate, the hardness in both cases being uniform across the plate thickness. The properties are summarized in Table 3, where the improvement in toughness produced by reageing at 230 °C was similar to that obtained by Lynch [2]. Furthermore, recognizing that correlations between yield strength (%) and hardness are imprecise, comparison with Lynch's [2] hardness and strength data suggests that the decrease of 13 Hv units ( ~ 8%) after re-ageing corresponded to ~ 9% reduction in yield strength.

3.2. S - L fatigue behaviour Despite the improved toughness obtained by re-ageing, the near threshold fatigue behaviour of the re-aged material was slightly inferior to that of the as-received plate when tested in desiccated air. This is shown in Fig. 2, where the cyclic crack propagation rates, daMN, of the two tempers are compared at AK values ~<3 MPa m ~/2, and in Table 4 where the corresponding AKth values and daMN rates at a AK value of 2.5 MPa m ~/2 are listed. However, tests in the two aqueous environments revealed no detectable difference in fatigue beTable 3 Microhardness and S - L fracture toughness (KIc) of 8090-T8771 plate As-received (T8771)

Re-aged (230 °C for 600 s)

Hardness (Hv)

Kic (MPa rn 1/2) Hardness (Hv)

Kjc (MPa rn I/2)

159

14.7

29.8

146

• * As-received o o Re-aged

10-8

10"9

I

1

I

I

I

I

I

I I I

10

A K ( M P a . m u2) Fig. 2. S-L fatigue crack growth behaviour of as-received alloy and re-aged alloy in desiccated air. haviour between the as-received and re-aged tempers, as shown in Fig. 3 for AK values ~<4 MPa m 1/2. The desiccated air data from Fig. 2 are represented as broken and unbroken lines in Fig. 3, from which it is very clear that fatigue cracking occurred most rapidly in 1 M A1C13. The behaviour in water was a little more complex. For example, Table 4 and Fig. 3 show that the highest AKth was obtained in distilled water, but the cyclic crack rates in this environment very quickly exceeded those in desiccated air as the AK rose above 2 MPa m ~/z. Thus, when compared with desiccated air, distilled water produced slight crack growth retardation near the fatigue threshold and accelerated crack growth at slightly higher AK values, whereas 1 M A1C13 accelerated crack growth under all conditions of testing. As AK increased to values /> 3 MPa m ~/2, Fig. 3 shows that S - L fatigue cracking approached power law behaviour of the type da/dNoc(AK) n (where n ~ 4 ) , similar to that observed in the as-received T - L orientation in aqueous environments at AK~> 5 MPa m 1/2 [1]. Other fatigue characteristics of as-received S - L and T - L orientations are compared in Table 4, where realtime crack velocities, da/dt at a AK value of 2.5 MPa m 1/2 are presented in consideration of the different loading frequencies of 30 Hz and 80 Hz used for the S - L and T - L tests respectively. Orientation comparisons were not made at A K > 3 MPa m ~/2 because the T - L specimens were subject to out-of-plane cracking and subsequent crack growth retardation in desiccated air [1].

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X, Chen, D. Tromans / Materials Science and Engineering A206 (1996) 163 175

Table 4 Stress intensity thresholds (AKth) and near-threshold crack rates ", (da/dN) and (da/dt), of 8090-T8771 plate Orientation

S S S T T T

L L L Lb Lb Lb

Environment

As-received (T8771)

Desiccated air Distilled water 1 M A1CI~ Desiccated air Distilled water I M AICI~

Re-aged (230 °C for 600 s)

AKth ( M P a m I:)

d a / d N ~ (ram/cycle)

da/dt ~ (rams -1)

AKth (MPa m t'2)

da/dN ~ (mm/cycle)

da/dt " (mm s i)

1.6 1.8 0.97 1.7 1.7 1.05

5.8 × 10 3.3 x 10 1.0x 10 2.3x10 3.2x 10 5.0 x 10

1.74x 9.9x 3.0x 1.84x 2.56× 4.0 x

1.3 1.8 0,97

1,4x 10 5 3.3x 10 5 1.0 x l0 4

4.2x 10-4 9.9x 10 -4 3.0 x 1 0 ~

6 ~ 4 * 6 6

10 --4 10 -4 10 -~ 10 4 104 10 4

Crack rate at A K = 2.5 MPa m ~e. b Data from Ref. [1].

In the as-received temper, A g t h values in the S L orientation were within 0.1 MPa m 1'2 of the corresponding T - L values, with the 1 M A1C13 producing the lowest AKth in both orientations. Also, at A K = 2,5 MPa m j/2, the real-time crack velocities (da/dt) in the aqueous environments were always higher for the S - L orientation, being highest in 1 M A1C13. Table 4 shows that the S--L orientation produced da/dt velocities that were ~ 4 times higher than T L velocities in distilled water and ~ 7.5 times higher in 1 M A1CI3, whereas the real time S - L velocities in desiccated air were very similar to those of the T L orientation. These observa-

10-3

(R=0.5, ~22 °C, 30 Hz, S-L) 10-4

ca

10 4

f

10-6

#7

//~

/

" 10-7 "

/~

•

1 M A1C13, Ecorr

•

Distilled Water

1;

--

I

Re-aged:

[ y

A 1 M A I C l 3,Ecorr

Desiccated

Air

I I I

10"8

As-received:

311

I I

/q4

/

/

,~

j,

,,

o

Distilled Water

- - - Desiccated Air 10"9

I

I

I

1

1

I

I

I

I I

10 A K ( M P a . m la)

Fig. 3, S L fatigue crack growth behaviour of as-received alloy and re-aged alloy in distilled water and 1 M A1CI3. Data from Fig. 1 in desiccated air are included as broken and unbroken lines for comparison purposes.

tions were consistent with the greater tendency for S L splitting phenomena to occur during T - L fatigue testing in 1 M A1C13 and for there to be a general absence of S - L splitting in desiccated air [1]. The Ecorr during S - L fatigue testing in 1 M A1CI3 was close to - 0 . 7 8 Vsc E, identical to that observed during T - L testing [1]. This underscored the validity of the comparisons in Table 4 because it showed that the S - L and T - L data were being compared under identical electrochemical testing conditions.

3.3. S - L fractography The fatigue cracks were well behaved and remained in the initial S - L crack plane where they followed a predominantly intergranular path under all test conditions. There was no significant difference in fractography between the as-received and re-aged tempers, but there was a difference in appearance between fatigue surfaces produced in distilled water and those produced in the other environments. Surfaces obtained in distilled water contained a corrosion product that was readily visible in the SEM and, near AK, h, the corrosion product was worn away in places to leave dark-looking features consistent with fretting or rubbing marks produced by opposing crack surface contact (closure effects), as shown near AK, h in Fig. 4(a). In contrast, fatigue surfaces produced in the other two environments were relatively free from solid corrosion products, as shown for desiccated air in Fig. 4(b). The absence of corrosion products in 1 M AIC13 was consistent with the lower pH and higher solubility of the oxide in this environment. Surfaces produced in I M AICI3 and desiccated air mainly exhibited tear ridges associated with the intersection of out-of-plane grain boundaries with the primary S - L intergranular crack plane, as shown in Fig. 4(b). This was confirmed by the fact that the dimensions of the tear ridge patterns were reasonably consistent with the dimensions of the narrow elongated appearance of grains in the S L crack plane. At higher

168

X. Chen, D. Tromans / Materials Science and Engineering A206 (1996) 163-175

the true effect of water was to promote cracking then higher crack rates should have prevailed just above AKth in distilled water, relative to the rates in desiccated air, consistent with the behaviour in Fig. 3. 3.4. Intergranular corrosion behaviour

(a)

(b) Fig. 4. SEM micrographs of intergranular S-L fatigue fracture surfaces near AKth: (a) re-aged alloy in distilled water showing presence of a corrosion product; (b) re-aged alloy in desiccated air. Direction of crack propagation is from right to left of each micrograph.

magnifications it was often possible to detect some larger grain boundary precipitates on the intergranular fracture surfaces, as shown in Fig. 5. It is widely recognized that corrosion products can induce crack closure effects, thereby reducing the effective cyclic stress intensity AKCn and causing crack growth retardation [17]. Such effects were consistent with the appearance of near-threshold fatigue surfaces in distilled water (Fig. 4(a)) and with the observation that the highest AKth occurred in the same environment (Fig. 3). Corrosion product-induced closure effects usually are most noticeable close to AKth [17]. Consequently, at slightly higher AK values, the crack tip opening displacements at R-ratio 0.5 overcame this type of closure and allowed a more natural closure-free comparison to be made between the effects of distilled water and desiccated air on fatigue behaviour. Thus, if

Surfaces parallel to the T-plane of the as-received and re-aged materials were polished metallographically and then exposed to 1 M A1C13 for 170 h in an unstressed condition, after which they were sectioned normal to the T-plane and parallel to the L-plane. In this manner, the extent of intergranular corrosion along the S-plane of the two materials could be compared. The sections showed that the degree of intergranular penetration in adjacent regions was quite variable in each temper with the average value being two to three times larger in the as-received temper than the re-aged condition, as shown in Fig. 6. Based on the deepest penetrations, the maximum intergranular corrosion rates were 1.5 x 10 9 m s - 1 and 9 x 10- lO m s - ~ for the as-received and re-aged tempers respectively. These values were equivalent to fatigue cracking rates of 5 x 1 0 -8 and 3 x 1 0 s mm/cycle at 30 Hz, and were directly comparable to the lowest recorded cracking rates in 1 M AIC13 in Fig. 3, indicating that the lower AKth in the A1C13 environment (see Table 4) was directly attributable to intergranular corrosion processes. Such localized corrosion phenomena may be explained only by associated differences in alloy chemistry between the grain interior and the GB region. More detailed observations of the intergranular corrosion slots in the SEM showed that relatively large ( ~ 1 pm) Cu-containing particles remained uncorroded, as shown by the EDX line scan for CuKc~

Fig. 5. SEM micrograph of intergranular S L fatigue fracture surface of re-aged specimen, obtained near AKth in desiccated air, showing numerous intergranular precipitates on the surface. Directionof crack propagation is from right to left of micrograph.

X. Chen, D. Tromans / Materials Science and Engineering A206 (1996) 163- 175

(a)

(b) Fig. 6. Light micrograph of sectioned specimens after being subjected to intergranular corrosion in 1 M AICI 3 for 170 h: (a) as-received alloy; (b) re-aged alloy. Grain boundary penetration occurred along the S-plane.

radiation in Fig. 7. It is known that Cu-enriched regions in AI Cu alloys are more noble (less active) than Cu-impoverished regions [18]. Consequently, the more active behaviour of grain boundaries may be related in part to localized changes in the distribution of the Cu solute that was present as a normal alloying element in the 8090 alloy (see Table 1). Examination of the profile of the slot in Fig. 7 showed that the propagating tip was no wider than ~ 0.3 ltm, indicating that any local solute concentration gradients responsible for preferential grain boundary corrosion were probably contained within ~0.15 /~m of the boundary. 3.5. Microstructure and solute distribution

Microstructural and chemical composition studies were conducted on thin foils using TEM and STEM techniques. The foils were prepared from the bulk

169

material with the plane of the foil parallel to the T-plane, so that the S - L grain boundaries were normal to the plane of view. All specimens exhibited a very fine intragranular dispersion of metastable 6'-phase (A13Li) precipitates that were 10-30 nm in diameter. No obvious differences in the amount or distribution of the cS'-phase could be detected between the as-received and re-aged tempers, as shown in Fig. 8. More quantitative comparisons were affected by unknown variations in foil thickness between individual TEM specimens, due to larger numbers of 6'-phase precipitates being observed in the field of view of thicker foils, but small changes in precipitate distributions between the two tempers must have occurred in order to account for the decreased hardness of the re-aged alloy (Table 3). Examination of the boundary regions revealed the presence of numerous, larger Cu-containing GB precipitates whose distributions were reasonably similar in both the as-received and re-aged alloys. The GB precipitates were approximately lenticular in shape with a diameter in the boundary plane that was usually of the order of ~ 200 nm (0.2 pm) and a central thickness of 70 nm (0.07 Itm), although larger and smaller precipitates were frequently present. A typical example is shown for the as-received temper in Fig. 9. The overall similarity in TEM microstructure between the two tempers was consistent with the observations of Lynch [2] and Blankenship and Starke [4]. Semi-quantitative EDX analyses were conducted with a narrow 50-70 nm (0.05-0.07 pro) electron beam in STEM mode in order to detect any variations in Cu concentration in the grain boundary region. Concentrations were measured at discrete positions along a line traversing the boundary by controlling the position of the electron beam. The results for both tempers are shown for a traverse through GB precipitates in Fig. 10(a) and for a traverse between GB precipitates in Fig. 10(b). Each point in Fig. 10 represented the average of four separate analyses. Consistently, the GB precipitates showed a significantly higher concentration of copper than the adjacent matrix (Fig. 10(a)L Also, both tempers showed a depletion of copper along the boundary between the precipitates, with an accompanying solute concentration gradient that extended over a distance of ~ 150 nm (0.15 /tm) on either side of the boundary (Fig. 10(b)). The width of the concentration gradient was in striking agreement with the anticipated value based on the dimensions of the tip of the propagating intergranular corrosion slot in Fig. 7, providing strong support for a relationship between intergranular corrosion and copper solute concentration gradients in both tempers. However, it was not possible to determine whether depletion of Cu in the actual boundary plane was significantly different between lhe two tempers because the lower limit of the sampling volume was limited by the dimensions of the electron beam to a diameter of approximately 50 -,70 nm.

170

X. Chen, D. Tromans / Materials Science and Engineering A206 (1996) 163-175

(a)

(b)

Fig. 7. SEM plus EDX line scan analysis of intergranularly corroded slot propagating along the S-plane in the re-aged alloy: (a) SEM micrograph showing corroded slot, uncorroded Cu-containing particle and location of line scan traversing the particle; (b) CuK~ line scan showing peak count rate coincident with position of particle.

Fig. 10(a) appeared to indicate that the re-aged GB precipitates had a higher copper concentration, but selected area electron diffraction patterns of GB precipitates in both tempers were reasonably consistent with T2-type phases of A16CuLi 3 [19] and AlsCuLi 3 [20], corresponding to 26 and 29 wt.% Cu respectively. Consequently, the apparent difference in Cu concentration was probably a sampling problem arising from the relative contributions of the precipitate and matrix to the EDX spectrum. Under these circumstances, sample volumes containing larger precipitates will tend to fill the sampling volume more completely and give a higher Cu signal. Magnesium analyses could not be conducted by the STEM-EDX method because the MgKc~ peak was barely detectable above the background, probably due to its low concentration (Table 1) and interference from the strong adjacent A1K~ peak. Therefore, within the limitations of the analytical technique, it was concluded that the Mg was likely to be distributed relatively uniformly with no significant concentration gradients. Lithium was incapable of being detected by the EDX method because of limitations imposed by the physics of the solid state detector, so attempts were made to search for the presence of Li concentration gradients by EELS techniques. The energy loss signal corresponding to the K-edge for Li ionization was not used because of strong interference from the A1 plasmon loss peaks [21]. Instead, attention was directed to the energy loss arising from plasmon excitation, where it is known that the quantized energy (Ae) lost is given approximately by

AE oCn 1/2

(3)

where n is the free electron density [21]. Thus, in the 8090 alloy, segregation of elements such as Cu, Mg or Li would be expected to change the position of the A1 plasmon peaks by decreasing Ae, as observed in binary alloys of A1-Cu [22], A1-Li [23] and A1-Mg [24]. However, since the EDX analyses in Fig. 10(b) showed a Cu depletion at the boundary, and there appeared to be no significant Mg gradients, any decrease in AE in the boundary region was more likely to be associated with Li enrichment. The EELS spectrometer was calibrated with respect to the first plasmon peak (AE~) of pure A1 (99.9995%) at 15.3 eV and repeated measurements on pure A1 after calibration gave inherent system errors of _+ 0.3 eV in the position of AE~. Hence only changes in position of the peak outside these errors could be attributed to the effect of solutes. In practice, repeated measurements at discrete points along lines traversing the boundary between GB precipitates in both tempers of the 8090 alloy gave a range of AE~ values between 14 and 15 eV, indicating a real decrease of ~ 0-1 eV in the position of the plasmon peak due to solutes. The reported decrease in AE~ in A1-Li binary alloys is 4CLi eV, where Cei is the atomic fraction of Li solute [21,23]. Therefore, assuming that the observed change in position of the first plasmon peak in the 8090 alloy was due primarily to Li, the results suggest that Li concentrations between 0-0.25 atomic fractions were present, which compare with an average Li atomic fraction of 0.086 based on the composition of the alloy (Table 1).

X. Chen, D. Tromans /Materials Science and Engineering A206 (1996) 163 175

171

Unfortunately, in terms of solute concentration gradients and segregation effects, the results were inconclusive because there was no systematic relationship between changes in AEj and sampling position relative to the GB in either alloy temper. Again, in similar manner to the STEM-EDX analyses, the minimum sampling volume was limited by the electron beam diameter (50-70 nm) and the presence of large numbers of smaller diameter (10 30 nm) cS'-phase precipitates (A13Li), with a Li atomic fraction of 0.25, undoubtedly influenced the EELS measurements and probably accounted for the larger (1 eV) decreases in Ae~.

4. Discussion Fig. 9. Bright field TEM micrograph of grain boundary precipitates in as-received material. Plane of view parallel to the %plane.

4. 1. General behaviour

The results showed that the near-threshold cyclic

crack propagation rates (da/dN) in the S - L orientation were higher than in the as-received T L orientation under all test conditions and the near-threshold crack velocities (da/dt) in the S - L orientation were either

18

16

--4t-

As-received

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Distance from the Grain Boundary (nm)

(a) 9

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As-received

7

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Re-aged

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Fig. 8. Dark field TEM mlcrographs showing fine intragranular dispersion of (~'-phase: (a) as-received alloy; (b) re-aged alloy. Plane of view parallel to the T-plane.

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Distance from the Gram Boundary (rim)

Fig. 10. STEM-EDX analyses of copper distribution across the grain boundary region of as-received and re-aged alloys: (a) traverse through grain boundary precipitate; (b) traverse between grain boundary precipitates.

172

X. Chen, D. Tromans / Materials Science and Engineering A206 (1996) 163-175

comparable to or significantly higher than those in the T - L orientation. The greater sensitivity of the S - L orientation to fatigue was probably related to the intergranular crack path, which contrasted with the predominantly transgranular crack path in the T - L orientation [1]. There was no beneficial effect of re-ageing on S - L fatigue behaviour in the low AK region, but the accompanying twofold increase in K~c was expected to allow S - L cracks to remain sub-critical to higher AK values before Kmax equalled Kic. Re-ageing produced lower average intergranular corrosion rates than the as-received alloy when exposed to 1 M A1C13. However, the maximum rates were fairly similar with a ratio of re-aged/as-received that was close to ~ 0.6. This similarity correlated with the similar appearance of the Cu concentration gradients in the GB region of both tempers (Fig. 10(b)).

4.2. Desiccated air The slightly deleterious effect of re-ageing on nearthreshold S - L fatigue behaviour in desiccated air appeared to correlate best with the accompanying decrease in yield stress, as indicated by the hardness measurements. Electrochemical phenomena were not expected in the absence of an aqueous medium and the mechanism of cracking was likely to be dominated by reversible slip processes in the crack tip region. Estimates of the reversed (cyclic) plastic zone (Ary) at the crack tip, based on the work of McClintock [25,26] and Lindley [27], may be represented by a general equation of the following form [1]:

A r y ~ p C ( AK~ 2 \ fly/

(4)

where C ~ 0.15, and p has a value of 1 for plane stress conditions and 1/3 for plane strain conditions. Therefore, if Ary controlled fatigue cracking in desiccated air, such that the reversed plastic zones were of equal size at Agth , then the following relationship should have prevailed: (AKth)An ,-, (0"y)AR (AKth)aA (ay)aA

(5)

where the subscripts AR and RA refer to as-received and re-aged tempers respectively. Using the S - L AK,h values for desiccated air (Table 4), Eq. (5) predicted a (O'y)AR/(O'y)RAratio of 1.6/1.3 (i.e. 1.2), whereas the ratio based on an estimated 9% decrease in Oy after re-ageing was ~ 1.1. These two ratios were sufficiently similar to corroborate the dominance of O-y-Controlled slip phenomena in desiccated air. Despite the fact that S - L fatigue cracking of the as-received material was intergranular, whereas T - L fatigue cracking was predominantly transgranular, Table 4 shows that the low AK S - L and T - L cracking

velocities (da/dt) were very similar in desiccated air. Thus, the intergranular S - L crack path in desiccated air may have been more a manifestation of the pronounced crystallographic texture commonly found in commercial AI-Li alloys and not necessarily indicative of weak S-plane grain boundaries. Typically, these alloys have been reported to exhibit a preferred {l10}(1T2> orientation, particularly at the mid-thickness plane, where the rolling plane (S-plane) was parallel to {110} and the rolling direction (L-direction) was parallel to ( 1 i 2 ) [28,29]. Studies by Hull [30] on sections of the same plate as that used in the present study showed (111) pole figures at the mid-plane that were consistent with the strong {110} (1 i 2 ) textures reported in the literature [29]. Therefore, assuming that slip occurred on {111 }<1i0> systems in the f.c.c, matrix of the 8090 alloy, the most favourably oriented slip systems in the presence of a tensile stress normal to the (110) plane (i.e. the S-plane) during S - L fatigue were (111)[011] and (lli)[101], both of which had Schmid factors of 0.408 [31]. Furthermore, assuming that fatigue cracking proceeded by restricted slip reversibility (RSR) on conjugate variants of {111 }, as proposed for austenitic stainless steel [31] and the 8090 alloy [1,32], then equal amounts of crack advance by RSR on the conjugate systems (111)[0il] and (1 li)[101] would have led to an overall macroscopic crack plane parallel to (110) (i.e. the S-plane) with a crack propagation direction parallel to (1i2> (i.e. the L-direction). Thus, in principle, RSR was able to produce a macroscopic crack plane parallel to the boundaries of the flattened grains in the S-plane and, if the RSR processes occurred close to these boundaries and the RSR increment on each conjugate slip plane was very small (i.e. a few Burgers vectors), cracking would have appeared to follow an intergranular path.

4.3. Aqueous environments The results indicated a strong correlation between copper-depleted regions at the GB and preferential intergranular corrosion. These observations, together with the known effects of Cu solute on retardation of corrosion activity of single-phase A1-Cu alloys in aqueous solutions [18] and the observed similarity between real-time intergranular corrosion rates and fatigue cracking velocities close to AKth in I M A1C13, all lead to the conclusion that S - L fatigue cracking in the aqueous environments was strongly influenced by corrosion dissolution phenomena that were directly related to the inhomogeneous distribution of Cu in the GB region. This was further supported by the observation that the gcorr during fatigue in 1 M AIC13 was - 0 . 7 8 Vsc z, close to the pitting potential of the as-received alloy [1]. It is known that Cu in solid solution raises the pitting potential of A1, so that Cu-depleted areas at the

X. Chen, D. Tromans / Materials Science and Engineering A206 (1996) 163-175

grain boundary will pit preferentially and lead to intergranular corrosion [18], as observed in this study and a previous TEM study on aged Al-4 wt.% Cu [331. Furthermore, under threshold conditions in 1 M AIC13, where AKth = 0.97 MPa m 1/2, Eq. (4) showed that the radius of the cyclic plastic z o n e , Ary, was ~ 200 nm, which was directly comparable to the width of the Cu-depleted zone at grain boundaries in both tempers, as seen in Fig. 10(b). Thus the combination of cyclic deformation, RSR processes and localized anodic dissolution could all have occurred within the dimensions of the Cu-depleted zone to accentuate intergranular cracking. The present study conveyed no definitive information on the heterogeneous distribution of active elements such as Na, K and/or Li at or near grain boundaries. Segregation effects involving one or more of these elements could have explained the low S - L toughness of the as-received alloy, as proposed by Lynch [2] and Webster [3]. For example, atomic absorption analyses of the 8090 alloy (Table 2) showed that there were sufficient alkaline impurities (Na plus K ~> 10 p.p.m.) to account for the low S - L toughness in the as-received temper, based on Webster's correlations. However, such segregation might have been expected to promote intergranular dissolution so that redistribution of the active segregant upon re-ageing [2] should have produced a difference in aqueous intergranular S - L fatigue behaviour. In fact, there was no difference between the two tempers, as confirmed by Fig. 3 and Table 4. There are three possible explanations for the similar behaviour: 1. The Cu-depleted grain boundary region dominated fatigue behaviour in the aqueous environments. 2. The improvement in toughness on re-ageing was unrelated to changes in the degree of segregation, but was caused by a change in deformation characteristics, as proposed by Blankenship and Starke

[4]. 3. The presence of segregant at grain boundaries was profoundly affected by deformation processes in the cyclic plastic zone, so that a quasi-steady state redistribution of segregant was established during fatigue, irrespective of the initial temper. With respect to the third explanation, it is widely recognized that the concentration of point defects, such as vacancies, increases during plastic deformation. Friedel [34] discussed several mechanisms by which deformation-induced vacancies may form, including the annihilation of dipole dislocations and movement of jogged dislocations. The analyses of Mecking and Estrin [35] showed that the relative excess vacancy concentration produced during deformation of AI increased with decreasing temperature, becoming several orders of magnitude higher than the equilibrium concentration at room temperature. Hence, repeated

173

deformation in the cyclic plastic zone at the tip of the fatigue crack in the 8090 alloy very likely produced and maintained an excess vacancy concentration that influenced the migration (diffusion) of solute atoms over small distances in the plastic zone. Thus, since S - L fatigue cracking was intergranular, the cyclic plastic zone around the boundary would have affected the segregation and redistribution of solutes in the boundary plane. In this manner, the crack tip would always have been propagating into a region whose degree of GB segregation was relatively independent of the temper condition.

4.4. Dissolution and hydrogen effects The low AKth in the A1CI3 solution correlated very well with the intergranular corrosion behaviour of the alloy and the influence of Cu concentration gradients on boundary dissolution. However, at AK values slightly above threshold (e.g. 2 MPa m l / 2 ) , the intergranular corrosion rates alone were insufficient to account for the rate of fatigue cracking, This situation was even more restrictive in distilled water, where intergranular corrosion rates were considerably smaller. Under these circumstances it was necessary to invoke a synergistic effect between anodic dissolution processes and cyclic loading or to seek an alternative explanation for the higher fatigue cracking rates in the aqueous environments relative to desiccated air. A suitable alternative was a hydrogen embrittlement process related to the generation of hydrogen by corrosion processes at the crack tip, as proposed previously to account for S - L splitting during T - L fatigue testing [1]. The very active Ecorr ( - 0 . 7 8 VSCE) observed during fatigue in 1 M AIC13 (pH 2.0) was equivalent to a cathodic overvoltage of - 0 . 4 2 V with respect to the reversible potential for hydrogen evolution [1], demonstrating that the cathodic reaction accompanying anodic dissolution at the crack tip was the evolution of hydrogen: 2H + + 2e

= H 2

(6)

A portion of the hydrogen would have been absorbed into the metal surface in monatomic form, which then diffused into the metal lattice under the influence of its activity gradient. If the surface concentration, Cs, of absorbed hydrogen was held reasonably constant by corrosion processes, diffusion theory [1,36] shows that the lattice concentration would have reached ~ 0.5Cs at a distance x below the surface (crack tip) after a time t that was determined by Eq. (7): X ~ ( D t ) 1/2

(7)

where D is the diffusion coefficient of monatomic hydrogen in the metal lattice.

174

X. Chen, D. Tromans / Materials Science and Engineering A206 (1996) 163-175

Using a value for D = 1 x 10-13 m 2 s 1, considered by Chen and Duquette [37] to be appropriate for A1-Li alloys, and t=3.333 × 10 -2 s, corresponding to the duration of one load cycle at 30 Hz, Eq. (7) shows that x = 5.77 x 10- 5 mm. This distance then determined the minimum distance over which hydrogen diffusion was significant during each load cycle and, therefore, corresponded to the minimum da/dN crack rate below which hydrogen embrittlement processes could have contributed to fatigue cracking. Inspection of Fig. 3 shows that the tests in aqueous environments reached cyclic crack rates of 1.6 × 10-4 mm/cycle, only twofold larger than the estimated hydrogen-influenced rate and supported the likelihood of a hydrogen embrittlement contribution to cracking. The influence of hydrogen on fatigue became even more likely in view of the fact that S - L cracking followed grain boundaries in the S-plane. These boundaries were subject to intergranular corrosion so that local enhanced hydrogen evolution occurred in the very region through which the crack propagated. Furthermore, grain boundary diffusion rates are normally expected to be higher than those of lattice diffusion [36], so that diffusion of hydrogen along the boundary ahead of the crack tip could have been larger than that estimated from Eq. (7), with a corresponding increase in the effect on da/dN rates. Also, the larger effect of 1 M A1C13 relative to distilled water on the enhancement of fatigue cracking rates was readily attributable to the different rates of corrosion in the two environments and the accompanying difference in the kinetics of hydrogen evolution, which controlled the amount of hydrogen (Cs) absorbed into the metal surface at the crack tip. Previous studies [1] indicated that S - L splitting during T - L fatigue tests in aqueous environments was related to hydrogen effects, consistent with the S - L cracking behaviour in the present work. The precise mechanism by which hydrogen contributed to intergranular cracking has not been resolved, although it is possibly related to electronic interactions between monatomic hydrogen and alloy atoms that result i n changes in bonding energy between the metal atoms. These effects could have arisen from the fact that quantum mechanical considerations lead to the conclusion that hydrogen in a metal changes the electronic structure of the host material in the local region [38]. This, in turn, should affect the electronic interactions (i.e. electron density) between the neighbouring host atoms and affect their local bond (cohesive) energy. The effect of changes in bond energy on fracture processes may be interpreted in terms of changes in the cohesive energy [39] or changes in surface energy [40] in analogous manner to steels. The tendency to form hydrides gives some indication of the possibility of electronic interactions between hydrogen and local metal atoms. It is known that A1

forms a hydride (A1H3), although it is not as stable as the Li hydride (LiH) [41]. Also, Na and K form halides which have the NaCl structure, similar to that of LiH [42]. If GB segregation of either Li, Na or K was present in the cyclic plastic zone, irrespective of the initial temper condition, interactions between monatomic hydrogen and these atom species may have played a contributory role during intergranular fatigue cracking in both the as-received and re-aged temper conditions.

5. Conclusions

Rising AK tests on S - L oriented 8090 plate in the as-received T8771 and re-aged conditions, plus comparison with previous T - L tests, allowed several conclusions to be made regarding cracking in the near threshold and low AK ( ~<3 MPa m 1/2) region of Stage I fatigue behaviour 1. Fatigue cracking of 8090-T8771 plate in the S - L orientation followed an intergranular path in desiccated air and aqueous environments. Crack rates (da/dN) and crack velocities (da/dt) in the aqueous environments were significantly higher in the S - L orientation than in the T - L orientation. However, crack velocities in the two orientations were similar when tested in desiccated air. 2. A 230 °C re-ageing treatment that doubled the S - L value of Klc had no beneficial effect on the Stage I S - L fatigue behaviour in aqueous environments, and had a slightly deleterious effect in desiccated air. 3. Intergranular S - L fatigue of the as-received and re-aged tempers in desiccated air was controlled by slip reversibility and the crack path was strongly influenced by the crystallographic texture of the alloy and grain shape anisotropy. 4. A Cu-depleted region was present at grain boundaries in the as-received and re-aged tempers that strongly influenced intergranular corrosion behaviour and intergranular S - L fatigue in the aqueous environments. Threshold behaviour correlated with the intergranular corrosion rates, whereas the enhanced cracking rates above the threshold were attributed to hydrogen embrittlement processes associated with the localized corrosion generation of hydrogen. 5. The study gave no definitive information on GB segregation of active elements, such as Li, Na or K, in the as-received and re-aged tempers. If differences were present prior to fatigue testing, they either had no effect on the resultant fatigue behaviour or were minimized by redistribution of segregant in the cyclic plastic zone.

X. Chen, D. Tromans / Materials Science and Engineering A206 (1996) 163-175

Acknowledgements T h e a u t h o r s w i s h to t h a n k t h e C a n a d i a n D e f e n c e R e s e a r c h E s t a b l i s h m e n t Pacific ( D R E P ) a n d the N a t u ral S c i e n c e s a n d E n g i n e e r i n g R e s e a r c h C o u n c i l o f C a n a d a for f i n a n c i a l s u p p o r t . A l s o , d i s c u s s i o n s w i t h D r J. M o r r i s o n o f D R E P are g r a t e f u l l y a c k n o w l e d g e d .

References [1] [2] [3] [4] [5]

[6] [7] [8] [9] [10] [11] [12]

[13]

[14]

[15] [16]

[17] [18]

X. Chen and D. Tromans, Mater. Sei. Eng., A189(1994) 45- 59. S.P. Lynch, Mater. Sci. Eng., ,4136 (1991)25-43. D. Webster, Adv. Mater. Proc., 146 (5) (1994) 18-24. C.P. Blankenship Jr and E.A. Starke Jr, Metall. Trans. A, 24A (1993) 833 841. H.E. Boyer and T.L. Gall (eds.), Metals Handbook, Desk Edition, American Society for Metals, Metals Park OH, 1985, pp. 4 81. M. Guttman, Mater. Sei. Eng., 42 (1980) 227 232. C. Lea, M.P. Seah and E.D. Hondros, Mater. Sci. Eng., 42 (1980) 233 244. S. Takayama, T. Ogura, S.-C. Fu and C.J. McMahon Jr, Metall. Trans. A, IlA (1980) 1513-1529. J. Kameda and C.J. McMahon Jr, Metall. Trans. A, 12A (1981) 31 37. P. Doig, D. Lonsdale and P.E.J. Flewitt, Metals Sci., 16 (1982) 335 344. M. Militzer and J. Wieting, Acta Metall., 37 (1989) 2585--2593. J. Morrison and Z.H. Zhai, in J. Morrison (ed.), Proc. Workshop Research and Development on Aluminum-Lithium Alloys, DREP Report 93-1, Defence Research Establishment Pacific, FMO, Victoria, B.C., 1993, pp. 288-304. A S T M E 399-78a, Plane Strain Fracture Toughness Testing q/ Metallic Materials, Annual Book of ASTM Standards, Part 10, American Society for Testing and Materials, Philadelphia PA, 1980, pp. 580 601. W.F. Brown and J.E. Srawley, Plane Strain Crack Toughness Testing ~!/ High Strength Metallic Materials, STP 410, American Society for Testing and Materials, Philadelphia PA, 1966, pp. 71 73. D. Singbeil and D. Tromans, J. Electrochem. Soc., 128 (1981) 2065 2070. A S T M E 647-787", Constant-Load-Amplitude Fatigue Crack Growth Rates Above 10 ~" m/cycle, Annual Book of ASTM Standards, Part 10, American Society for Testing and Materials, Philadelphia, PA, 1980, pp. 749 767. R.O. Ritchie, Mater. Sci. Eng., AI03 (1988) 15-28. M. Yasuda, F. Weinberg and D. Tromans, J. Electrochem. Soc., 137 (1990) 3716 3723.

175

[19] A.K. Mukhopadhay, D.S. Zhou and Q.B. Yang, Scripta Metall. Mater., 26 (1992) 237-242. [20] Powder Diffraction File, International Center for Diffraction Data, PA, 1991. [21] D.B. Williams, Norelco Reporter, 30 (I) (1983) 16 42. [22] P. Doig and J.W. Edington, Philos. Mag., 28 (1973) 961-970. [23] D,B. Williams and J.W. Edington, Philos. Mag., 30 (1974) 1147 1153. [24] G. Hibbert, J.W. Edington, D.B. Williams and P. Doig, Philos. Mag., 26 (1972) 1491-1494. [25] F.A. McClintock and G.R. Irwin, Fraeture Toughness Testing and its Applieations, ASTM STP 381, American Society for Testing and Materials, Philadelphia PA, 1964, 84-113. [26] F.A. McClintock, in O.F. Devereux, A.J. McEvily, and R.W. Staehle (eds.), Corrosion Fatigue, National Association of Corrosion Engineers, Houston TX, 1972, pp. 289 302. [27] T.C. Lindley, in L.H. Larsen (ed.), Subcritical Crack Growth due to Fatigue, Stress Corrosion and Creep, Elsevier Applied Science, London, 1981, pp. 167 213. [28] A.K. Vasudevan, W.G. Fricke Jr, R.C. Malcolm, R.J. Bucci, M.A. Przystupa and F. Barlat, Metall. Trans A, 19A (1988) 731 732. [29] P.S. Pao, L.A. Cooley, M.A. Imam and G.R. Yoder, Scipta Metall., 23 (1989) 1455 1460. [30] D.A. Hull, in J. Morrison (ed.), Proc. Workshop Research and Development on Aluminum-Lithium Alloys, DREP Report 93-1, Defence Research Establishment Pacific, FMO, Victoria, B.C., 1993, pp. 324-347. [31] C. Fong and D. Tromans, Metall. Trans. A, 19A (1988) 27652773. [32] X.J. Wu, W. Wallace, M.D. Raizenne and A.K. Koul, Metall. Mater. Trans. A, 25A (1994) 575-588. [33] A. Garner and D. Tromans, Corrosion, 35 (1979) 55-60. [34] J, Friedel, Disloeations, Addison-Wesley, Reading MA, 1964, pp. 116 123, 273-274. [35] H. Mecking and Y. Estrin, Seripta Metall., 14 (1980) 815-819. [36] P.G. Shewman, Diffusion in Solids, McGraw-HilL New York, 1963, pp. 13 14, 164 175. [37] G.S. Chert and D.J. Duquette, Metall. Trans. A, 23A (1992) 1551 1562. [38] R.P. Messmer and C.L. Briant, in R.A. Oriani, J.P. Hirth and M. Smialowski (eds.), Hydrogen Degradation oJ Ferrous Alloys, Noyes, Park Ridge, N J, 1985, pp. 140-166. [39] R.A. Oriani, Ber. Bunsen-Gesellscha[t Phys. Chem., 76 (1972) 848 857. [40] D. Tromans, Acta Metall. Mater., 42 (1994) 2043 2049. [41] D. Tromans, Scripta Metall. Mater., 27 (1992) 217-222. [42] R. Speiser, in R.W. Staehle, J. Hochmann, R.D. McCright and J.E. Slater (eds.), Stress Corrosion Cracking and Hydrogen Embrittlement of Iron Base Alloys, National Association of Corrosion Engineers, Houston TX, 1977, pp. 226-2413.