NATIONAL BUREAU OF STANDARDS REPORT

9821

INVESTIGATION OF THE DIRECTIONAL

EFFECTS IN THE STRESS CORROSION OF

ALUMINUM ALLOYS

Final Report

April 22, 1968

U.S. DEPARTMENT OF COMMERCE

NATIONAL BUREAU OF STANDARDS
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NATIONAL BUREAU OF STANDARDS REPORT

NBS PROJECT

3 120445 April 22 1968

Final Report

NBS REPORT

982 1

INVESTIGATION OF THE DIRECTIONAL EFFECTS

IN THE STRESS CORROSION OF ALUMINUM ALLOYS

by

J . K ruger - G . M . U giansky

J . R. Am brose

Contributors:

H . L . Logan - S. W. Stiefel

E. Escalante - A. W. Ruff

A. C. F raker - L . P . Skolmck

for National Aeronautics and Space Administration

George C. Marshall Space Flight Center

Huntsville. Alabama

Contract H-2151A

Control 1-6-54- 0 1046 -0 1 (IF)

IMPORTANT NOTICE NATIONAL BUREAU OF STA

for use within the Government, f and review. For this reason, the whole or in part, is not authori; Bureau of Standards, Washingtor

the Report has been specifically f

Approved for public release by the

Director of the National Institute of

Standards and Technology (NIST) on October 9, 2015

accounting documents intended

ubjected to additional evaluation

listing of this Report, either in Office of the Director, National

the Government agency for which

Dies for its own use.

U.S. DEPARTMENT OF COMMERCE

NATIONAL BUREAU OF STANDARDS

ABSTRACT

INVESTIGATION OF THE DIRECTIONAL EFFECTS IN THE

STRESS CORROSION OF ALUMINUM ALLOYS

Studies of the relative effects of various metallurgical and

electrochemical factors on the directional susceptibility to stress

corrosion cracking of a 7075-T651 aluminum alloy plate have been

conducted. The factors investigated that could have any possible

bearing on the directional susceptibility were as follows: 1) preferred

orientation; 2) grain morphology; 3) segregation; 4) precipitate and

dislocation configuration; and 5) electrochemical effects.

The role of grain morphology was found to be of paramount importance

in controlling crack propagation. It was found that a slow rate of

propagation along boundaries parallel to the applied stress, and the

fact that the intergranular crack must propagate primarily in this

direction accounts for the low susceptibility of longitudinally stressed

specimens.

This result suggests that the threshold stress for short transverse

specimens approximates the stress necessary to crack grain boundaries

normal to the applied stress, and the threshold stress for longitudinal

specimens approximates the stress needed to crack boundaries parallel to

the applied stress.

Preferred orientation was found to be of secondary importance when

compared to the effect of grain morphology. A high degree of preferred

orientation, rather than any specific slip plane orientations, was found

to increase the susceptibility of specimens with the same grain morphology'.

Electrochemical experiments showed that current flow was greater

in short transverse stressed specimens than in longitudinal ones showing

the creation of more active areas in the former. Polarization effects

due to stress were also minor when stressing longitudinally.

No directional effects related to segregation or precipitate or

dislocation configurations were found.

ACKNOWLEDGMENTS

We gratefully acknowledge the invaluable contributions of the

following: Mrs. M. Giles and Mr. D. Vieth of the National Bureau of

Standards for performing the electron probe microanalyses; Dr. H. Frankel

and Dr. T.D. Sciacca of the Goddard Space Flight Center of the National

Aeronautics and Space Administration for permission to use their X-ray

texture diffractometer; and the Japan Electron Optics Co., Ltd., for

making available their scanning type electron microscope.

This report was prepared by the Corrosion Section, National Bureau

of Standards under Contract No. H- 21 51 A "Investigation of the Directional

Effects in the Stress Corrosion of Aluminum Alloys" for the George C.

Marshall Space Flight Center of the National Aeronautics and Space

Administration. The work was administered under the technical direction

of the Propulsion and Vehicle Engineering Laboratory, Materials Division

of the George C. Marshall Space Flight Center with D.B. Franklin acting

as project manager.

iv
)
TABLE OF CONTENTS

Chapter Page

ABSTRACT ii

ACKNOWLEDGMENTS iv

I. INTRODUCTION 1

II. MATERIAL EVALUATION 4

III. EFFECT OF PREFERRED ORIENTATION ON SUSCEPTIBILITY 7

IV. EFFECT OF GRAIN MORPHOLOGY ON SUSCEPTIBILITY 17

V. EFFECT OF SEGREGATION ON SUSCEPTIBILITY 25

VI. EFFECT OF ELECTROCHEMICAL FACTORS ON SUSCEPTIBILITY... 28

VII. EFFECT OF PRECIPITATES AND DISLOCATION CONFIGURATIONS

ON SUSCEPTIBILITY 35

VIII. CONCLUSIONS 46

LITERATURE CITED 48

TABLES 52

FIGURES 59

V

Chapter I INTRODUCTION

Stress corrosion cracking may be defined as a complex interaction

between a sustained tensile stress and a corrosive attack resulting in

cracking and premature brittle failure of a normally ductile material.

The crack may follow either a transgranular (through the grain interiors)

or an intergranular (through the grain boundaries) path. In the aluminum

alloys, the crack path is characteristically intergranular.

During the fabrication of some wrought, high strength aluminum

alloy products, the grains become very elongated and do not significantly

recrystallize on further heat treatments (Figure 1). These products also

retain a high degree of preferred orientation (texture). The reported

(l

)

instances of stress corrosion cracking failures in aluminum alloys

have resulted from residual or assembly tensile stresses acting in the

short transverse direction relative to the grain structure (Figure 1).

Longitudinal tensile stresses, on the other hand, have rarely caused

stress corrosion problems. Hence, it is very important to understand

in more detail this great dependence of stress corrosion cracking

susceptibility on stress direction in the wrought, high strength

aluminum alloys

The purpose of the studies described in this report was to determine

the reasons for the great dependence in the stress corrosion cracking

susceptibility on stress direction in wrought, high strength aluminum

1
i t
alloys. Most of the studies concentrated on 7075-T6. However, some

studies were also carried out on 7075-T73, 2219-T37 and 2219-T87.

The strategy of attack was to explore all possible reasons and

factors that could influence the directional effects. This final

report describes in different chapters each of these possible sources

of directional differences. Many different and powerful tools of

modern materials science were required to do this. The possible

origins of di recti onali ty differences and the tools employed were as

follows:

1) The effect of preferred orientation— X-ray texture

diffractometer.

2) The effect of grain morphology-scanning electron microscopy

and optical microscopy.

3) The effect of segregation-electron probe microanalysis.

4) The effect of precipitates and dislocation configurations—

thin foil transmission electron microscopy (these studies were carried

out by Dr. A. Fraker and Dr. A.W. Ruff, Jr.).

5) Electrochemical factors—potentiostatic techniques.

On the basis of these many varied studies it was possible to

determine that the most significant factor influencing the directional

susceptibility to stress corrosion cracking was the effect of grain

morphology. The other factors play no role or, at most, a minor one

that is related to grain morphology.

2
r

•-

-

f- oq t >w

-
Publications and talks resulting from this work thus far are

the following:

Talks

1. Ugiansky, G.M., "Effect of Preferred Orientation on Stress

Corrosion Susceptibility" as panel member on Metallurgical Aspects,

Work Shop on Stress Corrosion of High Strength Aluminum Alloys, Naval

Air Engineering Center, Phi la., Pa., June 1967.

2. ligiansky, G.M., "Directional Effects on the Stress Corrosion

of a 7075-T651 Aluminum Alloy Plate", Engineering Materials Seminar,

University of Maryland, November 1967 and February 1968.

3. Ugiansky, G.M., Skolnick, L.P., and Stiefel , S.W., "Directional

Effects in the Stress Corrosion of a 7075-T651 Aluminum Alloy Plate",

AIME Annual Meeting, New York, New York, February 26, 1968.

Pub! i cations

1. Ugiansky, G.M., Skolnick, L.P., Kruger, J., and Stiefel, S.W.,

"The Rate-Controlling Step in Stress Corrosion Cracking", Scripta

Metallurgica, in press.

2. Ugiansky, G.M., "Directional Effects in the Stress Corrosion

Cracking of a 7075-T651 Aluminum Alloy Plate", M.S. Thesis at University

of Maryland, December 1967.

Work is now underway to prepare papers on the following:

1. Preferred orientation and grain morphology results.

2. Precipitate and dislocation studies..

3. The electrochemical studies.

3
.

'
Chapter II

MATERIAL EVALUATION

A. Introduction

In order to confirm that the material had the normal properties

for that alloy and heat treatment, both mechanical and stress corrosion

properties were checked.

B. Experimental Details

The material used for the research was a 2.5 inch thick plate of

7Q75-T651 aluminum alloy. The composition of the alloy was: A1 , 5.23

Zn, 2.13 Mg, 1.61 Cu, 0.184 Cr, 0.183 Fe, 0.040 Mn, 0.035 Ti , 0.035 Si.

The material was evaluated to confirm both the mechanical properties

and the stress corrosion susceptibility of specimens of three orientations (Figure 1) with respect to the rolling direction. For all material

evaluation tests, a specimen of dimensions shown in Figure 2 was machined

such that the reduced section was from the center of the plate. After

machining, the specimens were polished parallel to the axes with number

600 silicon carbide paper as a final preparation before testing.

Mechanical properties were determined for at least two specimens from

each direction using a Tinius-Olsen tensile test machine at a constant

strain rate of 0.005 in/min. Plotting of the engineering stress-strain

curves was done automatically utilizing an extensiometer attached to

the reduced section of the specimen. After testing, the 0.2% offset

yield strength was determined from the stress-strain curves and the

ultimate tensile strength from the maximum load.

4

In order to determine the stress corrosion susceptibility of the

material, specimens were dead loaded to 75% Y.S. (0.2% offset) to

eliminate any variation in susceptibility due to differences in

mechanical properties and exposed to an aqueous solution containing

0.3% NaCl , 3.0% ^C^Qy and 3.0% CrO^ in distilled H

20, solution pH = 0.9^ . Specimens were coated so that only the reduced section was

left exposed. The procedure followed was to pour the solution into the

glass cell surrounding the specimen just prior to loading. Failure

times were registered on a lapsed time indicator controlled by a microswitch placed beneath the load weights.

C. Results and Discussion

Table 1 shows that the maximum change in mechanical properties with

specimen orientation with respect to the rolling direction is only 10%,

whereas Table 2 shows a maximum difference of greater than two orders

of magnitude on the times to failure (tf

) of the stress corrosion specimens. This high dependence of stress corrosion susceptibility on specimen

(1) orientation is normal and as mentioned earlier, it was the purpose of

this research to explain why.

Incidentally, no failure occurred after 10,000 minutes for a longitudinally stressed specimen that had been room temperature aged for

*SP-5 Strippable Coating

Chemical Coatings & Engineering Co v, Inc.

Media, Pennsylvania

5
'
approximately one year before testing (shown in Table 2). The tests of

longitudinal specimens at various levels in the plate (discussed in the

next section) also showed a decrease in susceptibility after room temperature aging. These specimens did fail, however, probably due to the

higher testing temperature. Higher temperature overaging of the 7075

CD

alloy is known to produce immunity to cracking in 7075-T73 , but the

great effect of room temperature aging has not been, to the authors'

knowledge, discussed previously. It is suggested that any future

quantitative stress corrosion test of aluminum alloys be done so that

the precipitation in the alloy is held at a minimum after the specified

heat treatment is completed. This could easily be done by keeping

specimens refrigerated after aging and before subsequent tests are

carried out.

6

Chapter III

EFFECT OF PREFERRED ORIENTATION ON SUSCEPTIBILITY

A. Introduction

Over the years, many theories have been postulated to explain the

( 2- 11 ) stress corrosion cracking of aluminum alloys . The theories

differ widely from totally electrochemical to totally mechanical with

various combinations of the two in between.

( 5 ) One prominent theory is that promoted by Robertson and Tetelman

( F

)

and later by Parkins , in which it is stated that dislocation pile

ups at grain boundaries increase the stress level locally and, therefore, favor cracking at grain boundaries. Parkins has stated that

"In some cases these pile ups assist stress corrosion by producing

bursts of mechanical crack propagation, in others they may be the

cause, directly or indirectly, of intense electrochemical activity,

whilst in between there may be circumstances in which both mechanisms

help toward the total failure." Convincing evidence has recently been

presented by Holl^ and by Speidel^ which supports a dislocation pile

up mechanism for stress corrosion cracking of the aluminum alloys.

The localized stresses at grain boundaries due to dislocation pile

ups is directly related to the orientation of the slip planes with

respect to the grain boundary. Robertson, et al , have pointed out

(1 9 ) through an analysis made by Stroh that the stresses normal to the

grain boundary due to pile ups are maximized when the slip planes make

7
'
an angle of 70.5° with the plane of the grain boundary. This angle was

later substantiated by Smith and Barnby although they indicated a much lower dependence of stress on angle.

Because the normal stresses at the grain boundaries due to pile

ups are dependent or. slip plane orientation, susceptibility to stress

corrosion cracking is expected to be related to any preferred orientation of slip planes. (NOTE: Pole figures, as plotted in previous

reports, were found to be in error and were corrected for this report.)

B. Experimental Details

In order to determine the effect of preferred orientation on susceptibility, it was necessary to obtain specimens with different

textures, but with all other variables (grain morphology, composition

and mechanical properties) constant. In order to keep the grain morphology constant, the stress direction with respect to rolling (specimen

orientation) had to be held constant because of the highly elongated

grain structure. Preferred orientation was then determined as a function

of distance from the plate surface and other variables determined with

depth for constancy.

For the texture studies, disc shaped specimens (1.550 in. diameter x 0.125 in. thick) were machined with their flat surface parallel to the

rolling plane from various levels of the plate. The specimens were

polished on numbers 220, 320, 400 and 600 silicon carbide paper and

8
t
* number 600 microcut paper followed by an electropolish in a solution

of 20 ml. 70% perchloric acid and 100 ml. 95% ethanol (@ 0°C) to remove

the mechanically deformed surface layer. Each specimen was polished

o using five 10-sec. bursts (total 50 sec.) at 20 volts and 1 amp/cm .

For comparison purposes, a specimen having random crystal orientations was prepared by pouring a mixture of finely powdered aluminum

and parlodin into a mold to produce a disc shaped specimen of the same

dimensions given above.

The actual texture studies were made by the reflection method using

kk

a Siemens Texture Diffractometer . The complete X-ray installation

was comprised of the Kristalloflex 4 X-ray unit, the instrument cabinet

with radiation counting equipment and Kompensograph potentiometer

recorder, the texture diffractometer, and detector. When the entire

angle range was scanned (0 to 90°), the potentiometer recorded a curve,

which represented the reflection intensities in relation to the inclination angle a and aximuth \p. This curve was then evaluated in accordance

with stereographic projection.

Figure 3 shows the specimen used to determine stress corrosion

susceptibility as a function of specimen depth in the plate. The surface

preparation was identical to that of the specimens descr-i.h£d' above-.

*Buehler, Ltd. , Evanston, Illinois

**Siemans & Halske Aktiengesellschaft, Germany

9
' f

'
Masking off of all but 1/8 inch of the reduced section was carried out

using the coating mentioned earlier so that there was a maximum variation of only 1/8 inch in plate depth on any one specimen.

The method of stressing was identical to that mentioned previously.

Short transverse specimens were stressed to 50% Y.S. and longitudinal

specimens to 75% Y.S. (mechanical properties were determined as mentioned

earlier). As it was not necessary to compare failure times of short transverse specimens with those of longitudinal specimens, stress levels were

used that gave convenient failure times. However, the stresses were held

at a constant fraction of yield strength within the groups to eliminate

any differences in mechanical properties. The specimen cell arrangement

is shown in Figure 4 and the constant temperature system in Figure 5.

The complete system including the loading frame is shown in Figure 6.

A tubing pump was used to circulate 250 ml. of solution, which was

discarded after each test, through the cell at the rate of 400 ml/min.

and return it to the constant temperature system. The solution

(1*7 X8) temperature was held constant * at 35.0 ± 0.1 °C at the specimen

and was allowed to circulate around the specimen for GO seconds before

*Variable Speed Masterflex Tubing Pump

Cole Parmer Co.

Chicago, Illinois

10

loading. When failure occurred circulation of the solution was stepped

and automatically drained from the cell in order to prevent any further

corrosion. The lapsed time indicator also registered the time to

failure. A statistical analysis of variance was employed to determine

whether or not there was a significant change in susceptibility with

depth in the plate.

To determine changes in grain morphology with depth s specimens were

taken from various depths in the plate. Surfaces examined included

both the short transverse-long transverse and the short transverselongitudinal planes. The specimens were polished using standard

metal lographic techniques (through number 600 microcut paper and magomet

abraded polishing cloth) and etched with Keller's Etch to reveal the

grain structure. Grain boundary intercept distances (grain sizes)

were determined by several people independently to eliminate any

prejudices.

Emission spectrographic analysis of composition was made on material

taken from several levels within the plate.

C. Results and Discussion

As seen in Figures 7 through 11, the (111) pole figures changed

markedly with distance from the surface of the plate. The material

changed from nearly randomly oriented near the surface to the typical

*Buehler, Ltd., Evanston, Illinois

n

(19 ) aluminum rolling texture at the center of the plate.

Mechanical properties and stresses used for testing stress corrosion

cracking susceptibility at various levels in the plate are given in

Table 3.

The stress corrosion susceptibility of short transverse specimens

increases toward the plate center (Table 4) with an exception of 3/3"

below the surface. (This deviation from the trend will be discussed

later.) Table 5 shows that the susceptibility of longitudinal specimens

also increases at the plate center. The geometric mean of endurance

(20 21 ) time is reported as suggested by Booth, et al * . These data are

shown plotted in Figure 12. The statistical analysis of variance showed

in both cases that the above conclusions were significant.

Robertson, et al^

C/)

, have suggested that intergranular cracking will

occur through boundaries approximately normal to the stress at which slip

planes make an angle of approximately 70°. This angle was derived from

(12) the Stroh relationship , which shows that the stress, due to dislocation pile ups, normal to a plane making an angle e with the slip plane is

proportional to sin 0 cos 1/2 e. This function has its maximum at e = 70.5°.

It seemed logical that the preferred orientation could have oriented more

grains such that 9^70°. This, however, was not the case as determined

from the pole figures.

A careful analysis of the possible variations in slip plane orientations with respect to any particular crack plane showed that there will

12
I
always be at least one slip plane (111) making an angle between 54. /4°

and 90.0° to the crack plane. This leads to a possible variation of

only 10% in the function sin e cos 1/2 e. Furthermore, Smith and

(l 2 ) Barnby have shown that the stress normal to the plane, e degrees

from the slip plane, depends much less on e than indicated by Stroh.

They have shown that the tendency for cracking varies little when 0 <

e < 90°. Any change in susceptibility due to preferred orientation,

then, cannot be explained by a higher or lower fraction of grains being

oriented such that their slip planes make angles of approximately 70°

( 22 ) to the crack plane. Further evidence of this is shown by Logan .

He has seen that in ct-brass, which is susceptible to intergranular

cracking, both the angles between the (111) and the crack plane and the

angles between the (111) and the stress direction are completely random.

This shows that there is no relationship between slip plane orientation

and crack plane, and that the resolved shear stress on the slip planes

is also relatively unimportant. Therefore, any change in resolved shear

stress due to preferred orientation also could not account for the change

in susceptibility.

( 22 ) Robertson , however, has suggested that because a material having

a high degree of preferred orientation consists of lower energy grain

boundaries (lower angle) the corrosion rate and, therefore, the susceptibility to stress corrosion cracking would be less than for a material

13

consisting of randomly oriented crystals. Logan has also shown that

cracking, in fact, does occur primarily at grain boundaries where the

misorientation between slip planes is at least 13°.

Contrary to the above, the results of this research show that stress

corrosion cracking susceptibility increases with an increasing degree of

preferred orientation. It is believed that the following will clarify

this.

( 24 ) Davis and co-workers ~ have shown that for slip band continuity

across grain boundaries on aluminum bi crystals, the active slip planes

in the component crystals had to intersect the boundary on lines that

diverged by no more than approximately 15°. This angle correlates

approximately with the angle below which Logan observed no cracking

in a-brass.

It is believed, therefore, that material having a high degree of

preferred orientation, because it will contain more low angle grain

boundaries, will allow significantly more slip across grain boundaries

than material with random crystal orientations. Hence, the total

distance (L) between dislocation source and barrier (some high angle

boundary) will be greater in the oriented than in the random material.

Since (L) is a measure of the total number of dislocations piled up at

the boundary and the number, in turn, determines the magnitude of the

stress concentration, it is expected that material in which (L) is

greatest will be most susceptible to stress corrosion cracking. That

14
-

.

-
is, material with a high degree of preferred orientation would be

expected to be more susceptible than material with randomly oriented

grains.

In other words, preferred orientation makes some grain boundaries

approach being sub-grain boundaries or disappearing completely, thus

simulating material of larger grain size, greater (L), and greater

susceptibility. The above explanation of the dependence of susceptibility

on (L) has already been used to explain the greater susceptibility of

larger grained specimens 1 . (22 ) Logan's data can also be explained by the slip band continuity

across low angle boundaries. Since dislocation pile ups are believed

to be responsible for cracking, grain boundaries at which dislocations

do not pile up (low angle boundaries) would be immune to cracking.

Other variables (composition, mechanical properties, and grain size)

were also checked with depth in the plate to be sure that they did not

effect the stress corrosion susceptibility. Composition of the alloy

did not change significantly with depth. Table 6, and was ruled out as

a variable. Mechanical properties were also ruled out as a variable

because of the slight changes with depth seen in Table 3. In all tests,

however, any effects due to differences in mechanical properties were

minimized by testing at a constant fraction of yield strength. Grain

size and shape were constant throughout most of the plate interior, but

were not at 3/8" below the surface of the plate (Figure 13). This would

15
"

r
account for the relative high susceptibility of the short transverse

specimens taken from the 3/8“ level. Grain intercept distances in

the short transverse direction for different levels in the plate are

given in Table 7. This seemingly great dependence of susceptibility

on grain size led to the studies of the effects of grain morphology

on susceptibility explained in the next chapter.

16
.
Chapter IV

EFFECT OF GRAIN MORPHOLOGY ON SUSCEPTIBILITY

A. Introduction

Another factor thought to be important was the component of applied

stress normal to the grain boundaries which is dependent on orientation of

grains with respect to the stress. In a material having an elongated

grain structure, this is of particular importance.

Prestley^' , assuming that the susceptibility to stress corrosion

cracking is dependent upon the average resolved tensile stress acting

normal to the faces of an idealized grain, derived a mathematical model

to predict the variation in susceptibility to stress corrosion cracking

with the orientation of the test specimen in material having an elongated

(IS) grain structure. Haynie, et al , expanding Prestley's two-dimensional

model to a three-dimensional model based on an ideal elongated octahedral

grain, analyzed stress corrosion data for two aluminum alloys. They concluded that there is no significant difference in stress corrosion cracking

susceptibility of specimens stressed in the long transverse and short

transverse directions when the stress component normal to the grain

boundaries is the same. They also suggested that observed differences

as a function of stress direction are related to grain shape rather than

other metallurgical factors. However, as Haynie et al pointed out, changes

in other metallurgical factors were minimized by the use of thin plate

and sheet in their tests. Sprowls and Brown have pointed out that

17

Haynie's results should be regarded as tentative until other metallurgical

variables are adequately investigated.

Therefore, it was the purpose of the research in this section to

investigate the relative importance of the component of applied stress

normal to the grain boundary (grain morphology), which could affect the

susceptibility to stress corrosion cracking of a wrought, high strength

aluminum alloy.

B. Experimental Details

The specimen used for the effects of grain morphology studies (Figure

14) was designed specifically to fit into the specimen chamber of a scanning electron microscope. The specimens were extracted from the

center of the 7075-T651 plate such that the stressed face was in the

short transverse-longitudinal plane and the stress direction parallel

to the longitudinal direction. This particular orientation was chosen

because it is the grain morphology most immune to stress corrosion

cracking, and is ideal for observing crack initiation in an elongated

grain structure.

After a final light machine cut, the specimens were polished on

numbers 320, 400 and 600 silicon carbide paper, number 600 microcut

paper, on a magomet abraded polishing wheel and finally electropolished

for 30 seconds in the solution of perchloric acid and ethanol under the

conditions stated previously.

18

A small strain gage was then attached to the surface of the

specimen (Figure 15) to be stressed so that the relationship between

strain and deflection at the screw could be obtained. After obtaining

the strain-deflection curve, the stress-deflection curve (Figure 16) was

plotted with the help of previously obtained stress-strain curves.

Other specimens were then stressed to 75% Y.S. (0.2% offset) (56,500 psi)

and masked to leave only a 2mm square area exposed (Figure 17) to the

corrodent. These specimens were then corroded for various times from

30 seconds to 16 hours in the chromate solution. Non-corroded stressed

and corroded unstressed specimens were also examined for comparison.

Surface features were then observed in a scanning type electron microscope.

In some cases after cracking was observed, the stress on the specimen was

released (by removing the screw) so that cracks could be positively

identified.

Studies were also carried out with a high purity 7075 alloy forging

received from The Aluminum Company of America. Chemical analysis: 5.85

Zn, 2.57 Mg, 1.77 Cu, 0.25 Cr, 0.04 Ti , 0.01 Si, 0.00 Fe, 0.00 Mn.

After the miniature tuning fork specimens were machined from the

forging, they were heat treated to produce the -T6 condition (solution

treated to 900°F for 45 minutes in a salt bath, quenched in ice water,

*JSM Scanning Type Electron Microscope

Japan Electron Optics Lab., Ltd.

19

aged at 250°F for 26 hours).

These specimens were machined from the forging to produce longitudinal

and short transverse type specimens. After heat treatment, specimens

were given a minimum amount of mechanical polish, then electropol i shed as previously described.

Longitudinal and short transverse specimens were stressed at levels

ranging from 10 ksi to 70 ksi and corroded in the chromate solution for

times from 30 seconds to 10 minutes and observed with an optical microscope.

C. Results and Discussion

Observation of the surface condition of stress corroded specimens

utilizing the scanning electron microscope enabled study of the crack

initiation and propagation processes in the 7075-T651 alloy. Cracks

were seen to initiate in as short a time as 30 seconds in the longitudinally stressed specimens (Figure 18) although the same specimen was

not cracked while in the stressed but uncorroded condition. Successively

longer exposure times produced cracking as shown in Figures 19 through 24.

This cracking was found to occur primarily at the ends of the elongated

grain boundaries, normal to the applied stress as shown in Figure 20.

This Figure also points out the difficulty in studying early cracking at

low magnification. The same crack can clearly be seen in Figure 21 at the

higher magnification easily obtainable in the scanning electron microscope.

Cracking after 7 minutes exposure (Figures 22a and 23a) was confirmed by

releasing the stress on the specimens and observing closure of the crack

(Figure 22b and 23b). Stress corrosion cracks on the surface of specimens

exposed for 5 to 7 minutes were seen to have propagated to approximately

20

the same extent as cracks on a specimen exposed for 16 hours (Figure 24).

Cracking was also observed optically in 30 seconds on both short transverse and longitudinal specimens of the high purity 7075-T6 stressed at

60 ksi. (Figure 25 shows this for 1 min. exp.) Cracking observed in the

short transverse and longitudinal specimens was found to be always normal

to the applied stress at stresses as high as 70 ksi (Figure 26).

As found before, cracking proceeded along boundaries normal to the applied

stress and stopped when a boundary parallel to the stress was met (unfavorable

boundary), see Figure 27. At the lower levels of stress studied, cracking

was also seen to be present. On the longitudinal specimen corroded for 10

minutes at a stress of 20 ksi, cracking was observed (Figure 28). Cracking

was not as severe (pitting was more severe) as on specimens stressed to a

higher level, but it was significant because 20 ksi is well below the

threshold stress for failure of a longitudinal specimen.

Thus, cracking occurs at grain boundaries normal to the applied

stress in a short time (virtually instantaneously) regardless of the

type of specimen (i.e., short transverse and longitudinal). Cracking

occurs at these normal boundaries at low stress levels even on a longitudinal specimen. In the corrodent used, short transverse specimens

failed in approximately 5 minutes and longitudinal specimens in 24

hours. Since crack initiation is, however, very rapid for longitudinally stressed specimens, on those boundaries normal to the

stress direction, its long time to failure must be due to the fact that

crack propagation is difficult because most of the crack path lies along

boundaries parallel to the stress direction. Thus, an elongated grain

structure would be much more susceptible when stressed normal than when

21

stressed parallel to the principal axis of the grains. Figure 29, based

on the sort of crack initiation and propagation observed shows this

schematically. Actual micrographs of typical failures are shown in

Figure 30. The similarity between this figure and Figure 25 is

clearly seen.

In the short transverse direction, the propagation can always occur

in a plane essentially perpendicular to the stress direction. However,

when stress is in the longitudinal direction, initiation occurs on short grain boundary segments perpendicular to the stress direction,

but then the crack must somehow propagate along a direction parallel

to the stress before it can then connect another very short segment

that is perpendicular to the stress direction.

In summary, then, it is a combination of the facts that grain

boundaries parallel to the applied stress have a very low susceptibility

and that in the longitudinal specimens the crack must propagate almost

entirely along these paths of low susceptibility that accounts for the

relative low susceptibility of material stressed parallel to the longitudinal

direction. However, in short transverse specimens the crack may propagate

with very little deviation from boundaries perpendicular to the stress

(i.e., highly susceptible boundaries). Therefore, material stressed in

the short transverse direction is highly susceptible.

It is believed, then, that the threshold stresses for the different

grain morphologies 1 (different specimen orientations in a highly

elongated grain structure) can be related to the applied stress

22

necessary to crack grain boundaries at various angles to the stress

direction. If this is true, then the threshold stress for short transverse specimens approximates the stress necessary to crack grain

boundaries normal to the applied stress and the threshold stress for

longitudinal specimens approximates the stress needed to crack

boundaries parallel to that stress. Evidence for this is given by

the fact that boundaries normal to the applied stress crack at stress

levels as low as 20 ksi on either type of specimen.

The high susceptibility of short transverse specimens from 3/8"

below the plate surface (Figure 12) can be explained in light of the

grain size. Stress corrosion cracking is influenced by the amount that

the crack plane must deviate from a plane normal to the applied stress

in order to propagate through the specimen. Since the grains are smaller

along the stress axis at 3/8" below the surface (Table 7), the crack may

deviate less from a plane normal to the stress for specimens from the 3/8"

level than for specimens from deeper levels. Therefore, the material at

3/8" below the surface would be expected to be more susceptible than

material at lower levels in the plate.

Figures 31 through 33 show surfaces of unstressed specimens exposed

for 10 seconds and 5 minutes. General pitting has occurred to a much

greater extent on the unstressed specimens than on the stressed specimens

seen earlier (Figures 18 through 24). The stifling of general pitting

on specimens undergoing stress corrosion cracking may be due to the grain

23

areas which are cracking becoming more anodic than those areas normally

susceptible to pitting attack. The net effect would be inhibition of

general corrosion by an "internal cathodic protection mechanism." That

is, when the specimen is unstressed, small areas, e.g. weak sites in

the passive film, on the surface of the specimen are the anodes and

become pits. When stressed, the grain boundary becomes more anodic

than the interior of the grain (including the original weak areas in

the film). It thus acts as an anode that cathodically protects the rest

of the surface and stops pitting.

24
1
Chapter V

EFFECT OF SEGREGATION ON SUSCEPTIBILITY

A. Introduction

( 2 ) In 1940, Dix discussed the effect of segregation of alloy

elerrents on the stress corrosion of aluminum alloys. The larger

precipitates in the grain boundaries were through t to provide a susceptible path for the stress corrosion crack. Since then the

theory has been modified by many investigators as discussed by

Sprowls and Brown

It was the intention of this part of the investigation to determine

by electron probe microanalysis whether or not there was any variation

in segregation with orientation with respect to the rolling plane. That

is, this investigation was directed at determining the part that segregation plays in the directional susceptibility of the alloys tested.

B. Experimental Details

Flat, polished specimens were prepared from both the 7075-T651 and

-T73 material for electron probe microanalysis . Specimens were taken

perpendicular to the rolling plane from the center and from near the

surface of the plate. At each level, specimens were taken from both

the short transverse-long transverse and short transverse-longitudinal

planes.

Electron probe microanalysis was also made of the vicinity of a

25
.
crack on a flat specimen which had undergone stress corrosion cracking

in the previously mentioned chromate solution.

C. Results and Discussion

The target current images of areas with inclusions are shown in

Figure 34 for the specimens near the surface of both the 7075-T6 and -T73

conditions. The light areas indicate inclusions of lower atomic number

and the dark indicates inclusion of higher atomic number than the matrix.

In Figure 35, the magnesium X-ray images, the light areas are high in

magnesium content. These areas correspond to the light particles in

Figure 34 and, as determined by other X-ray images, are also high in

silicon. It was also determined that the dark inclusions of Figure 34

were low in magnesium and silicon.

Figure 36, the iron X-ray images, shows iron to be highly segregated

at the dark particles of Figure 34. Copper and manganese were also

segregated at these inclusions. Aluminum and zinc were found to be

depleted at all inclusions.

The chromium X-ray image, Figure 37, shows that there are zones

depleted of chromium that do not correspond to the inclusions of

Figure 34.

Titanium, not shown, also follows the chromium. These depleted zones

are seen to be more prominent in the 7Q75-T651 than in the -T73 material

and are believed to correspond to subgrain boundaries and may be somewhat

responsible for the higher susceptibility of 7075-T651

.

26
-
No major differences could be found between the two planes of

material examined. The only difference found in the electron probe

microanalysis between the material from the center and from near the

surface of the plate was that the inclusions and zones near the surface

were approximately twice as large as those near the surface. This is

believed to be due to the more highly deformed nature of the material

near the surface. No relation, whatsoever, could be found between any

segregation and the direction of rolling of the plate.

Electron probe microanalysis of the crack vicinity on the stress

corrosion specimen revealed a significant increase in chromium concentration in the vicinity of the crack, Figure 38. However, further

studies with specimens cracked in the chromium-free chloride ion

corrodent revealed no such effects. That is, no relationship could

be found between the crack path and any previous segregation of alloy

elements.

27
*
Chapter VI

EFFECT OF ELECTROCHEMICAL FACTORS ON SUSCEPTIBILITY

A. Introduction

The purposes of these studies were: 1) to learn something about

which factors in the environment control cracking, such as pH, solution

composition, and potential, so that some sort of cracking mechanism can

be arrived at and related to directional i ty ; and 2) to measure corrosion

rates under stress and without stress with respect to direction and to

learn if there is enhancement of corrosion rate depending on direction

and stress, as measured by current flow.

B. Experimental Results

Three specimen designs were used: 1) 0.75 x 0.75 x 0.030-inch square

specimens (Figure 39), capable of being stressed parallel to either axis

in a given plane; 2) tuning fork specimen (Figure 40) with a 0.375 x 0.500-inch face; 3) miniature tuning fork specimens of the same design

as those used in the scanning electron microscopy studies (Figure 14).

In each instance, specimens were machined from the center portion of the

plate in order to eliminate those effects resulting from the variation

in stress corrosion susceptibility with plate depth.

The procedure for surface preparation of the specimens prior to

exposure has been previously mentioned in this report.

Square specimens were clamped in a stainless steel bracket (Figure 39)

28

and all but an 1/8-inch square area in the center of the specimen

coated with electroplater's stop-off lacquer. Tuning fork specimens

were stressed before mounting in the support bracket and coated with

stop-off lacquer. Specimen assemblies were then placed in the appropriate electrochemical cell (Figures 39, 40b, and 41).

Two cell systems were used during the course of the investigation.

The initial design consisted of a pyrex glass body with a built-in

Luggin capillary and platinized platinum counter electrode. The cell

was later modified so that self-stressed tuning fork specimens could

be used. In addition, the Luggin capillary of the reference electrode

could be positioned to within 0.1mm of the specimen surface to minimize

IR drop in the electrolyte. The potential of the specimen surface was

maintained at a selected value by means of a potentiostat. Current

measurements were then made as a function of specimen potential. During

the course of exposure, the corrodent (500 ml) was circulated at a constant

flow rate of 600 ml per minute and directed at the specimen surface to

sweep away gas bubbles which would interfere with current measurements.

Electrochemical studies of the role of environment were made using

the three corrodents: 1) a 0.3% NaCI , 3.0% I^Cr^Oy, 3.0% CrO^ solution

whose pH was 0.9; 2) 3.5% sodium chloride, and 3) a 1.0N sodium nitrate

solution, both of whose pH was adjusted to 0.9 with hydrochloric and

nitric acid, respectively. These solutions were adjusted to these low

pH values because initial experiments showed that excessive film would

29
'
build up on tne specimen surface in the neutral solutions during anodic

polarization.

The effects of applied stress and chloride additions to a normally

non-aggressive environment upon longitudinal and short-transverse

stressed specimens were studied using the miniature tuning fork specimens.

A 1/8" diameter aluminum rod threaded at one end and simulating

the tightening screw was inserted into the specimen, thereby permitting

application of stress external to the electrochemical cell. All but the

polished surface of the specimen was coated with successive layers of

"Glyptal" lacquer and electroplaters stop-off lacquer.

Specimens were exposed for five minutes to environments of 500 ml

of 1 . ON sodium nitrate (0.02N in aluminum nitrate) which also contained

from zero to 100 ppm chloride ion. Solutions were recirculated at the

flow rate of 600 ml /minute at room temperature. Specimens were stressed

to 75% of yield either prior to exposure to the environment or within

the environment after electrically steady-state conditions were obtained.

In either case, optical micrographs were made before and after exposure,

and current measurements in polarized systems were made during the course

of exposure. In each environment, open circuit corrosion potential

measurements were made on both stressed and unstressed specimens.

Specimens were polarized to +0.100 volts vs SCE which correspond

to a potential within the active region of the polarization curve, where

30
*

'

i

i
film growth could be expected to occur.

C. Results and Discussion

Studies were carried out to measure the corrosion current associated

with crack propagation both as a function of time and stress direction.

A simple experiment for measuring the current and potential drop across

a cell composed of two identical tuning fork specimens immersed in the

chromate solution--one stressed to 75% of yield, the other unstressed--

has been performed.

Results from this measurement of open-circuit current flow in the

cell composed of a stressed and unstressed specimen were as follows:

for a longitudinal stressed specimen, the current density was 20 microamp/

2

cm and was relatively constant over the exposure period. Regular fluctuations were noted, but these could be due to hydrogen evolution which

would cause changes in the electrode surface area. Current measured in

the short transverse cell system rose steadily from an initial level of

2 2

20 microamp/ cm to 80 microamp/ cm over a 4 minute period until failure

occurred. Since the current is some measure of the amount of surface

available for reaction, this experiment shows that the short transverse

specimen, as expected from the microscopic studies, has more of its grain

boundaries favorably oriented for such reactions than the longitudinal.

Polarization curves were constructed for the 70 7

5 - T6 5 1 alloy in each

of the environments previously described. The potential was varied in

both an anodic as well as a cathodic direction, and current readings were

31
* }
made only after a steady-state value was reached. At least three sets

of data were obtained in each environment. The curves obtained are

shown in Figures 42 through 47.

The obtaining of these curves was the first step in our electrochemical studies of the role of directionality on the mechanism of

cracking. The open circuit potentials are labeled on these curves

for the chromate and for the chloride solutions and cracking does

occur at these potentials.

It is well known that specimens of the 7075-T651 alloy exhibit

shorter times to failure in the chromate solution, and much less severe

general pitting attack, than in the chloride environment. Furthermore,

on the basis of the polarization curves for this alloy in these solutions

(Figures 42 through 47) it is readily seen that 7075-T651 is considerably

more active in the 3.5% NaCl solution. From these observations, and the

previously noted stifling of pitting during stress corrosion, it is clear

that susceptibility to stress corrosion is not dependent upon general

corrosion susceptibility, but in all probability, depends upon stress

enhanced breakdown of a passive film (probably at the grain boundary)

followed by anodic dissolution at such grain boundaries.

In order to determine whether the role of the chromate under open

circuit conditions was to bring the system to a potential where cracking

occurs by a mechanism involving the formation of a passive film, potentiostatic studies in nitrate solutions containing no chloride were carried

32
I

'
out. By then setting the potential with the potentiostat at various

values, it was hoped that the potential for cracking could be determined.

We were unable, however, to obtain any evidence of stress corrosion

cracking in short transverse stressed specimens in nitrate solutions

under any conditions of applied potential and chloride concentrations up

to 100 ppm. However, addition of as little as 5 mg chloride ion (10 ppm

solution concentration) shifted the open circuit corrosion potential

nearly 400 mv in the active direction. This potential (-0.680 mv vs SCE)

corresponds to the corrosion potential in 3.5% NaCl where stress corrosion

cracking is known to occur. This observation would lead to the conclusion

that although small concentrations of chloride ion are needed to break

down the passive film generated on the metal in a nitrate environment,

such small concentrations are not sufficient to cause cracking. Further

work is necessary in the determination of these conditions of chloride

ion concentration and applied potential to cause cracking.

It was also noticed that the level of the corrosion current in

specimens, polarized to +0.100V, depended upon the mode of stress

application. Although the actual absolute current values were rather

unreproducible, the following trends were observed:

1. For specimens stressed prior to immersion, the corrosion

current was approximately 50% lower than the unstressed short transverse

specimens.

2. For specimens stressed, in situ, the current was perhaps 10%

33

greater than non-stressed specimens.

The explanation of this observation is unclear at this time, and

further experiments are necessary to elucidate this behavior.

Because it has not been possible to obtain reproducible results

in these electrochemical experiments, we were not able to tie them in

to optical and scanning electron observations of the surface to relative

electrochemical parameters to the difference in directional susceptibility,

as originally planned.

34

Chapter VII

EFFECT OF PRECIPITATES AND DISLOCATION CONFIGURATIONS

ON SUSCEPTIBILITY

A. Introduction

The purpose of this part of the study was to examine the dislocations, precipitate and grain boundary microstructures in various

commercial 7075 aluminum alloys and to see if these were related in

any way to susceptibility differences. Two tempers of this alloy were

studied, the T651 susceptible temper and the T73 resistant temper.

The technique applied was that of transmission electron microscopy of

thin metal sections taken from stress corrosion specimens of these alloys.

Both corroded and uncorroded specimens were examined in three different

sections, namely the longitudinal, long transverse, and short transverse

directions. It was hoped that these direct microstructural studies would

indicate possible explanation for differences observed in directional

stress corrosion susceptibility in this system, and at the same time

reveal any microstructural changes involved in the development of stress

corrosion resistance after suitable heat treatments.

Several recent studies of stress corrosion in alloys have concentrated

on the dislocation and precipitate micros true tures. It was shown by

(2 5 J Swann that alloys with cellular dislocation arrangements are resistant

to transgranular stress corrosion failure while alloys which exhibit planar

dislocation arrays are highly susceptible. Although aluminum fails inter35

showed that

( 8 ) granularly in stress corrosion, the studies of HoTT

susceptible underaged Al-Mg-Zn alloys do have planar dislocation bands

after plastic deformation. This type of dislocation arrangement is

missing if the alloys are given a further heat treatment which produces

immunity to stress corrosion cracking. [This observation also was made

/ 9 /? I by Ryum et al " who studied Drittleness vs. microstructure in Al-Mg-Zn

( 27 ) alloys.] Hornbogen has also emphasized the importance of dislocation glide bands, and has discussed the shearing of coherent precipitates

/ oo \ 29 ) with reference to the earlier work of Gleiter and Hornbogen . Speidel'

has elaborated on the idea that particle-shearing and dislocation pileup enhance stress corrosion cracking and he presents some supporting

experimental results. Earlier work by Thomas and Nutting^"' dealt with

zone formation, hardening, and with the importance of precipi tate-free

zones adjacent to grain boundaries in certain alloys. Much of the earlier

literature in this field is reviewed by Robertson and Tetelman 1 '5J. Jacobs concluded that stress corrosion resistance is lowered by an increased density of precipitates or dislocations with the precipitate

density being the most important factor. The stress corrosion studies

of Jacobs show electron micrographs of corroded surfaces. He has also

studied glide band structures in 7075 alloys and electron microscope

fractographs of cracked specimens.

B. Experimental Results

Samples were taken from the as-received commercially produced slabs

36

and from stress corrosion tuning fork shaped specimens which in some

cases had undergone stress corrosion testing at 75% of the yield stress.

No attempt has yet been made, however, to prepare foil specimens which

retain an original stress corroded surface. Wafers approximately 0.5mm

thick were cut from the stress corroded specimens by a spark discharge

slicer. These wafers were than spark planed to a uniform thickness of

0.3mm and electropol ished in a 1:2 nitric acid-methanol solution -25°C

and 8.5 volts to produce thin foils suitable for transmission electron

microscope studies. Foils taken from these specimens were mounted in

a rotating-ti 1 ting goniometer holder in the electron microscope and

examined at lOOkv accelerating potential. Selected area diffraction

pattern identification of precipitates was accomplished using the alloy

matrix diffraction spots as a local diffraction standard.

C. Results and Discussion

Grain Shape Anisotropy : As reported earlier, after optical microscope

examination of different sections of these alloys, a strong grain shape

anisotropy was found in the thin foil sections. Figure 48 is an electron

micrograph of a sample taken in the long transverse section from the T651

treatment. In this low magnification micrograph the grain shape anisotropy

is shown, the grains being elongated parallel to longitudinal direction

j of the material. In other sections of these alloys similar observations

are made although the amount of anisotropy varies with the section and

location in the slab. As shown by the roughly equiaxed grain in this

37

figure some variation in grain shape is noted in all sections. No

attempt has been made to determine lattice orientations from grain to

grain. The procedures which would involve selected area electron

diffraction are straight-forward but tedious. Previous X-ray

diffraction studies have indicated the nature of the preferred orientation in these materials and presumably those results would be substantiated by single grain studies of this type. In comparing the

precipitate and dislocation structures observed in the various sections

(to be discussed later in greater detail) it is clear that the high

degree of grain shape anisotropy has not led to corresponding differences

in these aspects of the microstructure. For example, the precipitate

density and dislocation density in elongated grains do not appear

appreciably different from the values found in equiaxed grains in any

of the sections. The degree of precipitation along the grain boundaries

does not appear to vary significantly between short and long boundary

lengths at the ends and sides, respectively, of the elongated grains.

These details will be illustrated in the thin foil micrographs and discussion which follows.

Dislocation and Precipitate Structures : In examining thin foils

made from various sections of these alloys we have employed bright field

transmission electron microscope techniques by tilting the area of interest

in order to control the diffraction contrast. Two general conditions are

of interest. First, by tilting the area away from the Bragg diffraction

38

condition, no contrast is developed at dislocations within the area,

however, precipitate contrast remains enabling one to determine precipitate sizes, shapes and densities. Figure 49 taken from a longitudinal

section of T651 at a relatively low magnification illustrates predominantly "precipitate contrast". That is, in only relatively few

areas are the diffraction conditions suitable for reviewing dislocations. One of these areas is indicated in the photograph. This condition

is emphasized to avoid misunderstandings concerning the distribution of

dislocations in the samples. Areas which appear to contain no or few

dislocations may in fact be unsuitably oriented for the diffraction

process required. This situation is usually indicated by somewhat

reduced contrast within those areas. The second technique involves

tilting rather near the Bragg diffraction condition to reveal dislocation

lines and at the same time reveal precipitate particles. This condition

is generally employed in high magnification studies due to the considerable

variation in lattice orientation in these polycrystall ine specimens. The

pi ate- like habit of the intermediate size precipitates and the tendency

for orientation is evident in Figure 49. The fine distribution of small

precipitates cannot be resolved at this magnification. We have identified these precipitates and many of the larger ones to be MgZn

2

in agreement with the results of other workers. These precipitates are known to

be anodic with respect to the matrix. Precipitate growth along the grain

boundaries is obvious in many areas and these grain boundary precipitates

39
"

i i
tend to be somewhat larger than those within the grains. No striking

examples of precipitate-free zones adjacent to the boundaries are found.

A thin foil section taken of the long transverse plane is shown in

Figure 50. The grain shape anisotropy is shown together with precipitate

structures in several grains and the dislocation and precipitate

structures in others. At this somewhat higher magnification the fine

distribution of small precipitates can be seen in the background. A

particularly interesting feature associated with this area concerns the

boundary at A. Selected area diffraction patterns taken on opposite

sides of this boundary show a very small if any lattice rotation. This

feature is apparently a low angle subgrain boundary oriented nearly

orthogonal to the long axis of the longitudinal grain structure. The

inset to this photograph shows a higher magnification view of this

particular area indicating several precipitates growing immediately

adjacent to this boundary. Structures of this type could represent

structural weaknesses in a material under the influence of an applied

stress. Since it has been previously reported that boundaries normal

to the longitudinal stress direction are generally sites of crack

initiation, a feature of this kind would be particularly suspect.

Another area taken from a long transverse section of the T651 alloy

is shown in Figure 51 . Two separate photographs have been obtained in

different contrast by tilting the sample. In Figure 51a the grain region

40

to the lower right shows only precipitate contrast while in Figure

<:\ b

51b the dislocations in this same area are revealed along with the

precipitates. It is interesting to note that the intermediate size

precipitates do serve as effective barriers for dislocations in this

material. The fine distribution of small precipitates is of course

the primary hardening constituent and its effect on the dislocation

lines is shown here as a general ragged appearance of the dislocation

images. Two triple grain boundary junctions are shown in this area

with advanced precipitation present at the boundaries. Several of

the grain boundaries are nearly orthogonal to the foil plane and

offer the opportunity to examine for precipitate-free zones adjacent

to them. One such boundary is shown at A in the figures and the small

precipitates appear to be present at a uniform density at all distances

from this boundary. In general, precipitate-free zones adjacent to

grain boundaries were not observed in any of the specimens examined

here, although occasionally such regions were observed. Further

examples of dislocation and precipitate structures in the long transverse

section of the T 651 alloy are shown in Figures 52 and 53. In Figure 52

at A, a dislocation is shown bowed around an intermediate size particle

indicating that the particle has disturbed the normal glide process of

the dislocation. A second example of this type of interaction is shown

at B and many other instances can be located in this field. The two grain

boundaries contained within this area can be seen to serve as effective

41

barriers for dislocation motion although the general dislocation

arrangement is randon and uniform. No tendency is observed for dislocation groups or planar arrays to form in the vicinity of these

boundaries. The large magnification of Figures 52 and 53 show clearly

the interactions between dislocations and the small precipitates in the

matrix. In Figure 53 two large dislocation loops are indicated at A.

These loops presumably arise as the dislocations attempt to move through

the lattice under the influence of stress and interact with the precipitate structures. Again there is lack of indication of glide band

formation in this region. Strong interactions between the dislocations

and grain boundaries are not observed.

A foil section taken in a short transverse plane is shown in

predominantly precipitate contrast in Figure 54. Several grains are

included in this region of the specimen and a considerable variation

in precipitate density is observed within the grains. Grain boundary

precipitation is indicated at several locations in this micrograph taken

from a specimen with the T73 treatment. The small hardening precipitates

appear to have grown considerably as a result of the thermal treatments

applied to this material in a comparison with the T651 treatment. The

dislocation density in this specimen and all others with the T73 treatment was appreciably smaller than that with the T651 treatment.

A higher magnification view of this specimen is shown in Figure 55,

a region illustrating a triple grain boundary junction. The small

42
ll-
precipitate distribution is clearly resolved in this case and for the

first time an orientation tendency is noticed among these precipitates.

Somewhat advanced precipitation at the grain boundaries is indicated.

As observed in the 1651 specimens, precipitate free zones adjacent to

the grain boundaries do not develop with the T73 treatment. This

particular region has been tilted away from the diffraction condition

so that dislocations contained within it are not in contrast. Another

example of a grain boundary region in this type of specimen is shown in

Figure 56. There are regions along this particular grain boundary where

the small precipitate density in the matrix is somewhat reduced and

other regions where there is no observable effect. Again the dislocation contrast has been suppressed in this figure.

The interactions between dislocations and the precipitates in

the T73 temper are illustrated in Figure 57 which is also taken from a

short transverse section. This micrograph is at a sufficiently high

magnification to clearly resolve individual precipitates in the background which have grown appreciably as a result of the thermal treatments applied to the alloy. There are many areas showing interactions

between the dislocations and the small precipitates and further interactions also with the intermediate size precipitates. The dislocation

distribution within the grains is reasonably random and no planar arrays

are observed. Not only is the dislocation density of the T73 treated

43

alloys much lower than that in the T651 treated alloys, but the dislocation lines appear straighter and less ragged possibly indicating

a reduced interaction with the precipitates. In view of the reduced

strength associated with alloys having undergone a T73 treatment these

observations are not surprising since they indicate that the precipitate

distribution is much less effective as a barrier to dislocation motion.

Conclusions from the microstructural studies of the commercial 7075

aluminum alloys used in this investigation are as follow. Both the T651

and the T73 treated specimens have high precipitate densities and electron

diffraction studies show that the majority of the intermediate sized and

small sized precipitates are MgZn^. The precipitates are randomly

dispersed throughout the grains and grain boundaries, and grain boundary

precipitation is only moderately advanced. The morphology of grains and

grain boundaries has not influenced the precipitate density or distribution

Sizeable precipitate-free zones adjacent to the boundaries are not a distinguishing feature in these materials are were rarely observed. The

principal differences in the T651 and T73 processed alloys are the dislocation density and arrangement and the size of the small precipitates.

The dislocation density is much greater after the T651 treatment than

after the T73. Dislocations are tangled and ragged in the T651 type

material while they are straighter after the T73 treatment. Dislocations

in the T73 specimen are few in number and are not easily observed. Growth

of the very small precipitates of the T651 specimen has occurred after the

44
• i^ad€

a »'!' 9

• i > : : ' £ aft* i' 3fiS ,.a ,y, f

; 't sqrDrt . -f j
T73 treatment. In the latter material, the formation of small directional platelets showing the growth relation of these precipitates to

the matrix is especially evident.

Thin foil sections taken normal to each of the three principal

directions showed no appreciable differences in dislocation density,

dislocation arrangements , or in precipitate density and distribution.

Therefore, the observed directional sensitivity in these alloys cannot

be explained in terms of these microstructural variables. The most

likely explanation still refers to the angle the grain boundaries make

with the applied stress direction as discussed earlier in this report.

45

Chapter VIII

CONCLUSIONS

A. The role of grain morphology was found to be of paramount importance

in controlling crack propagation. Observations of cracking in as short

a time as 30 seconds at grain boundaries normal to the applied stress on

the surface of longitudinally stressed specimens show conclusively that

it is the slow rate of crack propagation along boundaries parallel to the

applied stress that accounts for the low susceptibility of longitudinally

stressed specimens. These results suggest that the threshold stress for

short transverse specimens approximates the stress necessary to crack

grain boundaries normal to the applied stress, and the threshold stress

for longitudinal specimens approximates the stress needed to crack

boundaries parallel to the applied stress.

B. Preferred orientation was found to increase the susceptibility to

stress corrosion cracking in the 7075-T651 alloy by increasing the stress

at grain boundaries due to dislocation pile ups. However, its effect on

susceptibility is of secondary importance when compared to the effect of

grain morphology. A high degree of preferred orientation, rather than

any specific slip plane orientations, was found to increase the susceptibility of specimens with the same grain morphology.

C. The electrochemical experiments showed that current flow was greater

in short transverse stressed specimens than in longitudinal ones showing

4.6

the creation of more active areas in the former. Polarization effects

due to stress were also minor when stressing longitudinally.

D. No directional effects related to segregation or precipitate or dislocation configurations were found.

E. Pitting corrosion was found to occur to a much greater extent on unstressed than on stressed (undergoing stress corrosion cracking)

specimens. This phenomenon is explained by what is termed an "internal

cathodic protection mechanism."

47

LITERATURE CITED

(1) Sprowls, D.O. and Brown, R.H., "Resistance of Wrought High-Strength

Aluminum Alloys to Stress Corrosion", Technical Paper 17, Aluminum

Company of America, Alcoa Research Laboratories, New Kensington,

Pennsylvania (1962).

(2) Dix, E.H., Jr., "Acceleration of the Rate of Corrosion by High

Constant Stresses", Trans. AIME, 137 , 11 (1940).

( 3 ) Mears, R.B., Brown, R.H. and Dix, E.H., Jr., "A Generalized Theory

of Stress Corrosion of Alloys", Symposium on Stress Corrosion Cracking

of Metals, 329 (1944), ASTM-AIME.

( 4 ) Logan, H.L., "Film Rupture Mechanism of Stress Corrosion", J. Research

NBS, 48, 99 (1952), RP 2291.

( 5 ) Robertson, W.D. and Tetelman, A.S., "A Unified Structural Mechanism

for Intergranular and Transgranular Corrosion Cracking", Strengthening

Mechanisms in Solids, 217 (1962), ASM.

(6) Parkins, R.N., "Stress Corrosion Cracking", Metallurgical Reviews 9_,

35, 201 (1964).

( 7 ) Haynie, F.H. and Boyd, W.K., "Stress Corrosion Cracking of Aluminum

Alloys", DMIC Report 228 (1966), Defense Metals Information Center,

Battelle Memorial Institute, Columbus, Ohio.

(8) Hoi 1 , H.A. , "Deformation Substructure and Susceptibility to Intergranular Stress Corrosion Cracking in an Aluminum Alloy", Corrosion 23_,

6, 173 (1967).

48

( 9 ) Speidel , MoO., "Interaction of Dislocations with Precipitates in High

Strength Aluminum Alloys and Susceptibility to Stress Corrosion

Cracking", Paper presented at the International Conference on Fundamental Aspects of Stress Corrosion Cracking (1967), The Ohio State

Uni vers i ty.

(10) Sprowls, D.O. and Brown, R.H., "Stress Corrosion Mechanisms for

Aluminum Alloys", Paper presented at the International Conference on

Fundamental Aspects of Stress Corrosion Cracking (1967), The Ohio

State University.

(ID Jacobs, A.J., "The Mechanism of Stress Corrosion Cracking in 7075

Aluminum", Paper presented at the International Conference on Fundamental Aspects of Stress Corrosion Cracking (1967), The Ohio State

University.

(12) Stroh, A.N., "The Formation of Cracks as a Result of Plastic Flow",

Proc. Roy. Soc. 223 , 404 (1954).

(13) Smith, E.and Barnby, J.T., "Crack Nucleation in Crystalline Solids",

Metal Science J., 1_, 56 (1967).

(14) Prestley, J.S., Jr., "Effect of Specimen Orientation on Resistance to

Stress Corrosion Cracking of Aluminum Alloys--A Mathematical Model",

STP-425, Symposium of Stress Corrosion Testing (1967), ASTM.

(15) Haynie, F.H., Vaughan, D.A. , Phalen, D.I., Boyd, W.K., and Frost, P.D.,

"A Fundamental Investigation of the Nature of Stress Corrosion Cracking

in Aluminum Alloys", Technical Report AFML-TR-66-267 , Jan. 1967, Contract

49
jjisd

-

ii i
AF 33(6150-171), sponsored by the U.S. Air Force Materials Laboratory,

Nright-Patterson AFB, Ohio.

( 16) Sprowls, D.O., Lifka, B.W., King, W., Shumaker, M.B., Kelsey, R.A. , and

Vandenburgh, D.G., "Investigation of the Stress Corrosion Cracking of

High Strength Aluminum Alloys", Summary Report for the Period of

May 6, 1963 to July 6, 1965, Contract NAS 8-5340, sponsored by the

George C. Marshall Space Flight Center of the NASA, Huntsville, Alabama.

( 17) Romans, H.B. and Craig, H.L., Jr., "Environmental Factors Affecting the

Stress Corrosion Cracking Behavior of an Aluminum-Zinc Magnesium Alloy",

STP-425, Symposium on Stress Corrosion Testing (1967), ASTM.

( 18 ) Helfrich, W.J., "Influence of Stress and Temperature on Short-Transverse

Stress Corrosion Cracking of an Al-4.2 Zn-2.5 Mg Alloy", STP-425,

Symposium on Stress Corrosion Testing (1967), ASTM.

( 19 ) Barrett, C.S. and Massalski, T.B., "Structure of Metals", 556 (1966),

McGraw-Hill

.

( 20 ) Booth, F.F., Tucker, G.E.G., and Godard, H.P., "Statistical Distribution

of Stress Corrosion Endurance", Corrosion, 1_9, 11 , 390t (1963).

( 21) Booth, F.F. and Tucker, G.E.6., "Statistical Distribution of Endurance

in Electrochemical Stress Corrosion Tests", Corrosion, 2]_, 5, 173 (1965).

( 22 ) Logan, H.L., "Relationship Between Crystal Orientation and Stress Corrosion Cracking in Alpha and Beta Brasses", J. Research NBS, _56, 159 (1956),

RP 2662.

( 23 ) Robertson, W.D. , "Preferred Orientation as a Factor in Intergranular

Corrosion", J. Metals, 188, 790 (1950).

50
.

,
( 24 ) Davis, K.G., Teghtsoonian, E. and Lu, A., "Slip Band Continuity Across

Grain Boundaries in Aluminum", Acta Metal lurgica, T4, 1677 (1966).

( 25 ) Swann, P.R., "Dislocation Substructure vs Transgranular Stress Corrosion

Susceptibility of Single Phase Alloys", Corrosion 1_9, 102 (1963).

( 26) Ryum, N., Haegland, B., and Lindtveit, T. , "Brittleness and Microstructure

of Some Al-Mg-Zn Alloys", Z. Metallkde 58, 28 (1967).

( 27) Hornbogen, E., Discussion of reference ( 26).

( 28 ) Gleiter, H. and Hornbogen, E., "Theorie der Wechselwirkung von Versetzungen

mit Kohaerenten geordneten Zonen (I)", Phys. Stat. Sol. V2, 235 (1965).

( 29) Speidel, M.O., "Interaction of Dislocations with Coherent Particles and

Stress Corrosion Cracking of High-Strength Aluminum Alloys", Phys. Stat.

Sol. 22, K71 (1967).

( 30) Thomas, G., and Nutting, J., "The Aging Characteristics of Aluminum Alloys",

J. Inst. Metals 88, 81 (1959).

( 31 ) Jacobs, A.J., "The Role of Dislocations in the Stress Corrosion Cracking

of 7075 Aluminum Alloys", Trans. ASM 58, 581 (1965).

51

TABLE 1 MECHANICAL PROPERTIES AT THE PLATE CENTER (20°C)

Stress Direction Yield Strength (0.2% offset) Tensile Strength

(ksi

)

(ksi

)

Short Transverse 65.4 78.4

Long Transverse 70.7 83.3

Longi tudinal 76.7 87.7

52

r

TABLE 2

STRESS CORROSION SUSCEPTIBILITY AT THE PLATE CENTER (20°C)

Stress Direction Stress (ksi and % Y.S.

)

Average Endurance Time

(min.

)

Short Transverse 49.0 75% 4.8

Long Transverse 53.0 75% 143

Longitudinal 57.5 75% 834*

*No failure after 10,000 minutes for a specimen tested after room

temperature aging for 1 year.

53
f ,
TABLE 3

MECHANICAL PROPERTIES VERSUS DEPTH IN THE PLATE (20°C)

Stress

Di recti on Depth in Yield Tensile Stresses for Stress

Plate* Strength** Strength Corrosion Tests

(in.) (ksi j (ksi j (ksi ) [% Y.S.

Snort Transverse 0.375 68.5 78.3 35.5 52

0.625 65.2 75.8 35.7 55

0.875 66.4 77.4 36.5 55

1.125 65.6 77.0 36.5 56

1.375 65.2 74.2 36.2 5b

Longi tudi nal***

0.625 69.0 78.1 53.4 75

1.250 70.1 75.8 54.1 77

*Di stance between surface of plate and center of specimen gage length.

**0.2% offset yield strength.

***Room temperature aged for approximately 1 year before tests.

54

TABLE 4

SHORT TRANSVERSE STRESS CORROSION ENDURANCE TIMES

VERSUS DEPTH IN THE PLATE (35°C)

Depth in Plate (inches)

0.375 0.625 0.875 1.125 1.375

Endurance

3.2

4.1

4.2

7.5

8.4

12.0

Times 4.3 12.0

(minutes

)

4.9 12.4

6.5

9.4

19.2

5.7 5.4 4.0

6.3 5.7 4.6

7.0 6.0 4.9

8.2 6.1 5.0

10.6 7.9 5.8

20.0 11.3 b.8

Geometric Mean 4.9 11.4 8.7 6.8 5.0

(mi nutes)

Standard Devia- 1.5 3.8 3.7 1.7 0.6

tion of Log Normal

Distribution

(mi nutes

)

55

TABLE 5

LONGITUDINAL STRESS CORROSION ENDURANCE TIMES

VERSUS DEPTH IN THE PLATE (35°C)

Depth in Plate (inches)

0.625 1.250

539* 178*

187*

Endurance

Times

(mi nutes)

820

1551

1601

2412

548

692

919

972

2540 1074

3473 1935

Geometric Mean 1874 940

Standard Deviation

of Log Normal

Distribution

173 62

(mi nutes)

*Room temperature aged for approximately 4 months less than remainder

of longitudinal specimens.

56
-
TABLE 6

COMPOSITION VERSUS DEPTH IN PLATE

Element

Depth in

Plate (inches) Zn Mg Cu Cr Fe Mn Ti Si

Percent

0.375 5.17 2.14 1.62 0.182 0.188 0.040 0.033 0.035

±.03 ±.07 ±.02 ±.002 ±.005 ±.001 ±.001 ±.005

0.625 5.26 2.12 1.58 0.186 0.184 0.040 0.036 0.033

±.04 ±.03 ±.03 ±.002 ±.006 ±.001 ±.002 ±.006

1.250 5.26 2.13 1.63 0.184 0.178 0.040 0.036 0.036

±.03 ±.04 ±.02 ±.001 ±.008 ±.001 ±.0002 ±.007

Mean 5.23 2.13 1.61 0.184 0.183 0.040 0.035 0.035

±.02 ±.03 ±.01 ±.001 ±.004 ±.001 ±.0003 ±.004

Plus and minus values listed under the mean concentrations are the

standard deviations of means.

Each determination was run in quadruplicate.

57

TABLE 7

GRAIN SIZE VERSUS DEPTH IN PLATE

Depth in Plate (inches)

1.375

257

grains per inch in

the short transverse

direction*

Average number of

0.375 0.875

319 263

*Each number is the average of five different counts on each of

three specimens from that depth in the plate.

USCOMM-NBS-DC

58
.
Short

T r a n s verse

FIGURE 1

Composite photomicrograph of the grain structures in three

orientations at the center of a 2.5-inch thick plate of

7075-T651 aluminum alloy. Etched in Keller's etch. X 50

59

FIGURE 2

Specimen used for the material evaluation studies.

60

TOP OF

PLATE

A

FIGURE 3

Miniature specimen used to determine the mechanical properties

and susceptibility to stress corrosion cracking at various

depths in the plate of material.

bl

TENSILE

STRESS

FIGURE 4

Stress corrosion susceptibility testing cell.

b2

63

FIGURE

5

Recirculating

constant

temperature

corrosive

environment

system.

FIGURE 6

The complete stress corrosion susceptibility testing system.

Items shown include those in Figures 4 and 5 and the loading

frame.

64

RD

FIGURE 7

Absolute (111) pole figure for a plane of material at 0.26

inches below the surface of the plate. The constant intensity

contours are shown as multiples of the intensity of the (111)

reflection for a randomly oriented specimen.

65

FIGURE 8

Absolute (111) pole figure for a plane of material at

0.57 inches below the surface of the plate. The

constant intensity contours are shown as multiples of

the intensity of the (111) reflection for a randomly

oriented specimen.

66
,
FIGURE 9

Absolute (111) pole figure for a plane of material at

0.73 inches below the surface of the plate. The

constant intensity contours are shown as multiples of

the intensity of the (111) reflection for a randomly

oriented specimen.

67

RD

FIGURE 10

Absolute (111) pole figure for a plane of material at

0.89 inches below the surface of the plate. The

constant intensity contours are shown as multiples of

the intensity of the (111) reflection for a randomly

oriented specimen.

68

FIGURE 11

Absolute (111) pole figure for a plane of material at

1.20 inches below the surface of the- plate. The

constant intensity contours are shown as multiples of

the intensity of the (111) reflection for a randomly

oriented specimen.

69

ENDURANCE

TIME

(MINUTED)

3000

2000 —

1000 —

J

LONGITUDINAL

FIGURE 12

Stress corrosion endurance time versus depth of specimen

in the plate for both snort transverse and longitudinal

specimens. Extremes of data points are the standard

deviations based on the logarithms of endurance times.

70

FIGURE 13

Photomi crographs

plane (a, c & e)

plane (b,d & f)

(e & f)

Keller'

i nches

s etch.

of both the short transverse-longitudinal

and the short transverse-long transverse

at 3/8 (a & b), 5/8 (c & d), and 1 1/4

below the surface of the plate. Etched in

X 75

71
’
3-48 MACHINE SCREW

(WITH SHOULDER ROUNDED)

SECTION : A - A

n!nr> \V //

;« j'

/ 1 — '

r-r_vi

i 1 L

0.200

ALL DIMENSIONS IN INCHES

TOLERANCE i 0.001 THROUGHOUT

FIGURE 14

Miniature tuning fork stress corrosion specimen.

FIGURE 15

Miniature tuning fork specimen shown with strain gage in place

on stressed surface. X 5

73

2 4 6 8 10 12

DEFLECTION (in. X I0 3

)

FIGURE 16

Stress-deflection curve showing the stress on the test

area of the miniature tuning fork specimen versus the

deflection at the screw end of the specimen.

74

FIGURE 17

Miniature tuning fork specimen with the masking applied to

all but a 2mm square area on the stressed surface. X 4

75

ly

FIGURE 18

Scanning electron micrograph of a longitudinally stressed

specimen after 30 seconds exposure to the chromate solution. Note stress corrosion cracks. X 7,500

76

Stress

FIGURE 19

Scanning electron micrograph showing stress corrosion cracks

on the surface of a longitudinally stressed specimen after

2 minutes exposure to the chromate solution. X 15,000

77

Stress

FIGURE 20

Scanning electron micrograph of a longitudinally stressed

specimen showing stress corrosion cracks at the end of an elongated grain where the boundary is normal to the applied

stress (enclosed area) after 5 minutes exposure to the

chromate solution. Note the difficulty in observing the

cracks at this low magnification. The enclosed area is shown

at a higher magnification in Figure 21. X 750

78

Stress

FIGURE 21

Scanning electron micrograph of the area marked off in

Figure 20. Note that the cracking is primarily normal

to the applied stress. X 4,500

79

Stress

(b) lOy

FIGURE 22

Scanning electron micrographs of a longitudinally stressed

specimen after 7 minutes exposure to the chromate solution.

The specimen is shown in both the stressed (a) and unstressed (b) states. Note closure of the crack in (b).

X 4,500

80

(a)

FIGURE 23

Scanning electron micrographs of areas shown in Figure 22

in the stressed (a) and unstressed (b) states. Note closure

of crack in (b). X 15,000

81

Stress

lOy

FIGURE 24

Scanning electron micrograph of a longitudinally stressed

specimen after 16 hours exposure to the chromate solution

The general condition of the specimen was similar to that

of a specimen after 7 minutes exposure with little additional crack propagation. X 4,500

82

Stress

FIGURE 25

Stress corrosion cracks on short transverse (a)

and longitudinal (b) specimens after 1 minute

exposure to the chromate solution. Applied

stress = 60 ksi. (a) - X2,000; (b) - X3,000

83

FIGURE 26

Stress corrosion cracks on short transverse (a)

and longitudinal (b) specimens after 10 minutes

exposure to the chromate solution. Applied

stress = 70 ksi . X50

84

+-S

tress-*
; ’ ‘ . ' : ' / • r J >•. • : • :
FIGURE 27

A stress corrosion crack on a longitudinally

stressed specimen after 10 minutes exposure

to the chromate solution. Note that cracking

stops when unfavorable boundary is met.

Applied stress = 70 ksi. Keller's etch. X200

85

«-Stress+

FIGURE 28

Stress corrosion cracks on a longitudinal

specimen after 10 minutes exposure to the

chromate solution. Applied stress = 20 ksi.

X400

86

SHORT TRANSVERSE

CORROSIVE

ENVIRONMENT

FIGURE 29

Schematic representation of stress corrosion crack propagation through an elongated grain structure when stressed in

the s^rt transverse and in the longitudinal directions.

Note that in the short transverse specimen the crack may

oropaqate in a plane that deviates very little, while the

cra?k plane in the longitudinal specimen must deviate

greatly from a plane normal to the applied stress.

87
IS I

*
(b)

FIGURE 30

Photomicrographs of stress corrosion cracks

in (a) a short transverse specimen, X75, and

(b) a longitudinal specimen, XI 60. Note

similarity to Figure 29.

88

FIGURE 31

Scanning electron micrograph of the surface of an unstressed

specimen exposed to the chromate solution for 10 seconds.

Note the great amount of general pitting as compared to

specimens corroded under stress. X 1,500

89

* s

» 3 • - 4/

# .* |»-r

* .• < * ,' .'»> • f*9^m+ XT**

• . ' Pi * v *i5 •

• ^ *'*. *

’

• « v v »>•» .• #

^ :;. ,- n

^

< > «^.'

/ >* > ' • M

r:

1' > *7tr, i

SrT/.l* Y.K':>s.

*? *

I

(

*. •• •:

-. '•*v- -v;

*X) 'V.-v- • :•'". - r - ..y:- ,.

. 5

* «-*. v- y . ' 0?4. l€* .i V.\*7t „ W Jj*

lOy

FIGURE 32

Scanning electron micrograph of the surface of an unstressed

specimen after 5 minutes exposure to the chromate solution.

Note general pitting and intergranular corrosion. See Figure

33 for greater detail of enclosed area. X 1,500

90

ly

FIGURE 33

Scanning electron micrograph of the area marked off in

Figure 32. Note general pitting and intergranular corrosion. X 15,000

91

FIGURE 34

Target current images of areas with inclusions from a 7075 aluminum alloy

plate. The light areas indicate inclusions of lower atomic number and the

dark areas, inclusions of higher atomic number than the matrix. Both the

-T651 (a & b) and the -T73 (c & d) treatments are shown. Areas shown

include the short transverse-! ong transverse plane (a & c) and the shorttransverse-longitudinal plane (b & d). 7y/div.

92

(c) (d)

FIGURE 35

Magnesium X-ray images of areas shown in Figure 34. Light areas are high

in magnesium.

93

(a) (b)

(c) (d)

FIGURE 36

Iron X-ray images of areas shown in Figure 34. Light areas are high in iron.

94

(a) (b)

(c) (d)

FIGURE 37

Chromium X-ray images of areas shown in Figure 34. Light areas are high in

chromium.

95

(a) Target Current Plus CrKa

(c) iarget Current Plus CrKa

(b) Cr X-ray Scan

(d) Cr X-ray Scan

FIGURE 38

Electron probe microanalysis of areas along the path of

crack propagation. Figures (a) and (c) are target current

images with the CrKa X-ray rate meter trace superposed on

it. The dashed line is the locus of the electron beam as it was swept over the sample. Chromium X-ray area scans

are shown in Figures (b) and (d) in which light areas are

high in Chromium.

96

SCALE 1:1

FIGURE 39

PYREX GLASS

REFERENCE

ELECTRODE

— INPUT *

CLAMPING PLATE

TEFLON SEAL

0 RINGS

OUTPUT

COUNTER ELECTRODE

DRAiN

Initial electrochemical stress corrosion

cell design for externally applied stress

on square specimens.

97

FIGURE 40a

Tuning fork specimen in mounting bracket.

FIGURE 40b

Assembled electrochemical stress corrosion

cell for self-stressed tuning fork specimens.

98

LlI LiJ O Q 20

luq:

q: h

Ld o hh- Li_ Ld 3

Ijj LU _J Q_

drntmm or. lj ~ZL

CeL ZDOt—

i

L0_

99

Final

electrochemical

stress

corrosion

cell

design

for

self-stressed

tuning

fork

specimens.

CURRENT,

fJLQ/<Z

FIGURE 42

Cathodic polarization curve

3.0% K

? Cr2 07

- 3.0% Cr0

3

- SCE.

for 7075-T651 in 0.3% NaCl

Open circuit EMF = -440 mV vs

100

CURRENT,

yLLO/cm

FIGURE 43

Anodic polarization curve for 7075-T651 in 0.3% NaCl - 3.0% K

2 Cr2 07

- 3.0% Cr03>

101

CURRENT,

yLLO/cm

100,000 —

10,000 —

CvJ

1000 —

100 —

-MOO -1000 -900 -800 -700 -600 -500

POTENTIAL, mV vsSCE

FIGURE 44

Cathodic polarization curve for 7075-T651 in 3.5% NaCl

(pH = 0.9 with HC1 ) . Open circuit EMF = -670 mV vs SCE.

-400

102

800 -700 -600 -500 -400

POTENTIAL, mV vs SCE

FIGURE 4b

Anodic polarization curve for 7075-T651 in 3.5% NaCl

(pH =0.9 wi th HC1 )

CURRENT,

yCTCl/cnn

100,000

10,000

ftl

1000

100

-700 -600 -500 -400 -300

POTENTIAL, mV vs SCE

FIGURE 46

Cathodic polarization curve for 7075-T651 in 1 - ON NaNO^

(pH =0.9 with HN0

3

). Open circuit EMF = -260 mV vs SCE.

-200

104

CURRENT,

/JL

Q/C

100,000

FIGURE 47

Anodic polarization curve for 7075-T651 in 1.0N NaN0

3

(pH = 0.9 with HN0

3

)

10b

FIGURE 48

Grain morphology in a long transverse section of a stress corroded 7075-T651 specimen. M = 10,700.

106

FIGURE 49

Longitudinal section of a 7075-T651 specimen illustrating

precipitate distribution. M = 30,000.

107

FIGURE 50

A long transverse section of a 7075-T 651 specimen.

M = 34,000. The inset shows the area of a low

angle transverse boundary at M = 51,600.

108

FIGURE 51a

Grain boundary and precipitate structures in the

long transverse section of 7075-T651. M = 268,000.

109

FIGURE 51b

Area of 51a tilted to show dislocation contrast.

M = 268,000.

110

FIGURE 52

Dislocation and precipitate interactions in a long

transverse section of 7075-T651. M = 92,000.

Ill

FIGURE 53

An area adjacent to a grai nboundary in a long

transverse section of a 7075-T651. M = 155,000.

112

FIGURE 54

Grain and precipitate structures in a short transverse section of a 7075-T73 specimen. M = 30,000.

113

FIGURE 55

A triple grain junction region in the short transverse T73 specimen illustrating boundary precipitation

and general precipitate growth. M = 91,600.

114

FIGURE 56

Another boundary region in the T73 condition

specimen. M = 91 ,300.

115

FIGURE 57

Dislocation-precipitate interactions in a short

transverse section of the T73 specimen. M

= 268,000.

116