A critical evaluation of the stress-corrosion cracking mechanism in high-strength aluminum alloys

A Critical Evaluation of the Stress-Corrosion Cracking Mechanism in High-Strength Aluminum Alloys

SEONG-MIN LEE, SU-IL PYUN, and YOUNG-GAB CHUN

Attempts have been made to elucidate the mechanism of stress-corrosion cracking (SCC) in high-strength A1-Zn-Mg and A1-Li-Zr alloys exposed to aqueous environments by considering the temperature dependence of SCC susceptibility based upon the anodic dissolution and hydrogen embrittlement models. A quantitative correlation which involves the change of threshold stress intensity, Kzscc, with temperature on the basis of anodic dissolution has been developed with the aid of linear elastic fracture mechanics. From the derived correlation, it is concluded that the threshold stress intensity decreases as the test temperature increases. This suggestion is inconsistent with that predicted on the basis of hydrogen embrittlement. It is experimentally observed from the AI-Zn-Mg and AI-Li-Zr alloys that the threshold stress intensity, K~scc, decreases and the crack propagation rate, da/dt, over the stress intensity increases with increasing test temperature. From considering the change in SCC susceptibility with temperature, it is suggested that a gradual transition in the mechanism for the stress-corrosion crack propagation occurs from anodic dissolution in stage I, where the crack propagation rate increases sharply with stress intensity, to hydrogen embrittlement in stage II, where the crack propagation rate is independent of stress intensity.

I. I N T R O D U C T I O N

HIGH-strength aluminum alloys are well known to be susceptible to stress-corrosion cracking (SCC) in aqueous solutions. Anodic dissolution [1 7] could be supposed to be the mechanism from the fact that SCC is accompanied by preferential dissolution of grain boundary precipitates, which are more anodic than the matrix, and by the localized dissolution of the plastically deformed region within the precipitate free zone. ~,2~ Recently, many authors t8-18~ suggested that hydrogen embrittlement is in- volved in the SCC of the 7000-series aluminum alloys based on the following results: the change of SCC susceptibility with loading mode, t91 a similar potential dependence of SCC susceptibility to that of hydrogen permeability/~~ and a reduction in tensile ductility on pre-exposure to humid environments, tls,19~ Despite these observations, it is still not clear whether SCC is caused by anodic dissolution or by hydrogen embrittlement. This may be due to the amphoteric property t2~ of aluminum as well as difficulty in observing direct evidence of the SCC mechanism in aqueous solutions.

There is no doubt that SCC behavior is influenced in a complex manner by many variables, such as stress state (plane stress and plane strain), material parameters (chemical composition, strength, microstructure, etc.), and environmental factors (species, electrochemical potential, test temperature, etc.). Both anodic dissolution and hydrogen embrittlement are, in principle, distinct from the view of a fracture criterion. As a consequence, their SCC responses should be occasionally distinctive for specific variables. In addition, both mechanisms can si- multaneously operate in a material/environment system;

SEONG-MIN LEE and YOUNG-GAB CHUN, Graduate Students, and SU-IL PYUN, Professor, are with the Department of Materials Science and Engineering, Korea Advanced Institute of Science and Technology, Chongyangni, Seoul, Korea.

Manuscript submitted July 25, 1990.

however, it is supposed that either of them can predominantly operate in SCC under a certain combination of experimental variables. Considering these facts, it seems inappropriate to investigate the mechanism by using un-notched specimens for which the stress state is not clearly defined as a crack advances.

There are generally two distinguishable stages of stress- corrosion (SC) crack propagation for a precracked specimen. Stage I is more complex than stage II from the aspect of propagation kinetics, but the identification of the kinetics for stage II crack propagation provides an essential link for understanding the SCC mechanism. On the other hand, it is stage I crack propagation that de- termines the K~scc level. Thus, in this work, we will pay attention to the measurements of the threshold stress intensity due to SCC, K~scc, and the stage II crack propagation rate, (da/dt)n. The terms Ktscc and (da/dt)~i are regarded as measures of susceptibility to SC crack propagation in stage I and stage II, respectively. It should be noted that K~scc is an equilibrium parameter, whereas (da/dt)a is a kinetic parameter from the aspect of crack propagation.

For high-strength steels in aqueous solution or a gaseous hydrogen environment, t22-27~ it is well established that crack propagation is caused by hydrogen. In these systems, Ktscc generally increases as the test temperature increases, as illustrated in Figures 1 and 2, and (da/dt)~ increases up to moderate temperature; there is frequently a cut-off temperature above which the SC crack propagation rate is negligibly small or zero, as shown in Figure 3. In contrast, it was reported for A1-Zn-Mg alloys in aqueous solutions t7,8,111 that K~scc decreased whereas (da/dt)n increased as the test temperature increased. The contradictory results between the high- strength steels and aluminum alloys permit us to show the difference in mechanisms.

Therefore, in an effort to elucidate the SCC mechanism of high-strength AI-Zn-Mg and A1-Li-Zr alloys in

METALLURGICAL TRANSACTIONS A VOLUME 22A, OCTOBER 1991--2407

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T~

4

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Stress Intensity Factor-MPa. m vz Fig. 1--Hydrogen-assisted cracking propagation rate v s applied stress- intensity factor for the 200 ~ tempered AISI 4340 steel tested under a hydrogen pressure of 1.1 x 10 -1 MPa and at various temperatures: O - - O , 20 ~ @--@, 50 ~ rq--rq , 70 ~ I - - B , 90 ~ A - - A , 110 ~ A - - & , 120 oc.t271

aqueous solutions, their SCC susceptibility was experimentally determined as a function of test temperature. The relationship between Ktscc and temperature, T, is quantitatively derived on the basis of an anodic dissolution mechanism. The SCC mechanism is discussed by considering both the derived correlation between K~scc and T and the related experimental findings.

II. LITERATURE SURVEY

Of arguments on dissolution theory, the hypothesis that precipitate-free zone width is important to the resistance to SCC was early suggested and predicts that preferential deformation in the precipitate-free zone leads to preferential dissolution. This hypothesis is, however, open to speculation, as results among different authors are contradictory, t2,4,5] This may result from the difficulties in varying one parameter without a concomitant change in the other variables. A recent study by Pyun et al. u61 with careful treatment of microstructural variables showed that the nature of the matrix precipitates is of primary importance in SCC susceptibility. The significant role of the matrix precipitate has also been recognized on the basis of the dissolution mechanism. |6~ However, their observation u6) on the SCC susceptibility with matrix precipitate can be explained in terms of the relationship between hydrogen embrittlement susceptibility and slip planarity due to matrix precipitate obtained from cathodically hydrogen-charged aluminum a l loys . [28'29]

The SCC experiments of Green and Hayden I9~ con- ducted under different loading modes demonstrated that SCC in 7075 aluminum alloy is primarily caused by hy-

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50 60 70 80 I I

90

Stress Intensity Foctor--MPa. r~/~ Fig. 2--Hydrogen-assisted cracking propagation rate v s applied stress- intensity factor for the 450 ~ tempered AISI 4340 steel tested under a hydrogen pressure of 1.1 x 10 -1 MPa and at various temperatures: 0 - - 0 , 23 ~ 0 - - 0 , 34 ~ A - - A , 50 ~ A - - A , 70 ~

drogen embrittlement. They found that at a given nor- malized stress intensity (KJK1c and Km/Kmc), the 7075 aluminum alloy in a saline environment showed a shorter failure time in mode I than in mode III. In mode I, a triaxial tensile stress field exists ahead of a crack tip, and dissolved hydrogen is believed to collect in this triaxial region, which can lead to hydrogen embrittlement. On the other hand, such a triaxial stress field does not exist ahead of the crack tip of mode III, so there is no driving force for hydrogen collection and subsequent embrittlement.

A belief in the hydrogen mechanism has been raised by the Pyun group ~ with respect to hydrogen- recombination poison effects. Even though both the anodic dissolution rate and hydrogen reduction rate decreased for A1-Zn-Mg alloys in 3.5 wt pct NaC1 solution by the addition of hydrogen-recombination poison, the SC crack propagation period decreased with increasing hydrogen-recombination poison concentration. They uTl also discovered that the SC crack propagation time decreased when the specimens were both anodically and cathodically polarized. Such an observation has been in- terpreted as favoring the hydrogen mechanism, since Gest and Troiano u~ first proposed a hydrogen embrittlement mechanism based on the similar potential effects on SCC susceptibility and on hydrogen permeability.

Anodic dissolution theory has been based on experimental evidence to some extent, u-7j but a hydrogen embrittlement mechanism has been proposed, based on a variety of evidence, in recent works, t8-~8~ However, the examinations were partly made on un-notched tensile specimens. The tests with un-notched specimens would integrate over stages I and II of the propagation rate vs

2408--VOLUME 22A, OCTOBER 1991 METALLURGICAL TRANSACTIONS A

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.9 15

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I ( 5 6 i I i I ~ I ~ | i

2.0 2.4 2.8 3.2 3.6 4.0

Inverse Temperolure, 103/T Fig. 3 - Effect of temperatures on the stage II crack propagation rate for AISI 4340 steel tempered at various temperatures: 0 - - 0 , 200 ~ & - - & , 450 ~ and tested under a hydrogen pressure o f 1.1 x 10 -~ MPaJ27}

stress intensity curve as well as over the transition between the two stages, and thus, their results are presumably estimated as a whole. Considering that the SCC responses based on relevant mechanisms can be different for specific variables, such as applied stress intensity and test temperature, it is necessary to distinguish the mechanism in the stage I crack propagation from that in stage II.

III. Q U A N T I T A T I V E C O R R E L A T I O N S B E T W E E N Ktscc A N D T E M P E R A T U R E

A. Based on the Anodic Dissolution Theory

Crack propagation via anodic dissolution depends chiefly on metal dissolution at the crack tip. A reasonable correlation between the measured crack propagation rate and the dissolution current density was suggested, {3~ according to Faraday's law:

M da/dt = i a - [1]

zFd

where ia is the anodic current density at the crack tip, M is the atomic weight of the metal, z is the valency of the solvated species, F is the Faraday constant, and d is the density. If the crack tip remains bare, the maximum dissolution rate is given by the Tafel equation:

ib : io exp ( - olrFA ~I/RT) [2]

where i0 is the equilibrium current density, a r is the Tafel parameter describing the anodic dissolution behavior of the crack tip, A~7 is the potential difference between

electrolyte and metal, R is the gas constant, and T is the temperature. The presence of a passive film would result in a substantial decrease in the dissolution rate. No mat- ter which mechanism operates in SCC, it is reasonable to conclude that as a crack advances, bare surfaces are repeatedly formed at the crack tip by disruption of pas- sivating oxide film followed by the passivation on the s u r f a c e . [17] We assume that the instability of the passive film depends on straining at the crack tip. The incor- poration of this assumption into Eq. [2] yields the anodic current density, ia, a s follows:

i a = f ( K ) i b

= f ( K) i o exp ( - arFA n / R T ) [3]

where f ( K ) represents a function of stress intensity de- termining the extent to which the crack tip remains bare. Because our goal is to assess Klscc from the equilibrium aspect of crack propagation, it is reasonable to consider that the strain at the crack tip rather than the strain rate [3~ enters through f ( K ) .

With the aid of linear elastic fracture mechanics, the crack tip opening displacement, 6, in plane strain condition is generally given by

K 2 ( 1 - - l , 2 )

6 - [4] Etry~

where K is the applied stress intensity, v is Poisson's ratio, E is the elastic modulus, and Ors is the yield strength. The local strain in the vicinity of the crack tip, d e, can be represented in terms of distance directly ahead of the crack, x, and the crack tip opening displacement, 6 , as follows:[ 311

~p = S i x IS[ Combining Eqs. [4] and [5], one can obtain

K2(1 _ 1 2 ) d e [6]

E OrysX

Assuming that the extent of bare surface creation is lin- early proportional to the strain at the crack tip, one can regard the f ( K ) of Eq. [3] as the right-hand side of Eq. [6]. A constant strain should be maintained at the crack tip at K,h (that is, K~scc) necessary for the onset of crack propagation. In this case, the distance, x, is of the order of 2p, where /9 is the radius of the blunted crack tip. [32] Taking i~ in Eq. [3] as the threshold value, ith, at Kth, one gets the threshold stress intensity due to anodic dissolution, Kth(mD), as follows:

k \ i o / 1 - 1"2 exp

where /3 is the constant. Equation [7] states that the threshold stress intensity decreases as the temperature increases when crack propagation is caused by anodic dissolution.

B. Based on the Hydrogen Embrittlement Theory

The fracture by hydrogen embrittlement is mainly a stress-controlled process, t33~ and such susceptibility is

METALLURGICAL TRANSACTIONS A VOLUME 22A, OCTOBER 1991 --2409

enhanced when hydrogen enrichment is attained at a fracture zone and stress near the crack tip is intensified, based upon the concept of the critical hydrogen concentration and stress necessary for hydrogen embrittle- mentJ 33'34~ With increasing temperature, the maximum stress normal to the crack plane is reduced as a consequence of crack blunting due to enhanced plasticity. The equilibrium concentration of hydrogen within the stress field is given by/351

Ceq = Co exp ( ~riiVn/3RT) [8]

where Co is the initially uniform concentration, t r , /3 is the hydrostatic stress, and VH is the partial molar volume. Equation [8] indicates that the equilibrium hydrogen concentration at the trap site at which microcracks can initiate decreases as the test temperature increases. Hence, based on solubility consideration alone, the time required for the accumulation of a critical hydrogen concentration in the fracture zone is longer the higher the temperature. Eventually, a higher applied stress intensity is required for further crack extension at higher temperature, and thus, K,h is shifted toward higher values. This fact is experimentally demonstrated from some works on the hydrogen embrittlement tests for the high-strength steels exposed to gaseous hydrogenJ 23,241

Also, Gerberich and Chen [361 showed that the applied stress intensity equals the threshold, K,h, when the equilibrium concentration of hydrogen at the crack tip just achieves the critical concentration. With this concept, they obtained the threshold stress intensity due to hydrogen embrittlement, K,h(HE), under plane strain condition as follows:

RT K t h ( H E ) -- - - In (C~/Co) - Orys/20t [9]

aV.

where a is a material constant related to plastic con- straint (usually taken to be 1.25 cm -~ for high-strength steels), Cr is the critical concentration, and O-y~ is the yield strength. Since the yield strength generally decreases as the temperature increases and C~r is presumably not appreciably changed with temperature, it is readily seen from Eq. [9] that K,h(HE) increases with increasing temperature. This result is in contrast to that from Eq. [7] based upon the anodic dissolution model.

IV. E X P E R I M E N T A L P R O C E D U R E

A. Materials

The high-strength A1-Zn-Mg and A1-Li-Zr aluminum alloys used in this work were cast by ingot metallurgy in our laboratory. The SCC studies were made on double cantilever beam (DCB) specimens, as shown in Figure 4. Their chemical composition and thermal history are listed in Table I. Two different A1-Zn-Mg alloy specimens were prepared by varying quenching media. After the completion of the heat treatment, 5 mm on each side of the specimen were milled away to remove the residual stress caused by quenching, and then the final dimensions were obtained. The DCB specimens were loaded using a steel ball and bolt until a pop-in crack was mechanically introduced through the chevron notch

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steel ball~J"" ' ~ j~l

I" 127 detoil

./+ 25.,

(ram)

Fig. 4 - - G e o m e t r y of the double cantilever beam specimens used for stress-corrosion cracking tests.

starter. During SCC tests, the loading parts were insu- lated with epoxy to prevent galvanic interaction.

B. SCC Tests

For the SCC tests, the bolt-loaded DCB specimens were immersed in both 3.5 wt pct NaC1 solution and distilled water over a temperature range of 25 ~ to 90 ~ The SC crack propagation rate vs applied stress-intensity factor curves were obtained as a function of test temperature from the alloy systems. The stress-intensity factor, Kj, was calculated at the crack tip under mode I loading by using the following equation, t8~

AEh{3h(a + 0.6h) 2 + h3} 1/2 K,=

4{(a + 0.6h) 3 + hZa}

in which A is the diplacement of the DCB specimen at loading point, E is the elastic modulus, h is the half length of the specimen height, and a is the crack length. The crack length was measured from the crack tip to the loading line by an optical microscope with magnification 40 times.

V. E X P E R I M E N T A L R E S U L T S

A. Crack Propagation Rate vs Stress Intensity Curves

The crack propagation rate, da/dt, vs the stress-intensity factor, KI, relationships for the A1-Zn-Mg alloy specimen A are given in 3.5 wt pct NaC1 solution and distilled water in Figures 5 and 6, respectively. As the test temperature increases, the crack propagation rate vs stress-intensity curve is dramatically shifted upward, re- gardless of the environments. So, the threshold stress intensity decreased with increasing test temperature. This indicates that the onset of crack propagation requires a lower mechanical driving force as the test temperature increases. The stage II crack propagation rate increased with increasing test temperature.

The da/d t vs Kj relationships for the A1-Zn-Mg alloy specimen B are presented in 3.5 wt pct NaCI solution


Table I. The Specifications of the Aluminum Alloys Examined in this Work

Chemical Composition in Weight Percent

Zn Mg Li Cu Fe Si Mn Cr Zr A1

A1-Zn-Mg 4.0 2.8 - - 0.1 0.4 0.3 0.25 0.2 - - bal. AI-Li-Zr - - - - 2.0 2.0 . . . . 0.15 bal.

Deformation

Thermal History

Solution Treatment Quenching Medium Aging Condition

AI-Zn-Mg (Alloy A) AI-Zn-Mg (Alloy B) AI-Li-Zr

cold-rolled 1 h at 465 ~ oil 48 h at 120 ~ cold-rolled 1 h at 465 ~ water 92 h at 120 ~ forged 1 h at 520 ~ water 50 h at 200 ~

and distilled water in Figures 7 and 8, respectively. Only the stage II crack propagation behavior was observed in the specimen. It is also apparent that the stage II crack propagation rate increases as the test temperature increases up to 80 ~ irrespective of the environments.

Figure 9 also shows the effect of test temperature on crack propagation for the A1-Li-Zr alloy in 3.5 wt pct NaCI solution. The crack propagation is predominantly characterized by the stage I behavior. Similarly, it ap- pears that the threshold stress intensity decreases and the stage II crack propagation rate increases as the test temperature increases. It is noted that the crack propagation rates for the A1-Li-Zr alloy are lower by two orders of magnitude than those for the A1-Zn-Mg alloys. How- ever, the values are in agreement with the results obtained by Christodoulou et al. t371

B. Fractography

Figure 10 shows typical scanning electron microscopic fractographs of the A1-Zn-Mg alloy specimen A tested in 3.5 wt pct NaC1 solution at 35 ~ The analyses

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1,5 8

0

I ~i01

, i ~ ! i , i

l I J I , I I I a

5 I0 15 2 0 2 5 3 0

Stress Intensity Fact0r-MPa.~/z Fig. 5 - -S t re s s -co r ros ion crack propagation rate v s applied stress- intensity factor for the A1-Zn-Mg alloy specimen A tested in 3.5 wt pct NaC1 solution at various temperatures: O - - � 9 25 ~ 0 - - 0 , 35 ~ / k - - / k , 45 ~ & - - & , 55 ~

of the fracture morphology associated with crack propagation revealed that the SC crack path was exclusively intergranular and independent of the mechanical driving force over stage I (Figure 10(b)) and II (Figure 10(a)). The intergranular mode appeared in all of the AI-Zn-Mg alloy specimens, irrespective of solution treatment and test temperature. These results are in good agreement with other observations. [7.13,14] The forged A1-Li-Zr alloy specimens also showed the intergranular mode, regard- less of test temperature.

VI. D I S C U S S I O N

According to the foregoing theoretical analyses on the Kzscc, the threshold stress intensity dependence on temperature is quite different for the respective mechanisms. That is, based upon anodic dissolution, K~scc decreases and, on hydrogen embrittlement, increases with test temperature. So it is suggested that the temperature dependence can provide convincing evidence for the mechanism in the stage I crack propagation.

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17

Stress Intensity Factor- MPo ~/z 22

Fig. 7--Stress-corrosion crack propagation rate v s applied stress- intensity factor for the A1-Zn-Mg alloy specimen B tested in 3.5 wt pct NaC1 solution at various temperatures: ( 3 - - 0 , 27 ~ O - - O , 40 ~ A - - A , 60 ~ A - - k , 80 ~

In this study, it is experimentally shown from the present aluminum alloys that Klscc clearly decreases with increasing temperature. Thus, it is concluded that the anodic dissolution rather than the hydrogen embrittlement is chiefly operative in the stage I crack propagation for the high-strength A1-Zn-Mg as well as A1-Li-Zr alloys in aqueous solutions.

On the other hand, Speidel and Hyatt c8] and Speidel vii

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12

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A A A A A

0 0

A

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t

Stress Intensity

[ I

17

Factor- MPa. ~/z 22

Fig. 8--Stress-corrosion crack propagation rate v s applied stress- intensity factor for the A1-Zn-Mg alloy specimen B tested in distilled water at various temperatures: O - - O , 27 ~ O - - O , 40 ~ A - - A , 60 ~ & - - A , 80 ~

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I

-12 I 0 , I , t

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Stress Intensity Factor-MPa r4/z Fig. 9--Stress-corrosion crack propagation rate v s applied stress- intensity factor for the A1-Li-Zr alloy specimen tested in 3.5 wt pct NaCI solution and at various temperatures: O - - O , 30 ~ O - - O , 50 ~ A - - A , 70 ~ & - - & , 90 ~

observed that no effect of solution p H was found on the stage II crack propagation rate of precracked specimens from p H 11 to p H 6, whereas the solution p H vastly affected the stage I crack propagation rate. The acidity at the tip of a narrow SC crack remains around p H = 3.5 in aqueous solution, even if the bulk solution has a much higher p H value, t38] High hydrogen fugacity is maintained at the crack tip. Hence, it seems that the p H dependence of the stage I crack propagation rate cannot be satisfactorily explained by hydrogen mechanism; however, it is believed that the hydrogen embrittlement is responsible for the stage II crack propagation based on the mechanistic criterion for hydrogen embrittlement t33j (the combined effect of high hydrogen ion concentration and high stress intensity achieved in stage II). Such observations indicate that a single embrittling mechanism would not give a reasonable explanation for the different responses of crack propagation in stages I and II.

As a result of the SCC mechanisms, we suggest the following: a gradual transition of the SC crack propagation mechanism presumably occurs from anodic dissolution to hydrogen embrittlement as the applied stress intensity increases. The temperature dependence of Ktscc would indicate that the anodic dissolution model is responsible for the stage I crack propagation. In contrast, hydrogen embrittlement rather than anodic dissolution is responsible for the stage II crack propagation.

For the high-strength steels,[24.25.27] stage II crack propagation exhibits substantially different responses in two temperature regions, as shown in Figure 3. Based upon fractographic study, Gao et al. I251 have observed that the reduction in crack propagation rate with increasing temperature in the high-temperature region is accompanied by fracture mode transition from intergranular to micro- void coalescence. Thus, they have proposed a hydrogen


and temperature has been developed based upon the anodic dissolution model.

2. The temperature dependence of the threshold stress- intensity factor accounts for whether the stage I crack propagation is caused by anodic dissolution or by hydrogen embrittlement.

3. The mechanism for the SC crack propagation of high- strength aluminum alloys is presumably determined by applied stress intensity and, thus, gradually moves from anodic dissolution in stage I to hydrogen embrittlement in stage II.

4. From the invariance of the intergranular fracture mode and the monotonic change of crack propagation rate with temperature, it is suggested that the micromechanism of hydrogen embrittlement remains un- changed during the stage II crack propagation.

A C K N O W L E D G M E N T

The authors are indebted to the Korea Science and Engineering Foundation for financial support.

Fig. 10--Scanning electron microscopic fractographs for stress- corrosion cracking of the A1-Zn-Mg alloy specimen A tested in 3.5 wt pct NaC1 solution at 35 ~ (a) transition region between pop-in crack and stress-corrosion crack and (b) between stress-corrosion crack and mechanical crack. Arrows indicate the direction of stress-corrosion crack propagation.

partitioning model describing the different crack propagation responses.

However, Pyun and Lie t271 recently proposed that changes in the micromechanism of hydrogen embrittlement must be taken into account for the different responses in the stage II crack propagation rate. They concluded that the simultaneous changes in crack propagation rate and fracture mode with temperature could not be explained simply with hydrogen supply rate models ~22,251 but should include consideration of changes in bulk properties, such as yield strength and fractional surface coverage of hydrogen with temperature. Consid- ering the invariance of the intergranular mode and the monotonic change of crack propagation rate with temperature obtained for the A1-Zn-Mg alloys of this work, it is suggested that the micromechanisms remain un- changed during the stage II crack propagation.

VII. C O N C L U S I O N S

1. From the equilibrium aspect of SC crack propagation, a correlation between threshold stress intensity

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Documents

A critical evaluation of the stress-corrosion cracking mechanism in high-strength aluminum alloys