STRESS CORROSION CRACKING STUDIES

IN CARTRIDGE BRASS

by

Charles D. Easteal

A thesis submitted to the Faculty of Graduate Studies and Research in partial fulfilment of the requirements for the degree of Master of Engineering in Metallurgical Engineering.

Department of Metallurgical Engineering, McGill University,

August,l960. Montreal.

ACKNO'iJlLEDGEMENTS

The author wishes to express his

gratitude to Professer J.U. MacEwan for

the time and effort expended in directing

this research. Thanks are also due to

Dr. H.H. Yates for many helpful discussions,

to Mr. A.J. Ward for his help in preparing

materials and to the Defence Research Board

for financial assistance.

TABLE OF CONTENTS

1. INTRODUCTION

2. HISTORICAL REVIEW

2.1 General. 2.2 Corrosive Environments. 2.3 Al1oy Composition. 2.4 Paths of Cracking and Microstructure. 2. 5 Effects <iÛ Co1d-Work. 2.6 Thresho1d Stress.

3. THEORETICAL DISCUSSION

3.1 General. 3.2 Nature of Localized Corrosion. 3.3 Loca1ized Corrosion of Cartridge Brass. 3.4 Crack Propagation.

4. EXPERIMENTAL PROCEDURE

4.1 Introduction. 4.2 Materials. 4.3 Pre-corrosion Tests. 4.4 Loop Tests. 4.5 Metal1ography and Photomicrography

5. RESULTS AND DISCUSSION

Page

l

6

6 7

12 16 20 22

24

24 30 33 44

50

50 51 53 58 59

60

5.1 .Mechanica1 Properties of Materia1s. 60 5.2 Microstructure of Materials. 64 5.3 Loss of Strength Relative to Rolling Direction. 69 5.4 Measurement of Residual Stress. 73 5.5 Effect of Stress-Relief Annealing. 77 5.6 Paths of Cracking. 80 5.7 Relative Susceptibility of Strip Materia~ S5 5.8 Results of Loop Tests. 94

6. SUMMARY

Appendix I - The Hounsfield Tensometer. Appendix II- Residual Stress Measurement. References.

97

102 104 109

Fig.l

Fig.2

Fig.)

Fig.4

Fig.5

Fig.6

Fig.?

Fig.$

Fig.9

Fig.lO

Fie;.ll

Fig.l2

Fig.l3

LIST OF ILLUSTRATIONS

Page

Relevant part of Cu-Zn equilibrium 5 diagram.

Test specimen for pre-corrosion test. 56

Relationships between percentage reduc- 63 tion and ultimate tensile strength for cold-rolled cartridge brass strip.

Material A, longitudinal section, X200. 66

Material A, parallel to rolling plane, 66 X200.

Material B, longitudinal section, X200. 67

Material B, parallel to rolling plane, 67 X200.

Material C, longitudinal section, X200. 68

Material C, parallel to rolling plane, 68 X200.

Material A, exposed 48 hours, longitudi- 82 nal section, XlOO.

Material C, exposed 102! hours, longitudi- 82 nal section, XlOO.

Material A, exposed 21 hours, rolled 83 surface, XlOO.

Relationships between effective decrease 90 in cross-sectional area and time of exposure.

-1-

1. INTRODUCTION

Stress-corrosion cracking is currently defined

as the cracking resulting from the combined effect of

corrosion and stress {1). Cracks so produced are of a

brittle nature and may be intercrystalline or trans­

crystalline depending upon the alloy and corrosive atmos­

phere involved. Frequently the amount of corrosion associated

with the cracking is extremely small.

The types of alloys that may be made to stress­

corrosion crack are numerous. In fact it seems probable

that every alloy will so fail given the correct conditions.

It is fortunate that the corrosive atmospheres that cause

such cracking are relatively few. Austenitic stainless

steels in concentrated chloride solutions or steam contain­

ing chlorides; certain magnesium alloys in distilled water,

chromate or fluoride solutions; mild steel in nitrate

solutions; sorne aluminum alloys in chloride solutions;

copper-base alloys in moist ammoniacal atmospheres; are

sorne of the systems in which the phenomenon is observed {2 ).

In order that stress-corrosion cracking may occur,

surface or sub-surface tensile stresses must be present. It

is quite probable that shear and torsional stresses have an

effect but there is no record of compressive stresses having

-2-

caused this type of failure. In fact, the introduction of

compressive stresses by sorne such method as peening may be

used as a preventive measure.

The stresses causi ng cracking may be either

external or residual, the former resulting from applied

loads, and the latter remaining from fabrication and assembly

methods. This classification of the responsible stresses

gives rise to a further term that is commonly used. The

type of failure caused by the simultaneous action of corrosion

and residual stress upon sorne copper alloys is referred to

as nseason cracking". However, it is unlikely that the mech­

anism involved is any different from that associated with

external stresses.

-Of the two types of stress, that resulting from

applied loads is generally regarded as being less likely to

cause stress-corrosion cracking in practice. The reason

being that in any correctly designed structure the stresses

are known and may be kept to a safe working level. However,

residual stresses which may result from internal changes in

structure, cold-working, \veldi ng, shrink-fi ts or a~cidental

damage, may be of considerable magnitude and only calculable

by tedious methods. Even so, in all but the simplest cases

the local intensity and direction of these stresses may be

quite uncertain.

-3-

Most cold-working processes can generate sufficient

residual stress to cause stress-corrosion cracking. This is

particularly true of sinking, that is, drawing tubes without

a plug or internal mandrel, which may result in considerable

circumferential stresses. On the other hand the stresses

resulting from the cold-rolling of metal sheet or strip are

usually not large (J).

Susceptibility to cracking under the influence of

internal stress and corrosion was probably first recognised

in the alloys of copper and zinc ani copper and tin ( 4). The

importance and wide use of these alloys and the consequent

desirability of eliminating their tendency to crack led to

considerable investigation of the problem over periods up

to forty-five years. The brasses, in particular the cartridge

brasses containing approximately 70 percent of copper and

JO percent of zinc, have received much attention. This is

partly due to the fact that they are more prone to this type

of failure in certain environments than other copper alloys.

Also, they are well suited for drastic cold-forming operations

which may however, impart high residual stresses to the

finished shape.

Structurally, the cartridge brasses are incl uded

in the category known as the "alpha brassesn. The nalpha

brasses" can contain up to 37 percent zinc approximately

-4-

(see Fig.l) and consist entirely of the alpha-phase, which

is a primary solid solution of zinc in copper having a face-

centred cubic structure. Higher zinc contents give rise to

the appearance of the beta-phase, an intermetallic compound

of limited ductility which has a body-centred cubic structure.

The reduction of the stress-corrosion cracking

tendencies of internally stressed alloys may be accomplished

by decreasing the level of residual tensile stress. This

may be done by mechanical treatment s such as peening, roller

straightening or stretching, which cause sorne plastic deforma­

tion of the part or shape to occur. More conunonly, stress­

relief annealing is employed, which involves heating the

stressed mat erial to a lov" temperature for a short time.

The recrystallization temperature is not exceeded, so that

no decrease in hardness or strength occurs. In fact, in

sorne cases, an increase i n these properties is noted. The

time a nd temperature required depend on the severity of

deformation, alloy composition and acceptable stress level

after the treatment • For most cop:çe r alloys the time varies

between JO minutes and one hour, and the temperature from

150°C (J00°F) to J00°C (5750F).

de9C

-5-

4840 39 . .±.

1

1

•oc 1 +

1 /3'

1 1 1

WEIGHT PERCE'NTA6~ ZtNC

·--45

Fig. 1 Relevant part of Cu-Zn equilibrium diagram.

-6-

2. HISTORICAL REVIgw

2.1 General

Much experimental work has been performed in

connection with stress-corrosion cracking of copper alloys

as evidenced by the extensive literature related to the

subject. Sorne of the first work published is that of

Heyn (5). He refers to the cracking of cold worked copper

alloy shapes, in many cases years after the fabrication

had been completed. An example quoted is that of condenser

tubes, fabricated and stored, which had cracked without

having seen service. Heyn calls this "spontaneous cracking"

and attributes it to the residual stresses remaining from

the forming operation.

Moore, Beckinsale and Mallinson (6) examined a

number of brass parts that had cracked. They report that

certain factors were c ornmon to all the fai lures: a state

of initial stress existed; a degree of strain-hardening

due to cold-work was present; the surface showed signs of

chemical action which often appeared slight; and the crack

follov1ed an intercrystalline p~th. The authors state, hO\'J'ever,

that cold-work is probably not necessary for cracking but is

often the cause of the initial stress.

-7-

2.2 Corrosive Environments

The first substances recognised as having the

ability to cause certain cold-"l."lorked copper alloys to crack

were mercury and solutions of mercury salts. It is now

recognised that the failure so produced is not stress­

corrosion cracking. Nevertheless, immersion in one or

other of these substances has been commonly used to establish

whether or not the1avel of residual stress in a brass part

is sufficient to cause stress-corrosion cracking in service.

Moore and Beckinsale (7) say that the test was first used

in 1905 ~~d assign the credit for its development to

Dr. F. Rogers.

Heyn (5) states that corrodine; agents acting on

the surface of cold-worked copper alloys may cause cracking

provided that the attack is localized. He adds that sub­

stances causing general corrosion tend to relieve internal

stress by the removal of surface layers, and thus tend to

decrease the possibility of cracking. As far as brass is

concerned, he suggests that substances contained in air such

as carbonic acid with moisture, ammonia vapour and sulphurous

acid vapour may, after sufficient time, cause cracking.

Moore, Beckinsale and Mallinson (6) were the first

to attempt a comprehensive survey of the environments that

-8-

could cause cracking. They conclude that the cracking of

brass is never spontaneous. Such failure is always the

result of action by sorne substa~ce such as ammonia which

previously had often been regarded only as an accelerating

agent. In the course of their studies they subjected cups

spun from brass sheet to various corrosive atmospheres. They

report that of the substances tried only ammonia, ammonium

compounds and mercury are effective in causing cracking.

In a paper concerned with the incidence of firing

splits of small arms cartridge cases manufactured in India,

Grimston ( 8) reports th at the fai lures could be as cri bed

to stress-corrosion cracking caused by storing the cases

in wooden boxes which had been wetted with dilut e sulphuric­

acid pickling solution. In the discussion follo\ving the

presentation of the paper, Dickinson observes that he had

repro duced the cracking using drawn brass tube but that

similar tube stored i n glass i n the pre senc e or dllute

sulphuric acid had not cracked. Nevitt failed to reproduce

the cracking and states that for the cases to fail sorne con­

tamination from mercury or ammonia must have occurred. He

suggests that in the case of a loaded cartridge sorne con­

tamination by mercury from the detonator might take place.

Read , Reed and Ros enthal (9} have summarized sorne

of Johnston' s experiments on corrosive environments. ~v et

-9-

pyridine, wet trimethylamine and hydrogen sulphide did not

cause cracking, but sulphur dioxide, water vapour and air

did. Also, Johnston indicated that ammonia cannot cause

stress-corrosion cracking in brass in the absence of oxygen

and water vapour. This latter result has been confirmed in

part by Morris (10). He reports that a stressed specimen

partially immersed in a solution containing ammonia cracks

at the meniscus.

Rosenthal and Jamieson (11) report cracking of

stressed cartridge brass specimens stored with air, water

vapour and any one of a number of amines. Primary amines

appeared to be more active in promoting failure than secondary

or tertiary ones. HO\'!Tever, sorne doubt exists as to whether

it is the amines per se or ammonia formed by decomposition

of the amines which cause cracking.

Read, et al.(9) also mention work carried out by

Edmunds concerning the effects of hydrogen cyanide upon

stressed brass. Specimens were dipped in a one percent

aqueous solution of hydrogen cyanide and then suspended above

the surface. No decrease in t ensile strength was observed

after sorne 235 hours exposure.

Apart from the original hypothesis of Heyn (5),

namely, that carbonic acid might in the course of time assist

-10-

in the cracking of brass, little further investigation of

the effect of carbon dioxide appears to have been done

un til the work of Edmunds, Anderson and VJari ng ( 12). They

showed that the prese~ce of minute amounts of carbon dioxide

in moist ammoniacal atmospheres accelerated cracking. They

report that the effect is so marked that it may be that at

least a trace of carbon dioxide must be present for cracking

to occur. On the other hand, Read et al. (9) state

Johnston's view that large quantities of carbon dioxide in

an ammoniacal atmosphere prevent cracking altogether.

Sorne evidence of the nature of the substances

that cause the cracking of brass is given by chemical analysis

of the corrosion products. Moore, et al. (6) report the

presence of ru~onia, but not of nitrates, in the corrosion

products present at the fractured surfaces of bars which

had cracked whilst stored outdoors. Read, et al. (9) sub-

mit the results of a micro-chemical analysis reported by

Johnston. The analysis was performed on a white corrosion

product which had formed on brass cups exposed to an atmos­

phere containing ammonia, water vapour and oxygen. The

analysis indicated the presence of zinc only. No acid

radical being detected, it seems likely that the material

was zinc hydroxide or, alternatively, a complex of zinc

hydroxide and ammonia.

-11-

The production of a particularly heavy, dark blue,

crystalline corrosion product on brass that had been stressed

in an atmosphere containing amrnonia, air, moisture and carbon

dioxide is reported by Edmunds, et al. (12). An investigator

suspected the substance to be copper amino carbonate, Cu(NHJ) 4co3 •

This compound was synthesised and its colour and x-ray diffrac­

tion pattern were found to be identical to those of the corrosion

product. The authors state that using electron diffraction

techniques, basic zinc carbonate, 2ZnCOJ· 3Zn(OH) 2 has been

detected on other specimens.

-12-

2.3 Alloy Composition

Many investigators have established that the

tendency of brass to stress-corrosion crack inc reas es with

increasing zinc content. The question asto whether or not

pure copper will crack has never been definitely established

despite the many opinions expressed on the subject. Cook

(13) and Edmunds (14) summarize the existing evidence. On

the strength of this and their own experiments, they both

conclude that copper is immune to cracking in ammonia and

mercury within the limits of accuracy of all available tests.

In addition, Cook points out that the slight evidence of

failure of copper in service supports this conclusion.

Edmunds (14) quotes the opinion of Bassett to the

effect that copper-zinc alloys containing more than SO percent

copper had never been known to corrosion crack. Ho..,..;ever,

Moore, Beckinsale and Mallinson (6) found that whereas spun

cups of copper-zinc alloy containing 2.5 percent zinc withstood

the action of mercurous nitrate and ammonia, similar cups of

an alloy contaLüng ruxpercent zinc failed in both. Crampton

(15) used mercurous nitrate to test the susceptibility to

cracking of sunk tubes ma~ufactured from a series of copper­

zinc alloys containing from 6- 40 percent zinc. He states

that tubes containing less than 10 percent zinc were practically

-13-

immune to cracking, tubes containing between 10 and 20

percent zinc cracked on1y after prolonged exposure and that

tubes of higher zinc contents were very susceptible to crack­

ing.

In testing a number of a1loys by exposing them,

whilst under tension, to an atmosphere of ammonia vapeur

and air, Bulow (16) found the al1oy containing 35 - 36

percent zinc to be the most susceptible to cradking. He

states also, that where the applied stress was approximately

equal to the proportional limit, only alloys containing more

than 20 percent zinc showed a marked lack of resistance to

stress-corrosion attack. Edmunds (14) found that alloys

containing as little as three percent zinc wou1d crack in

ammoniacal atmospheres. Brasses containing 10 percent zinc

were slightly susceptible to mercury cracking whilst those

containing 20 percent zinc were high1y susceptible.

The relationship between zinc content and sensitivity

to cracking is well summarized by Cook (13). He states that

service failures due to stress-corrosion cracking are

practically limited to alloys containing 20 percent or more

of zinc, whilst alloys cont~ning less than 15 percent zinc

are almost immune. He adds, however, that whilst there is

much laboratory evidence to support this view, it has been

shown that certain circumstances can ca1se cracking in alloys

containing as lit tle as three percent zinc.

-14-

The effect of minor constituents upon the

cracking characteristics of brass has long been of interest.

In part this was because of the possibility that the crack­

ing tendency of brass might be caused or aggravated by sorne

common impurity which could be eliminated. On the other

hand it was felt that the addition of sorne element in small

quantities might increase resistance to cracking without

impairing the other properties of brass.

In reviewing the causes of condenser tube failure

in warships, Allen (17) states that most failures due to

stress-corrosion cracking occurred where the material of con­

struction was comparatively impure. Wilson, Edmunds, Anderson

and Peirce (18} quote the views of Fox and Jevons, both of

whom state that impurities have effects on season-cracking

resistance, although no supporting evidence is given. They

also mention the contrary opinion of Jonson who states that

chemical composition doe s not affect cracking except insofar

as the hardness of brass is affected.

Moore, et al. (6) endeavoured to r educe season­

cracking susceptibility of brass by small a dditions of man­

ganese. They report that cracks formed just as readi ly as

in pure brass when the al loy was submitted to the action of

mercury or ammonia. Crampton (15) notes that as far as the

cracking of brass tube was concerned, the presence of iron

-15-

and lead had no practica+ effect. Tin, on the other hand,

had a slight but distinct protective effect. This is in

contrast to the findings of Morris (10). He found that both

lead and tin lowered the resistance of 70:30 brass to stress­

corrosion failure.

The most comprehensive investigation concerning

the influence of added elements is tmt of Wilson, et al.

(18). Using 70:30 brass as a basis material, the effects

of adding thirty-six elements were studied. None of the

added elements was fourrl to increase the susceptibility of

brass, although silicon with a maximum effect at 1.5 percent,

increased the resistance of the alloy. Under certain circum­

stances, phosphorus, arsenic, barium, cerium, magnesium,

tellurium, tin, beryllium and manganese also impro ved resis­

tance to cracking . Another phase of this work involved the

preparation of extremely pure brass from high-purity zinc

and copper, special precautions being taken to prevent con­

tamination during melting . The cra cking behaviour of t he

high-purity brass was similar to that of commercial-purity

brass, indicating that the impurit ies normally present in

commercial brass do not a f fec t the susc eptibi lity to stress­

corrosion cracking of the alloy.

-16-

2.4 Paths of Cracking and Microstructure ·

In general, stress-corrosion cracks will propagate

in a plane that is perpendicular to the operative tensile

stress, and may follow either an intergranular or trans-

granular path. The au estion of the location of season-cracks ...

in brasses has been summarized by Bassett (19). He states

that in alpha brasses cracks are usually intergranular

although in sorne cases a f ew crystals on the line of break

may be fractured. In brasses where both the alpha and beta

phases are present cracks pass through the beta, whilst in

beta brasses the cracks are transgranular.

Graf (20) quetes the view of Althof to the effect

that with increasing plastic deforma ti on cracks in alpha brass

become progressively more transgranular. Edmunds (14)

investigated the susceptibility to cracking of a single

crystal of high purity alpha brass containing approximately

73.5 percent copper. The single crystal proved to have only

slightly greater resistance to cracking in an ammoniacal

atmosphere than polycrystalline brass of the same composition.

The indication being, therefore, that grain boundaries are

not absolutely necessary for stress-corrosion cracking of

brass containing only the alpha phase, to occur.

Both Edmunds (14) and Graf (20) note the experiment

using a single crystal of brass containing 64 percent copper, as

-17-

performed by Wassermann. When stressed in a moist ammoniacal

atmosphere the crystal cracked, the fractured surfaces being

perpendicular to the stress direction. Numerous other

cracks developed none of whic h followed any definit e crystal

plane.

Perryman {21) observes that the addition of small

amounts of aluminum to cartridge brass tends to change the

mode of stress-corrosion cracking from intergranular to

transgranular. Robertson and Bakish {22) quote the results

of Whitaker who confirms this effect for additions af

aluminum and also for additions of tin or silicon.

Crampton (23) conducted experiments on tubes

prepared from single crystals of cartridge brass by both

sinking and drawing. When tested in mercurous nitrate

solution the tube failed to crack in 24 hours. Tubes fab­

ricated by the same me th ods from polycrystalline brass of the

same composition failed within one minute. In a further

test involving a single crystal of high purity alpha brass,

Edmunds (14) found that the surface application of mercury

in no way affected the tensile strength and elongation.

However, the strength and duc ti lit y of polycrystalline

specimens were decreased consirl erably by this treatment.

This information indicates then a fundamental dif­

ference between the cracking of cartridge brass in ammonia

-18-

and tha. t in mercury or solution of mercury salt s. Cracking

in ammonia occurs preferentially at grain boundaries but

under certain circumstances can propagate in a transgranular

manner. Cracking in mercuzy is only an intergranular phenomenon.

In considering the effect of grain s ize upon crack­

ing tendencies, sorne of the first published views were thœ e

of Campbell (24). He inspected a large number of season

cracks in tubes and found all types of grain: large, small,

heavily worked and lightly worked, at the fracture surfaces.

He concludes that grain size and shape had lit tle or no

influence on the failures.

More recent work by Morris ( 10) sho\'lred th at increase

in grain siz e lO\'ITered resistance t o stress-corrosion cracking

in annealed brass. Croft (25) investigated the cracking ten­

dencies of brass wire, externally loaded and exposed to the

action of mercurous nitrate solution. He reports that whether

the wire was annealed, cold-drawn, cold stretched or stress­

relief annealed, increase in grain size increased the suscep­

tibility to cracking. The effect was so pronounced that he

suggests that stresses of a high magnitude are probably necessary

to cause season cracking in service of fine-grained metal.

In subjecting cartridge brass specimens of varying

grain size to an externally a pplied tensile stress of

-19-

10,000 p.s.i. in an ammoniacal atmosphere, Edmunds (14)

established a definite decrease in time to failure as the

grain size increased in the range 0.02 to 0.08 mm. Suscep­

tibility to cracking in the presence of mercury varied in

the same manner and was even more pronounced. Edmunds suggests

that the magnitude of the effect should be greater when the

cracking is due to the presence of mercury, in view of tŒ

completely intergranular nature of this ~pe of failure.

-20-

2.5 Effects of Cold-Work

Moore, Beckinsale and Mallinson (6) tested speci­

mens of annealed and cold-worked brass strip in a tensile

machine after exposure to an ammoniacal atmosphere, and

calculated the changes in ultimate tensile strength and

elongation. They report that increasing cold-work decreases

the deterioration of tensile properties for a given time of

exp os ure. Re ad, Reed and Ros en thal ( 9) point out hm'lever,

that this result is difficult to interpret. The reason being

that the three main effects of c old-work na mel y, work-hardening,

introduction of residual stress and change of grain shape,

are not separately evaluated.

In discussing a paper by Burns Read and Tour (26),

Bassett and Tour agree that the amount of cold-working to

which a specimen has been subjected is not a determining

factor in season cracking. Both suggest, however, that

residual strains caused by variations in working can be

decisive.

Croft ( 25) reports th at the susceptibility to stress­

corrosion cracking of brass wires depends to sorne extent upon

the method employed to finish them. He showed that wires

finished by cold-drawing were much less resistant than those

fi a ished by cold-stretching. He found also that the cracking

-21-

tendency of cold-drawn wires was at a maximum for a particular

degree of cold-reduction, 20 percent in the case of the material

studied.

Croft and Sachs (27) state that the susceptibility

to cracking of cartridge brass decreases as the degree of

cold work increases, providing residual stress is absent.

They suggest that this effect may be due to progressive frag­

mentation of the material giving rise to a smaller average

grain size. Cracking tendency is thus reduced, in the same

way that decrease in conventional grain size reduces it in

annealed material.

Edmunds (14) records a series of tests performed

on cold-rolled and recrystallized cartridge brasses. He

reports that with high applied stresses cold-rolled material

had superior resistance to cracking in ammoniacal atmospheres.

With low applied stresses the recrystallized material was

superior. It is suggested, however, that the apparently

lower resistance of the cold-rolled brass at low applied

stresses may be due to the presence of residual stresses at

the surface.

-22-

2.6 Threshold Stress

The existence or otherwise of a threshold stress,

that is, a stress level below which stress-corrosion cracking

cannot occur even after long periods of exposure, is debatable.

Moore, Beckinsale and Mallinson (6) subjected

specimens of annealed cartridge brass strip to the action

of mercurous nitrate solution and applied tensile stresses

of different values. They report that specimens failed within

minutes if the stress was greater than a value of the order

of 13,000 - 18,000 p.s. i. For stresses lor.-rer t han this

cracking did not occur within 14 days. The authors concede

th at somewhat lower stresses existi. ng in thin surface

layers may cause cracking.

These findings are confirmed by Crampton (15)

who states that the presence of mercury salts \"Till not cause

drawn brass tube to crack providing tœ residual stress at

exposed surfaces is less than 12,000 p.s.i. appruximately.

Similarly, Sachs and Espey (28) evaluate the threshold stress

for mercury cracking of brass tube as bei ng between 12,000

and 15,000 p.s.i.

Cook, in discussing a paper by Hudson (29), states

that a 11nealed and work-hardened brass specimens exposed to an

atmosphere of wet ammonia for th ree months were "complet ely

-23-

rotten and quite unsuitable for testing". If Cook's results

are valid then the concept of a threshold stress for stress-

corrosion cracking must be false.

Morris (10) partially immersed specimens of annealed

brass in arnmonia water, loaded them in tension, and measured

the breaking time for various values of applied stress. Stress

as ordinate was then plotted against breaking time as abscissa

on ordinary cartesian co-ordinate paper. At high stresses the

curve sloped downwards steeply. At low stresses the slope

of the curve decreased but, apparently, did not approach any

asymptotic value.

Edmunds (14) subjected unstressed, recrystallized

cartridge brass specimens to a moist ammoniacal atmosphere

and reports serious deterioration within a day. ün the basis

of his own work, and that of other workers Edmunds concludes

that if the corrosive conditions are such that a liauid . corrosion product forms on the specimen, then no threshold

stress exists.

-24-

3. THEORETICAL DISCUSSION

3.1 General

Any acceptable theory of stress-corrosion crack-

ing must explain why a normally ductile material may be

rendered brittle by the combined action of stress and

chemical environment. Furthermore, it should be capable of

explaining why alloys are susceptible only in specifie corrodents.

A suitabletheory should also explain the rate of cracking which

although low for a purely mechanical process, such as cleavage,

is high for simple chanical a ttack (30).

Early theories were quite inadequate as they em­

phasized either the chemical or the mechanical aspects of

failure and ignored their co-operative effect. A major

advance in the understanding of stress-corrosion cracking

was made by Dix (31). He postulat es that in order t hat stress

shall accelerate damage to an alloy in a corrosive environment

there must exist in the alloy a susceptibility to selective

corrosion along continuous paths , such as grain boundaries.

In addition, the stress must act in a di rect ion s uch th at it

tends to pull the alloy apart along these continuous paths.

Further work performed by Dix and his associates

led to the pre sentatio~ of their concept of a generalized

mechanism of stress-corrosion cracking (32). The following

-25-

quotation summarizes their interpretation of the cracking

proc ess:

"If attack penetrates preferentially along a narrow

path, it would appear axiomatic that a component of tensile

stress normal to the path would create a stress concentration

at the base of the localized corroded path. The deeper the

attack and the srnaller the radius at the base of the path,

the great er would be the stress co ne En trati on. Su ch a con­

dition would act to pull the metal apart along the se more or

less continuous localized paths. At sufficient concentration

of stress, the metal might start to tear apart by mechanical

action. Since it has been observed that a scratched metal

surface is anodic to an unscratched metal surface, the tear-

ing action described above would expose fresh metal, unprotected

by films, to the action of the corrosive environment. Because

this freshly exposed metal is more anodic, an increase in

current flm-'1 from the base of the localized path to the

unaffected surface would be expec ted and hence there would

be an acceleration of corrosion. Further corrosion would

result in further tearing of the metal and, as a result,

increased rate of penetration would occur b ecaœ e of the

mutual effect of the corrosion environment and the tensile

stress."

Subsequent work in t:œ field of stress-corrosion

cracking has not led to any serious change in the abov e

-26-

view of the process. However, sorne difference of opinion

exists as to the relative importance of corrosion and

stress in causing stress-corrosion failure.

One body of opinion, represented by the views

of Champion (JJ), holds that the main role of stress is

to rupture surface films. Stress-corrosion cracking is

thus regarded as being electrochemical rather than mechani­

cal. Film-rupture theories emphasize the rate of film for­

mation as compared with the rate of stress concentration.

It is suggested that if a dangerous stress concentration is

reached before film formation is completed, deformation occurs,

the existing film is danaged ani further cracking takes place.

On the other hand, completion of the film before the critical

stress concentration is reached will prevent rapid crack

propagation. From this viewpoint, suscepti bility to stress­

corrosion cracking is dependent upon the film-forming ten­

dencies and the formability of the alloy.

Hoar and Hines 04) regard the relat.i. ve importance

of corrosion and stress as be ing variable, dependi ng upon t œ system concerned. They believe that the cracking of austenitic

stainless steels is completely electrochemical, the function

of stress being to render the material at the tip or base of

the crack anodic and in a favourable state for preferential

dissolution.

-27-

Although there is no doubt that deformation can

rupture films and that stress can cause anodic behaviour at

the tip of a crack, it seems unlikely tha t corrosive attack

alone can cause the rapid failures due to stress-corrosion crack­

ing that are c ornmonly observed. Also, the britt le nature of

these failures seems to indicate that stress has sorne more

important mechanical role in causing them. In view of

the se circumstances it seems probable th at stress and cor­

rosion mutually initiate a deformation process that results

in brittle failure.

Keating ( 4} developed a the ory of the mechanism

of stress-corrosion cracking on the above basis. Later,

Gilbert and Hadden (3 5} prop osed a more detailed theory to

account for the stress-corrosion cracking of an alU111.inum

alloy containing seven percent magnesium. This latter theory

was then expanded by Harwood {2) to provide a comprehensive

description of the mechanism of stress- corrosion cracking

in alloy systems in general.

According to Harwood, the most probable process by

which stres s-corrosion cracking occurs is as f ollows:

"1. Localized elect rochernical corrosion occurs along

narrow paths prod uc ing trench-like fissures . It is most

likely that the advancing edge of these fi s sures have

extremely sharp radn of curvature, possibly of atomic dimen-

-2S-

si ons . . . . . . . . . . More than one such crevice may be produced,

but one usually sharpens and deepens to a greater extent than

the others.

2. As the fissure grows deeper and sharper a stress

concentration is developed at its tip .••••••••• At a

sufficiently high stress, localized plastic deformation

occ urs at the ti p of the fissure. This deformation which

is limited to the region ahead of the apex of the not ch,

initiates a britt le crack ••••••••••

3. Depending on the geometry of the specimen, rigidity

of the loadi ng fixture, test conditions, and certain energy

considerations inherent in brittle crack propagation, a crack

may propagate through the entire specimen, causing instant­

aneous (ca ta clysmic) failure or it may stop after pr ogressing

a finite distance ••••••••••

4. Mechanical extension of the c revice exposes clean

metallic surfaces, and the corrosive agent is immediately

drawn into the crack by capillary action. A period of rapid

corrosion then follows . It may weil be that this stage of

rapid corrosion aids in the penetration of the crack, but

lateral corrosion will also occur, resulting in branching

at each point of arrest. It seems reasonable t o believe that

the ma j or fact or i n penetration of a crack is the r esult of

mechanical action rather than electrochemical attack.

-29-

Acceleration of corrosion rate, as a result of

exposure of unfilmed metal surfaces to a corroding environ­

ment, rapidly decreases owirg to polarization and re-formation

of films, caused by electrolyte concentration changes at the

narrOlJI tip of an arrested crack.

6. Conditions similar to stage l no'!Jv prevail again,

and slow localized corrosion continues until a sufficiently

high stress concentration is produced which initiates deforma­

tion and crack formation. The entire cycle of events is

repeated until failure occurs because of crack propagation,

or the reduction of the load-bearing cross-sectional area.n

The hypothetical mechanism described above involves

two distinct stages. The first of these is a period of local­

iz ed corrosive attack which is then followed by the cracking

stage. In the event that limited failure occurs, then a period

of rapid corrosion is also involved. It is n~1 proposed to

consider each of the two main stages in sorœ detail, emphasis

being placed on the mecha~ism of stress-corrosion cracking in

th e cartridge brasses.

-30-

3.2 Nature of Localized Corrosion

The nature of the localized-corrosion stage of

stress-corrosion cracking has been well established as being

electrochemical. It has been demonstrated by Priest, Beek

and Fontana (36) for magnesium based alloys, Edeleanu (37)

for aluminum based alloys, Parkins (3S) for steels and Mears,

Brown and Dix (32) for austenitic stainless steels, brasses

and a magnesium based alloy, th at application of cathodic

protection prevents the initiation of cracking. It has been

shown also that cathodic protection will arrest cracking once

it has started (36).

The effectiveness of cathodic protection in preventing

cracking of brass is demonstrated by the work of Mears, et al.

(32). Stressed specimens of cold-~~rked 70:30 cartridge brass

were c oupled t o pieces of various sheet metal and immersed in

concentrated arrunonium hydroxide solution containing 2S percent

NH3. The period of time after which failure occurred was

measured in each case. The results of these workers are presented

in Table 1.

The results indicate that if the contacting metal

is sufficiently anodic to the brass then cracking may be prevented.

Less anodic metals reduce the rate of cracking whilst cathodic

materials, such as nickel, increase it.

-31-

TABLE I

Contacting Metal Potential Difference Time ta Failure

Zinc -0.829 volts Did not fail 1615 hours.

Cadmium -0.416 " Tin -0.280 456 hours

Le ad -0.073 360

70:30 brass o.ooo 193

Nickel +0.010 145

A further significant fact is provided by the work

of Parkins (38). He established that the minimum applied

current needed for complete cathodic protection was higher

for a cold-worked steel specimen stressed to 54 ,000 p.s.i.

than for a specimen of the same steel that had not been cold­

worked and was stressed to 40,000 p.s.i. Thus it was demon­

strated that stress and/or cold-~·lOrk is of importance to the

corrosion reaction.

Further evidence of the electroc hemical na tu re of

in

stress-corrosion cracking is provided by the f act that in s orœ

cases failure may be pre vented by the use of corrosion inhibitors

(39).

-32-

Intense localized attack of an electrochemical

nature presupposes the existence, at the surface of the alloy,

of a number of large cathode - small anode local cells. It

would appear from the lit erature that the follm'ling may pro vide

small anodic regions conducive to the formation of local cells

and resultant electrochemical corrosion:

l.

2.

).

A metallurgical phase present in small proportion.

Grai:1 boundaries where the Energy is high due

to intense disorder.

Grain boundaries or other substructural boundaries

at which solute atoms may segregate.

Regions where surface films have been ruptured.

Areas where plastic deformation has occurred.

As far as the locali zed corrosion of cart ridge brass

is conc erned, sorne di ffi cult y ha s be a1. experienced in determin­

ing the exact nature of the local anodic areas.

-33-

3.3 Localized Corrosion of Cartridge Brass

As previously noted the path of stress-corrosion

cracks in cartridge brasses is preferentially intergranular,

although cold-work and the presence of certain alloying elements

tends to cause cracking to become transgranular. Consequently,

if the theory of stress-corrosion cracking as oultined so far

is to be considered correct, then at the intersection of

grain boundaries wit h the surface of cart ridge brass, anodic

regions must exist. It is pertinent therefore, to consider

the evidence for, and possible causes of, anodic behaviour at

grain boundaries in cartridge brass.

Dix (40} reports an experiment performed by Brown.

Two large-grained specimens of 70:30 brass were taken: the

grain boundary zones on one were masked with wax or varnish;

on the other the grain areas were masked. The specimens were

then immersed in a one percent ~~onium hydroxide solution.

It was found that an open circuit potential difference of

0.035 volts existed between the grain areas and the grain

boundary zones, the latter being anodic to the forrœr. Upon

closing the circuit a cur rent of 0.19 milliamperes flowed

between the specimens.

A possible cause of this anodic behaviour at the grain

boundaries could be the presence of a second phase. This is

believed to be the case for many alloy systems in which stress-

-34-

corrosion cracking occurs in an intergranular manner. Either,

a phase precipi tated at the grain boundaries, or narrow

regions in close proximity to the grain boundaries which

have been depleted of solute atoms, are anodic to the bodies

of thegrains. Localized electrochemical corrosion may then

take place, resulting in stress-corrosion cracking, as for

example in sensitized austenitic stainless steels and magnes­

ium-aluminum alloys.

Although there is no microscopical evidence that

grain boundary precipitates exist in cartridge brasses, it

is obvious that sorne attention should be paid to this

possibility. It has been shown (18) that no added element

will increase the susceptibility of cart ridge brass, therefore

any such precipitated phase must occur in the copper-zinc

system. It has been suggested by Harrington and Jester (41)

that precipitation of the beta-phase may take place in what

are normally regarded as homogeneous alpha brasses, under

certain conditions. The authors point out that assuming the

same composition, the beta-phase has a higher densi~ than

the alpha-phase. Consequently at high pressures the alpha-

alpha plus beta phase boundary of the c opper-zinc equilibri um

diagram {see Fig.l) will be shifted to la-rer zinc concentrations.

It is surmised that the residual stresses from cold-working would

provide the high pressure necessary for this shift to occur.

-35-

The theory was proposed to account for the increase in hardness

and electrical conductivity that occurs when colà-1.vorked

alpha brass is subjected to annealing below the recrystalliza­

tion temperature. However, even if this theory is correct,

the precipitation of the beta-phase cannot be regarded as an

important factor in the stress-corrosion cracking of cartridge

brass since annealed brass, in which no residual stress

pattern exists, is susceptible to such crackin~ when under

the influence of externally applied stresses.

An alternative proposal to accoWlt for the anodic

behaviour of grain boundaries in cartridge brass is the

disoràer thereat due to orientation differences between adjacent

grains. However, the same disorder exists at the boundaries

in pure copper, which may be regarded as immune t o stress­

corrosion cracking. As the addition of small amounts of

phosphorus, arsenic, antimony, aluminum, silicon, nickel or

zinc in solid solution in copper gives susceptible alloys (42),

it would appear that the alloying elements are specifically

involved in causine; the potential difference between grains

and grain boundaries.

It seems, therefore, that localized corrosion at

grain boundarie s in cartridge brass is caused by a composition

difference between the bodies of grains and their boundaries.

-36-

Several attempts have been made to establish this difference

experimentally for various solid solutions. The results,

generally, have been inconclusive. This has been primarily

due to the fact that small excess quantities of solute were

determined as differences between large concentrations (22).

Dean and Davey (43) studied solid solutions of

copper in zinc. Metal was removed from grain boundaries

by electrolytic etching and analysed for copper spectrographi­

cally. Their results appear in Table 2:

TABLE 2

Copper Content

Specimen Number Wet Analysis of Spectrographie Analysis

Bulk Sample

Grai n Body Grain Boundary

3a 1.07% 1.1% O. $% 3b 1.07 1.1 0.7 6a 2.07 2.1 1.6 6b 2.07 2.1 1.$ 7a 1.65 1.6 1.5 7b 1.65 1.6 1.4

The results shcw clearly that the grain boundary

material had a distinctly lo\'rer copper content than metal

removed from the bodies of the grains.

-37-

The authors suggest tentatively that their results

might be interpreted in terms of the Gibbs isotherm, whicb

may be expressed as follows:

where u =

c =

u = ... c dt RT • dC

••••••••••••••••• ( 1)

the excess of solute in the surface layer per

square centimetre of surface.

the concentration of solute in the solution.

t = the surface tension of the solution.

T =

R =

the absolute temperature.

the gas constant.

From Eq. (l) it may be seen that if the surface

tension increases with increase in solute concentration, then

u is negative. Therefore the surface concentration is less

than that in the bulk of the solution. If it could be shawn

that the surface tension of copper is great er t han tha t of

zinc at 410°C (the equilibrating temperature used by the workers),

then use of the Gibbs isotherm might be justified.

The only relevant surface tension data (44) quoted

by the authors is as f ollows:

Surface tension of zinc at 450°C = 755 dynes/cm.

Surface tension of copper at ll40°C = 1120 dynes/cm.

-38-

This evidence may not be regarded as gi ving a certain

indication as to the relative surface tensions of copper and

zinc at 410oc, for the following reason. The surface tension

of copper a9parently increases with temperature whilst that of

zinc decreases (45). However, as a general rule metals with

higher melting points have the higher surface tensions.

The authors point out that if their reasoning is

correct, then an excess of zinc might be expected at grain

boundaries in sol id sol ut ions of zi ne in copper.

An investigation was carried out by Clifton and

Smith (46) upon the composition differences in a bronze

containing 1.4 percent tin. They eut slices 0.02 mm. wide

from positions progressively removed from a grain boundary.

They found that there was less than 0.1 atomic percent differ-

ence in the composition of the slices.

More recently, Thomas and Chalmers {47) investigated

the segregation of radioactive polonium, present to the extent

of one part in 1010 , in a lead alloy containing five percent

bismuth. Segregation at grain boundaries was sho\m to occur,

its extent being dependent on the orientation of adjacent

grains and the temperature at which equilibrium was attained.

The concentration of polonium decreased rapidly with increase

in temperature, and it was only in the high angle boundaries

that the segregate persisted when temperatures were high.

-39-

Postulating that in a solid solution the component

with the lower surface tension is positively adsorbed at the

grain boundary, Speiser and Spretnak (48) have suggested

one of the most probable theories of localized grain boundary

corrosion in alpha brasses. Assuming that copper has a higher

surface tension than zinc, the latter will be positively

adsorbed at the grain boundaries. In ammonia solutions, the

following reactions involving complexions are possible (49):

Eo = - 0.45 volts Cu 2 [zn(NH3)J ++ 2 Zn + 8 NH3 + Oz + 2 H20 = + 4 OH- ••••• (3)

Eo Zn

= - 1.43 volts

where E0 = standard electrode reduction potential.

Zinc is thus shown to be the more active electro-

chemically in ammonia solutions, and being preponderant at

the grain boundaries would account for the anodic behaviour

reported by Dix (40).

The authors point out t hat a solid solution of zinc

in copper exhibit s a large negative deviation from Raoult's

law, so that t he effective concentration of zinc as far as

chemical re actions are concerned is considerably less than

the mole fraction. Consequently, the concentration of ziœ

must be large for rapid grain boundary corrosion to occur.

-40-

This would verify the summary provided by Cook (13) to the

effect that service falures of cartridge brasses are prac­

tically limited to alloys containing 20 percent or more of

zinc, although certain circumstances can cause cracking in

alloys with as little as three percent zinc.

The above theory, of course, takes no account of the

accelerating effect of small amounts of carbon dioxide in

the corrosive atmosphere. However, it is likely that carbon

dioxide may take part in reactions involving complex ions

similar to Eqs.(2) and (3), the standard electrode reduction

potentials for the se reactions being even more divergent than

those stated.

Assuming that the above mechanism provides a

plausible theory of localized intergranular corrosion; the

problem remains of the origin of the localized corrosion

which gives rise to transgranular cracking. The initial

attack must be i~ the body of the grain since there is no

evidence that a transgranular crack can develop from an

intergranular corrosion fissure, or vice versa (2).

Investigation of the cracking of single crystals

of brass (14) (20) has not revealed any crystallographic

dependence of fracture, the crack surfaces being per­

pendicular to the stress direction. However, observation

of the initiation of localized corrosion in brass is not

-41-

easy due to the tarnishing effect of the ammoniacal

atmosphere.

Ideally, transgranular cracking studies should

be performed on an alloy on ~1hich opaque films are not fo:rmed

and which require no re-polishing before inspection, as this

may result in loss of sigaificant evidence. Bakish and

Robertson (50} found that copper-gold alloys subjected to the

attack of aqueous ferric chloride, which dissolves copper

preferentially, were suitable for observing structure­

dependent activity. These workers prepared single crystals

of a sol id solution allo y contai ning 48.9 percent copper,

corresponding to the formula Cu3 Au, and applied the corrosive

solution to a polished surface which was viewed under a micro­

scope. They found that at least two types of structural

site became active. One type was due to imperfect growth of

the crystal and was too small to be analysed in detail. The

ether type of active site was produced by deformation.

By viewing areas of the crystals that appeared to be

free of growth imperfections, i t w as observed that at small

strains of about two percent individually resolved sites of

activity appeared. These sites were invariably associated

with deformation structure, most of them appearing in clusters

of slip bands. Occasionally a si te appeared in a single slip

band. Further straining of the crystals revealed that each

active si te nucleated a small crack w hic h grew with increase

-42-

in strain in a direction perpendicular to the principal

stress axis. Further proof was thus provided that whereas crack

nucleation is structure-dependent, crack propagation depends

only on the stress direction.

Although the exact nature af the structure at an

active site in a slip cluster is not known Bakish and Robertson

(51) were able to obtain sorne indication. A single crystal of

the copper-gold alloy was strained approximately five percent.

As before, active sites appeared in the surface traces of slip

bands when the crystal was immersed in ferric chlorid e solu­

tion. The crystal was then sectioned along a slip plane and

it was seen that copper had been preferentially removed from

two types of structural path:

(i) Traces of a second set of slip planes in the

primary slip plane.

(ii) Curved traces which were possibly strained regions

associated with the g eneration of dislocation loops

during plastic deformation.

Although the work described applies to an alloy in

the copper-gold system, it seems likely that the inferences made

will apply equally well to the alpha brass es, due to the simi­

larities in behaviour between the two. Consequently it is

assumed that the active sites that will give rise to localized

corrosion i n t he bodies of g rains of cartridge brass will dev­

elop at growth imperfections or in clusters of slip bands.

-43-

In addition sorne sites may occur at single slip bands.

-44-

3.4 Crack Propagation

Although the electrochemical nature of the localized

corrosion stage seems well established, cons i de rab le doubt

remains r egardinG the mechanism of crack propagation. Basic­

ally, as noted in Section 3.1, the problem is whether crack

propagation is purely electrochemical or electrochemical­

mechanical in nature.

Supporters of the electrochemical theory hold that

the function of stress is to rupture surface films and/or

rend er material at the tip of the corrosion fissure or advanc­

ing crack more anodic. In fact, film rupture may play a more

fundamental part than merely to allow accelerated corrosion

to occ ur. The re is sorne evidence (52) (53) (54) tha t the

presence of surface films can reduce creep rate and pre­

sumably the rate at which other deformation processes occur.

It is postulated (55) (56) that adherent films act as barriers

to the movement of dislocations. Breaking or removing the

films allows the dislocations to proceed in their original

directions, thus producing deformation.

There is evidence (57) that stressed metal is anodic

to metal in the unstressed state, a fact which could cause

accelerated corrosion a t the tip of a crack. Edeleanu (5S),

whose views may be regarded as typical of those workers favour­

ing an electrochemical-mechanical mechanism, points out however,

-45-

that the difference in potential between defonned material

at the tip of a crack and the bulk of the alloy can only be

of the order of a few millivolts. He states that this small

voltage difference can not be expected to accelerate corrosion

to the necessary rate, especially as the solution in the crack

must have a high resistance. It would seem t h erefore, that

although sorne stress-assisted corrosion can occur, this alone

cannet be responsible for crack propagation.

The observed manner of propagation of stress­

cor rosion cracks is of interest. Gilbert and Hadd en (35),

Edeleanu (37) , a '1d Farmery ace ording to Evans {59), state

that stress-corrosion failure is discontinuous, the crack

propagating by a seri es of limited fractures. Pri est, Beek

and Fontana (36), by means of motion pic t ure microscopy,

showed that a plastic def ormation wave precedes the tip of an

advancing crack. In the course of microscope studie s of the

cracking of single crystals of alpha brass, Edeleanu (58)

obeerved that once a crack had started a slip line was usually

present at the tip. After a p eri od of time a faint extension

to the crack would suddently become apparent in its entirety.

At the same t i me a new slip step bec ame visible at t he new tip

or an exis ting slip step became more pronounced. The suggestion

is made that this evidence is consistent with a two-stage

mechanism of cr ack propagation. The au t hor propos es that a

stage of slow chemical embrittlement pre cedes rapid brittle

-46-

crack propagation through normally ductile material.

Forty (30) submits a tentativ e explanation of the

mechanism of the two stag es proposed by Edeleanu. He refers

to tm work of Graf ( 20) who concJ.udes th at in systems

susceptible to stress-corrosion cracking it is always the

least noble component of tre alloy which is dissolved prefer­

entially. In the case of cart ridge bras s stressed in an

ammoniacal environmen.t, dezincification will occur, thus inject­

ing vacancies into the active regions of the alloy surface.

These vacancies may form voids, either by aggregation or by

reaction with dislocations. This restricts t he plastic defor­

mation of the brass so that a crack can form.

Forty then notes that the crack wil l propagate in

a brittle manner providing that it has a vel oci ty great er

than that of dislocations at the tip of the crack which are

under the inf~uence of the concentrated s t resses. It is

proposed that if sorne co ndition exists that restricts the

velocity of these dislocations, then crack propagation can

occ ur. As stress-corrosion c r acki ng is belie ved not to occur

i n pure metals, it would appear that solid solution hardening

mieht exert the necessary restrictive effect. More specifically,

the cri terion for brittle failure could be t hat t he material

should exhibit a pronounc ed yield-point, since only t he initial

motion of the dislocations need be slow.

-47-

Accordiag to Cottrell (60), the yield point in

steels is best described in terms of an "atmospheren of carbon

atoms forming around dislocations. The upper yield point

is associated with the hie;h stresses necessary to force dis­

locations away from their atmospheres, in order that yielding

may occur. For steels the yield point is sharp, since the

retarding effect of the atmospheres is removed immediately

yielding occurs. For this reason, Forty thinks that this type

of yield point mechanism cannet account for the propagation

of stress-corrosion cracks unless these have a very high

velocity.

The author observes that a more prolonged restrictive

effect could be provided by a mechanism originally proposed

by Fisher (61). This mechanism considers that in the case of

substitutional solid solutions with high solute concentra­

tions, phase demains of short-range arder can slow down the

motion of dislocations . Passage of a dislocation through a

domain ch anges the configuration ac ross the slip plane and a

more random arrang ement of atoms of higher energy is produced.

A yield point can be expected as further passage of dis­

locations will result in a progressive increase in disorder

which will lessen the restrictive behaviour of the domain.

As a number of dislocations, depending upon the domain size

and the Burgers vector of tœ d i slocation will have to pass

before the strengthening effect is completely overcome, it

seems that slower moving cracks will be able to propagate

-48-

in this type of alloy.

This theory of crack propagation being possible,

providing that the movement of dislocations at the crack tip

can be retarded by phase domains of short-range order may

be used to explain one of the observations of Edeleanu (58).

As previously described, he observed that cracks seemed to

progress intermittently, being halted by existing slip bands

which sometimes became heavier. The inference is that the

material between slip bands is sufficiently strong or hard to

allow brittle fracture, whereas in the slip bands the alloy is

soft er and will deform rather th an c le ave. Presumably, the

short-ran8e or der which limits the motion of dislocation in

the regions between slip bands is partially or c ompletely

destroyed in the slip bands themselves, due to the prier

passage of dislocations. Consequently, when a crack reaches

a slip band the relative freedom of movement of dislocations

allows relaxation of stress to occur by plastic deformation

and brittle cracking stops. Further pro gress of the failure

can only be initiated by chemical embrittlement at the tip

of the crack.

The observations of Edeleanu (58) and the theory of

Forty (30) suggest that susceptibility to cracking might be

dependent, to sorne extent, on the amount of cold-work t o \'l'hic h

the alloy has been subjected. If the number of slip bands is

-49-

increased by pr evious working, then the number of stages in

cracking, and consequently the time to failure,is i:tcreased.

Eventually, the slip bands will themselves work-harden due to

the generation and interaction of dislocations so that they

will no longer provide a barrier to crack propagation. There­

fore it is to be expected that a curve s hovving the relationship

between time to failure and increasine cold-work would rise to

a peak corresponding to the maximum number of slip bands being

present. Thereafter the curve should decline owing to the

hardening of the slip bands.

The the ory of crack propagation so far di scuss ed has

applied to the particular case of traœ granular cracking.

However, if we accept the concept of large-angle grain boundaries

as postulated by Evans (59) then intergranular cracking can be

explained in the sarne terrns. Evans suggests that grain boun­

daries are mainly regions of disorder, but intermittently,

ttbridges" of c ont inuous ordered ma terial may link the adjacent

grains. Assurning that cracking is initiated by chemical

embrittlement of an intercrystalline nature, the crack could

propagate along the grain boundary due to the hindering effect

of disordered material upon dislocations. At the continuous

bridges no restriction of dislocation movement would occur and

as a result the crack would halt. Further progress would

require that chemical ernbrittlement of the bridge should take

place.

-50-

4. EXPERIMENTAL PROCEDURE

4.1 Introduction

The practical work undertaken was a study of the

stress-corrosion cracking behaviour of cold-rolled cartridge

brass strip, a material wh ich cracks readily. The inves­

tigation involved the application of pre-corrosion tests, where­

by specimens were subjected to the corrosive environment for

a defini te period prior t o mechanic al tes ting. In the se tests

the operative tensile stresses which tended tocause crackine

were residual, having remained from the cold-rolling of the

strip. Associated with these tests, metallographie techniques

were used to establish the rnicrostructural features exhibited

by stress-corrosion cracks in the strip material.

A subsidiary investigation was performed using the

"loop test". This is a test of a type which is commonly

used in industry to measure the relative susceptibility of

alloys to stress-corrosion cracking , vmilst under the influence

of applied stresses. The main purpose of this investigation was

to establish th e reproducibility with which the test could be

applied to cartridge brass strip, and to determine, if possible,

any controls that would improve the precision of the results.

-51-

4.2 Materials

The test materials used were three different

gauges of cartridge brass strip supplied by Noranda Copper

and Brass Ltd., Montreal East, P.Q.

All three materials met the requirements, with

regard to composition and properties, of cartridge brass sheet,

strip, plate, bar and dises as specified in ASTM Sta'1.dard Bl9-55.

Hence, the composition of the test materials was in

keeping with the follm'ling specifications:

Cu 68.5 - 71.5%

Pb (max.) -

Fe (max.) -

0.07%

0.05%

Total elements other

than Cu & Zn (max.)- 0.15%

Zn re mai nd er.

Each of the three gauges of strip had been reduced to

a thickness of 0.200 in. and annealed before the final cold

reduction. The thickness of each strip, along with its per­

centage of cold-reduction and identifying symbol, are recorded

L'l Table 3.

Each strip so received was six inches wide and

eut into sections sorne three feet long. From these, a

-52-

number of specimens six inches long by approximately half an

inch wide were eut, using a pcwer saw with adequate water cool­

ing. Specimens were prepared with the long dimensions either

parallel or perpendicular to the rolling direction.

TABLE 3 t f

Identifying 1'hickness Pere en tage Symbol Reduction

A 0.083 in. 58.5% B 0.063 68.5 c 0.026 87.0

-53-

4.3 Pre-corrosion Tests

Accelerated stress-corrosion cracking tests for

brasses have been widely used in industry. These tests can

be divided into two types: those employing mercury or

aqueous solutions of mercury salts as the corrosive medium

and those employing ammoniacal atmospheres. The mercury

test cannot be regarded as absolutely suitable for two

reasons. The first of these is that mercury cracking is

not, correctly speaking, stress-corrosion cracking . Sec ondly,

there appears to be a threshold stress associated with mercury

cracking which is not apparent with ammonia cracking. Con­

sequently th ere is no certainty that a brass part which does

not crack during the course of a mercury test, will not crack

in the presence of a rrmonia during service.

Basically, two methods of producing the nec essary

atmosphere for an ammonia test are available. Cylinders of

the required gases or alternatively aqueous ammonia solutions

can be used. The former system is most flexible as the con­

centrations of a number of gases can be precisely regulated.

On the other hand, the second method is simpl er as far as the

apparat us required is concerned. As it was only necessary to

ensure that conditions were reproducible, a test of the second

type was used.

-54-

The test selected has been described by Jamieson

and Rosenthal (62) and entitled by them the "Aqua Ammonia

Test". The authors point out that in a closed system con­

taining an ammonium hydroxide solution, the vapour pressures

of H20 and NH3 are a function of the temperature and concentra­

tion of the solution. If the temperature is held constant

the partial pressures of H2o and NH3

are a function of solution

concentration only. If the system is vented to the atmos­

phere by a capillary tube the partial pressure of air is then

determined by the diff erence between atmospheric pressure

and the sum of the partial pressures of H2o and NH3 above

the solution. Although independent control of the partial

pressure of each gas is not exercised, it is only necessary

to control the concentration of the a:nmonium hydroxide solution,

the temperature and the ratio of solution volume to the con­

tainer volume in order to determine the composition of the

corroding atmosphere.

Two types of contai ners were used for this test.

The first type was a glassconical flask closed with a rubber

stopper containing a glass capillary tube. At the begi nning of

each test a measured volwne of ammonium hydroxide solution

(28-29 percent NH3, S.G. = 0.9016 at 60°F) was placed in the

dry flask. One specimen only was tested at a time, and this

was placed in an almost vertical position in the flask being

supported by the glass neck.

-55-

The second type of container was a 10 in. diameter

laboratory desiccator with a central opening in the cover.

The opening was fitted with a rubber bung containing a glass

capillary tube a11d at the cormnencemen t of each test a measured

quantity of the solution was poured into the space below a

perforated ceramic plate. Up to three specimens were tested

concw~rent.ly, being supported in a horizontal position by two

glass rods which rested on the perforated plate.

The test specimens were prepared as follmvs. The

burrs from the sawing were removed and the centre of the

specimen was necked by fi ling. In most cases, but not all,

the minimum width at the neck was 0.335 in. ± 0.005 in. The

specimen was then degreased by placing it for approximately

one hour in warm water containing a household detergent.

After degreasing the specimen was washed thoroughly in cold

water and dried. For a distance of approximately two inches

at each end the strip was covered with an insulating material.

This was usually bees-wax although in so me tests paraffin

wax or cellulose tape was used. The specimen (Fig.2) was then

ready to be placed in the corrosion vessel. All tests were

performed at room temperature which did not vary more than two

degrees Centigrade from 25°C.

After exposure to th e corrosive atmosphe re for a

definite period the specimen was removed and thoroughly rinsed

Fig. 2

-56-

T

insulating material

Test speciman for pre-corrosion test.

-57-

in colà water. The insulating material vras renoved, the

strip was washed successively in water and acetone, and then

dried in a blast of warm air.

The specimen was then tested on the Hounsfield

Tensometer (see Appendix I). The value of theload required

to fracture the specimen, along with the ultimate tensile

strength of the material as measured on the same machine,

were used to determine the amount of damage it had suffered

due to stress-corrosion cracking.

-58-

4.4 Loop Tests

The conduct of these tests was based on the

procedure detailed by Thompson (3) •. The test specimens were

samples of strip six inches long by approximately half an

inch wide. The edges of each specimen were smoothed with a

file. The specimen was then bent around a three quarters of

an inch diameter mandrel until the ends touched, whereupon

they were fastened ~lith copper wire. The loop of strip thus

formed was degreased in a war.m detergent solution for approxi­

mately one hour and thoroughly washed in cold water. The

specimen was then dipped in distilled water before being sus­

pended by a copper wire hook in a glass battery jar. A

known volume of concentrated ammonium hydroxide solution was

introduced into the jar and a loosely fitting cover was placed

over the top. The time for fai lure to occur was noted.

Although the test has the advantage of simplicity

and allows the progress of cracking to be studied it has the

disadvantag e that the level of stress in the specimen is

unknown. Consequently, its value lies mainly in its use for

comparis on purposes.

-59-

4. 5 l\1etallography and Photomicrography

Specimens of the test materials tha t were to

be examined under the microscope were prepared by grinding

wi th emery pap ers follo\.'red by polishing on laps wi th carbo­

rundum and alumina.

Etching of the brass specimens was difficult, the

problem being to outline the structure satisfactorily with­

out causing excessive darkening of the deformed metal.

Eventually, the etchant used was a mixture of a 50 percent

solution of concentrated ammonium hydroxide in vvater and a

three percent solution of hydrogen peroxide in water, in

the ratio of ten to one.

Photomicrographs were taken on Kodak 1Jï plates in

a Bausch and 1omb Metallograph, using a blue filter. The

plates were developed in Kodak Dl9 and fixed in Kodak Acid

Fixer. Prin ts "'Jere made on Kodak F2 paper us ing Kodak MQ

Developer and Acid Fixer.

-60-

5. RESULTS AND DISCUSSION

5.1 Mechanical Properties of Materials

As a preliminary to conducting the pre-corrosion

tests i t was necessary t o establish the ultimate tensile

strength of all three materials, both in and at right angles

to the direction of rolling. Specimeœ were trinrned and

necked by filing and the minimum width measured by micro­

meter. As no corrosion was involved it was not considered

necessary to keep w ithin the tolerances for minimum width

that were set for the pre-corrosion test pieces.

Each test piece was pulled in the Hounsfield

Tensometer and a facsimile of the stress-strain diag ram

obtained. The maximum load applie d during tre test was

obtained from the graph. From the known thickness of the

strip and tre measured width at the neck before testing, the

ultimate tensile strength was c alcuJa ted. The relevant data

is subrnitted in Table 4.

The results indi cate that for each material the

direction in which the ultirnate tensile strength is measured

makes a considerable difference to tre value obtained. The

difference increases with increase in cold-1:rork ing , so that

in the cas e of the most heavily reduced material, c, the

differentia! amounts to sorne 15,000 p.s.i.

-61-

TABLE 4

Test Material Direction of Width at Maximum U.T.S. No. Specimen Neck Load

C3 A Parallel to dir- 0.352 in. 1.22 ton 93,500 p.s.i. ection of rolling

B4 A Perpendicular to direction of rol- 0.355 1.29 98,200 ling

C5 B Parallel to dir- 0.340 0.93 97' 300 ection of rolling

Bl B Perpendicular to direction of rol- 0.336 0.98 103,800 ling

C7 c Parallel to di r-ection of rolling 0.351 0.43 105,600

B8 c Perpendicular to direction of rol- 0.314 0.44 120,700 ling

This variation of strength wi th direction of t esting

in sheet materials has been reported by Gohn and Arnold (63).

Their data for the nearest equivalent to cartri dge brass

tested namely, "best spring brass" containing 72 percent copper

and 28 percent zinc is summarized in Table 5.

For purposes of comparison the two sets of data are

plotted in Fig . 3. It will be noted that although Gohn and

Arnold' s measurements were made over a different ran ge of

percentage reduction, where the curves overlap, the differ­

entiai i s approximately the sane.

-62-

TABLE 5

Percent age Direction of Specimen U.T.S. Reducticn

37.1% Parallel to direction of rolling 74,200 p.s.i.

37.1 Perpendicular to direction of 76,800 rolling

60.5 Parallel to direction of rolling 94,4.00

60.5 Perpendi cular to direction of 99,600 rolling

Fig.3 also contains a curve plotted from data

relating ultimate tensile strength to percentage reduction,

as quoted in the ASivi Metals Handbook (64). The figures

given relate specifically to flat products, however, the

relationship between the axis of the test specimen and

direction of rolling is not given. The position of the

curve would indicate that the values of ultimate tensile

strength quoted are for directions parallel to the rolling

direction.

UTS, psi. 12o,ooo

110,000

100,000

90,000

80,000

70,000

-·-·---~--0-0-

-IHl-

30

Perpendicul:-ar to direction of rolling (C.D.E.)

Paralle1 to direction of rolling (C.D.E.)

PerpendiGular to direction of rolling (G & A)

Parallel to direction of rolling (G & A)

ASM Handbook

40 50 60

~

./ •

•

Fig. 3 Relationships between percentage reductio~ and ultimate tensile strength for cold-rolled cartriQse brass strip.

70 80 90 % ReductiQll.

l OO

1 0'­\..ù 1

-64-

5.2 .r-1icrostructure of !vlaterials

Longitudinal sections and sections parallel to the

plane of rolling were prepared from each material. The sec­

tions were polished, etched and photographed at 200 magnifica­

tions in accordance with the procedure outlined in Section

4. 5. The photomicrographs of each section appe ar as Figs.

4 - 9.

Inspection of the microstructure of each material

revealed, as was to be expected, a marked deformation struc­

ture, which was mœt pronounced in material C. The longi­

tudinal sections showed that considerable flattening and

elongation of tœ.grains had occurred. Many of the grains

were so deformed that the et chant had darkene d them con­

siderably and no features were discernible. The more lightly

etched grains, in many cases, showed traces of very pro­

nounced slip bands.

The sections taken parallel t o the plane of

rolling showed that cons i d.erable slipping had oc curred.

Most grains showed traces of heavy slip bands which tended

to occur in a direction that was approximately perpendicular

to the di rection of rolling . An interesting feature of the

specimen taken from material B (Fi g .7) is that one grain

i s divided by a line running i n a dire ction th at i s nearl y

parallel to the direction of rolling; one one si de of the

-65-

line heavy slip bands appear, on the other no slip bands are

visible. Presumably, this is an example of an annealing twin

resulting from the previous strain-anneal history of the strip.

In one half of the twi.n a set of slip planes is suitably

oriented so that slipping can occur. On the other side of

the twinning plane, the orientation of the se planes is such

that slipping has not taken place, Other evidence of the

existence of annealing twins is present in the micro­

structures of all three materials.

-66-

•

Fig.4 Material A, longitudinal section, X200

Fig.5 Material A, parallel to rolling plane, X200

-67-

Fig.6 Material B, longitudinal section, X200

Fig.? Material B, parallel to rolling plane, X200

-68-

Fig.8 Material C, longitudinal section, X200

Fig.9 Material C, parallel to rolling plane, X200

-69-

5.3 Loss of Strength Relative to Rolling Direction

Specimens of a 11 three ma teri als, taken both

parallel and perpendicular to the direction of rolling,

were subjected to pre-corrosion tests. The specimens

were exposed individually in conical. flasks for 24 hours,

the percentage volume occupied by the concentrated ammonium

hydroxide solution being $.7 percent.

In order t o get a me as ure of the dan age due to

stress-corrosion cracking that had occurred, a quantity

designated "apparent tensile strength" was calculated. The

apparent tensile strength was derived as follows:

Apparent tensile strength = Maximum load Cross-sectioŒll area of specimen at fracture

The lower the calculated apparent tensile strength

compared with the ultimate tensile strength of tœ material,

then the greater the damage due to cracking.

Generally, when tested on t he Hounsfield Tensometer

the specimens broke at the narrov1est section. Consequently,

for these specimens the value of apparent tensile strength

was calculat ed from the measured width at the neck. If, for

any reason, a specimen failed other than at the neck, the

width of the specimen was measured at the fracture and this

value was used in calculating the apparent tensile strength.

The data resulting from this test is presented in Table 6.

The relevant strength da ta for the uncorroded ma terials i s

Test Material Direction of Speci-No. men

C4 A Parallel to direc-tion of rolling.

B5 A Perpendicular to direction of rollin8 .

ClO B Parallel to direc-tion of rolling.

B7 B Perpendicular to direction of rolling

Cà c Parallel to direc-ti on of rolling.

B9 c Perpendic ular to direction of rolling.

~.........-..._ ___ -- ---- --- -- - - - -

TABlE 6

Width Width at at Fracture

Neck

0.334 in. 0.334 in.

0.336 0.336

0.336 0.336

0.294 0.294

0.334 0.375

0.323 0.323

------ ·--- - - -- - --

Maximum Ap:ça.rent Load Tensile

Strength

0.68 tons 54,900 psi

1.19 95,600

0.91 96,400

0.86 104,000

0.355 81,700

0.45 120,000

U.T.S.

93,500 psi

98,200

97,300

103,800

105,600

120,700

1

-

1 -,J 0 1

-71-

included in this table for purposes of comparison.

Inspection of the results indi ca tes that, for each

material, the specimen eut parallel to the direction of rolling suff­

ered the greater loss in strength. Inspection with a low-powered

microscope provided the reason. Cracks had developed in a trans­

verse manner in these specimens ' whereas the cracks in the speci-

mens taken perpendicular to the direction of rolling were

longitudinal. It is concl uded therefore that as stress-

corrosion cracks progress along planes that are approximately

perpendicular to the opera ti v e tensile stresses, tre n the

residual stresses in the cartridge brass strip tested must

lie in the direction of rolling.

The results obtained are in direct contract to

those of Czochralski and Schreiber (65). In conducting

pre-corrosion tests of a similar nature on rolled brass sheet

of different compositions they found that deterioration in

tensile strength occurred more quickly in specimens taken

perpendicular to the direction of rolling. This, presumably,

is due to the pattern of residualstress being different in

the materials used by the se workers. It is conceivable that

sorne type of cold-rolling operation could leave internai

tensile stresses acting transversely rather than longitudi­

nally.

-72-

It is üOt kno-.;..;n t o what extent tre material s used

by Czochralski and Schreiber had been reduced. It is likely

that they were testing less heavily worked materials than the

cartridge brass strip utilized for this research, since they

report that in the alpha brasses cracking was intergranular.

As will be shown in a subsequent section, cracking in ma ter­

ials A, B and C was substantially transgranular.

-73-

5.4 Measurement of Residual Stress

The residual tensile stresses acting at the

surface of each type of strip, sho1~ to be longitudinal by

the experimental results reported in Section 5.3, were

measured us ing a f orm of t:œ Modifi ed Anderson and Fahlman

technique ( see Appendix II). Two approximated values of

the stress in each material were obtained, one by measuring

deflection and applying Eq. (12), the other by constructing

the radius of curvature and applying Eq.(S). The datais

swmnarized in Table 7. The values of Young's modulus of

elasticity and Poisson's ratio used in the calculations

were obtained from the Metals Handbook (64) arrl Field

Foster (66) respectively.

Inspection of the results reveals that material

A, the least heavily reduced, possesses the highest residual

stress. The stress in material B is indicated as being

slightly greater than that in material C. This latter dis­

tinction is a fine one and may not be correct in view of

the admitted inaccuracies inherent in deflection techniques

of stress measurement.

The distribution of residual stress in cold-rolled

strip has been explained by Baldwin (67). He states that this

distribution is a function of the ratio between strip thick­

ness and length of contact in the rolli~ direction between

TABLE 7

Materia1 Thickness, Young 1 s Poisson's Length of Def1ection, t Modu1us Ratio Samp1e, y

l

A 0.083 in. 16 x 106 psi 0.376 21/32 in. 0.029 in.

B 0.063 16 x 106 0.376 21/32 0.020

c 0.026 16 x 106 0.376 21/32 0.048

2E1tl Radius of Sl= 12 Curvature,

~

21,800 psi 2olj8 in.

11,400 247/8

11 '300 103/4

-

E1t s1~

19,000 psi

11,800

11,250

L__ - - - ---- -

1 .......;]

f

-75-

roll and strip, and tte refore of tre ratio of strip thick­

ness to roll diameter. If this ratio is large, then plastic

defonnation of the stock is confined to the surfa ce regions.

If the ratio is small, t hen plastic defonnati on extends

through the entire thickness.

In the fir st case, the surface layers are extended

in the rolling direc"t<ion by a greater amount than the central

zone. The surface layers are thus constrained by the central

zone and are held in longitudinal compression. Conversely,

the central zone is held in tension by the extension of the

surface layers. The second case, which applies to thin strip

produced by large diameter rolls, commonly results in longi­

tudinal tensile stresses at the surfa ce. One explanation

that has been proposed (68) suggests tha t in rolling thin

strip the surface layers are restrai ned t o move at the peri­

pheral speed of the rolls by frictional forces whe reas the

central zone flmvs plastically and is virtually extruded to

a greater elongation. The centre zone thus holds tre

surf ace lay ers in tension, being i tself held in longitudinal

compress ion.

The second condition, namely a small ratio of

strip thickne s s to roll diameter, obviously must have

applie d to the fabrica tion of the cartridge brass strip used

i:1 this work, th us accounting for tre observed long itudinal

-76-

tensile stresses at the surf ace. It remains to be explained,

however, as to why the least heavily reduced material, A,

possesses the highest residual stress. Baldwin (67) states

that when strip-rolling is carried out with a number of passes

the residual stress pattern is c ont rolled largely by the final

pass. He cites as an example the case of a bearing bronze

whic h had the sarne stress distribution whether i t had a t \\0

percent final reduction following a single pass giving a 16

percent reduction or a number of passes each giving a two per­

cent reduction, to the same size. Thus it would appear that

sorne knowledge of the cold-rolling schedule to which the

materials bad been subjected is necessary to explain the

differences in residual stress that were found to exist.

-77-

5.5 Effect of Stress Relief Annealing

Four specimens of mat erial A, two parallel and

oo perpendicular to the direction of rolling, were stress

relief annealed for 30 minutes at 250°C {482°F). One of

each pair was then tested on the Hounsfield Tensometer. The

other specimens were subjected to pre-corrosion tests in

vented conical flasks for 24 hours. The conditions of

exposure corresponded to those of the tests reparted in

Section 5.3. The values of u1timate tensile strength obtained

for the stress relief a'1nealed specimens are presented in

Table 8, the strength data of the as-rolled mat erial being

included for purposes of comparison:

TABLE 8

Test Direction of \<Jidth at Maximum U.T.S. U.T.S. No. Specimen Neck load as-rolled

Dl Perpendicular to direction of rol- 0.355 in. 1.37 ton 104,200 98,200 ling. psi psi

D3 Paral1e1 to direc-tion of rolling 0.284 1.01 96,000 93,500

From the resu1ts it may be seen tl:at stress relief

% increase

6.1%

2.7

annealing increased the u1timate tensile strength of the material, an

effect which has been wide1y reported in the1iterature but so far

has not been explained satisfactorily. Further, it appears that

-78-

the percent age inc rease is more pronounced in the direction

perpendicular to the rolling axis.

The resul ts of the pre-corrosion tests on stress­

relieved material are summarized in Table 9. The results

of pre-corrosion tests of the same duration performed on

the test material in the as-rolled condition are included.

In order to compare the effect of corrosion upon the strength

of the different specimens, the apparent percentage decrease

in tensile strength based on the ultimate tensile strength

of the uncorroded material has been calculated in each case.

The results show a decided decrease in suscepti­

bility to stress-corrosion crackiag on the :r:art of the stress­

relief a~nealed rnaterial. It seerns possible that residual

stress can never be completely eliminated by stress-relief

annealing, however, by using a higher temperature or longer

time than the relief treatment performed for this experiment

it should be possible to reduce the residual stress to a level

such that cracking would only be initiated after a very long

exposure to the corrosive environment.

!

Test Condition Direction of Speci- Width at No. men Neck

D2 Stress-relief Perpendicular to 0.336 in. annea1ed rolling direction

B5 As-rolled 0.336 " n

D4 Stress-relief Parallel to rol- 0.336 annealed ling direction

C4 As-rolled tT n 0.334

~---- ----------- - ---------- ---- - -------

TABLE 9

Width at Maximum Apparent Fracture Load Tens ile

Strength

0.336 in. 1.29 tone 103,700 psi

0.336 1.19 95,600

0.404 1.12 75,000

0.334 0.68 54,900

-------- '-----

U.T.S.

104,200 psi

98,200

96,000

93 '500

Apparent 1

% Decrease i

in U.T.S. 1

0.5% 1

1

2.7

21.9

1

'

41.4 1

-

1 .....:] \.0 1

-80-

5.6 Paths of Cracking

Specimens of materials A, B and C which had been

exposed to the corrosive environment for various periods

were sectioned, and polished and etched in accordance with

the procedure outlined in Section 4.5

Longitudinal sections of specimens of material A

and material C wh ich had be en exposed for 48 hours and

102à hours respectively, and which exhibited pronounced

cracking were photographed at lOO magnifications. The

photomicrographs are presented as Figs. 10 arri 11. No

photomicrograph was made of any section of material B as even

in the specimen which had been exposed for the longest time,

namely 102! hours, the extent of cracking was small and dis­

played no features which are not illustrated by Figs. 10

and 11.

In addition, the rolled surface of a sample of

material A was polished, etched and exposed to the corrosive

environment for 21 hours. It was re-polished on the fine

lap, re-etched a~d photographed at 100 magn ifications. This

photomicrograph appears as Fig .l2.

The most noticeable feature of the cracks, and

one which is apparent in all three photomicrographs, is that

they are basically transgranular. Occasionally a crack may

-81-

follow a suitably oriented grain boundary for a short dis­

tance, but it soon reverts to a transgranular path. This

observation provides supporting evidence for the conclusions

of Althof who, as quoted by Graf (20), states that increas­

ing plastic deformation of alpha brasses results in a greater

tendency for cracking to become transgranular.

Further, Figs. 10 and 11 reveal that whereas in

its early stages a crack will tend to propagate along a

plane, that is approximately perpendicular to the stress

axis and plane of rolling, after proceeding . some distance

i t will branch.

In the case of rnaterial A, the branching results

in the progress of two cracks, each of which tends to

follow initially a plane which makes an approximate angle of

45° to the original crack. As the two branches propagate

further the angles they make with the original crack increase,

until they eventually advance in a plane that is parallel

to the rolled surface.

With material C, the branching results in the

forma ti on of two cracks each of which immediat ely begins to

progress parallel to the rolling plane.

It is presumed that the se chang es in the direction

of crack propagation result from alt erations in the local

distribution of residual stresses cau s ed by previous progress

of the cracks.

-82-

Fig.lO Material A, exposed 48 hours, longitudinal section, XlOO

Fig.ll Material C, exposed 102i hours, longitudinal section, XlOO

-83-

Fig.l2 Material A, e~osed 21 hours, rolled surface, XlOO

-84-

From Fig. 10 it may be seen that associated with

the crack in material A there are a number of rounded dark

areas. The same effect is noted to a lesser extent in the

region of the crack in material C, shCMn in Fig.ll. These

regions are thought to be sites of localized corrosion which

has occurred during the propagation of the stress-corrosion

cracks. Admittedly, from the photomicrographs it would

appear that some of these regions are completely unconnected

with the cracks. H<::JW"ever, it must be remembered that the se

photomicrographs provide only a t 1-vo-dimensional view of a three­

dimensional process, and that connection betv1een the cracks

and these regions may exist in ~lanes other than the ones

that are visible.

,•

-85-

5.7 Relative Susceptibility of Strip Materials

Pre-corrosion test specimens eut parallel to

the direction of rolling were prepared from all three mater-

ials and exposed to corrosive conditions for kno~'ln periods

before being tested on the Hounsfield Tensometer. The

corrosion vessel employed for these tests was a laboratory

desiccator in which concentrated a ,moniwm hydroxide solu-

tion occupied 5.5 percent of the volume, and in which the

specimens were supported in a horizontal position.

In previous tests damage due to stress-corrosion

cracking was measured in terms of the apparent tensile

strength of the corroded specimen. Hcwever, this procedure

could not be us ed for these tests as it do es not talœ the

variation in thickness of the materials into account.

Consequently, it was decided that stress-corrosion damage

should be measured in terms of "effective decrease in cross-

sectional area" which was calcula ted as follows:

Effective decrease in cross-sectional area = Cross-sectional area of _

specimen at fracture Cross-sectional area of unaffectec material at frac­ture

= (thickness x width at fracture)

Maximum load Ultimate tensile strength of as-rolled material.

-$6-

The results obtained for materials A, B and C are

reported in Tables 10, 11 and 12 respectively. A graph of

time of exposure against effective decrease in cross-sectional

area for each material appears in Fig. 13.

Inspection of the graphs of stress-corrosion damage,

as represented by the effective decrease in cross-sectional

area, against time , reveals that for each material there

is a certain period of time during which no significant

cracking occurs. It is presumed that this may be regarded

as an induction period during which a film of corrosive

moisture forms on the surface followed by intense localized

attack which eventually initiates cracking.

It appears that the induction period for material

A is the shortest, being approximately four hours. The

induction periods for materials B and C are approximately

20 hours and 40 hours respectively. Since it is unlikely

that the time necessary for film formation will vary from one

material to another the difference between these times must

be r elated t o the speed wi th which anodic a reas in the diff­

erent mat e rials a re s uff ici en t l y co rroded t o i nit i a te cracks.

It will be noted that in the materials studied

the induction period i a creases, both with decreas e of

r esidua l tens i l e s tress at the s urfac e and 1r1ith increase

of plastic deformation. Consequently, it appears that two

Material - A,

Test Time of Width at Width at No. Exposure Neck Fracture

Jl 4 hr. 0.339 in. 0.339 in.

J2 7~ 0.337 0.345

J3 12 0.338 0.344

J7 12 0.340 0.340

J4 16 0.329 0.340

F2 24 0.338 0.338

E4 48 0.335 0.388

1........__ -- ----- - - - --- ----- - - ---~

TABLE 10

Thickness = 0.083 in., U.T.S. = 93,500 p.s.i.

Area of Cross- Haximum Area of Cross- Effective Decrease section at 1oad section of Un- in Cross-sectional Fracture affected Material A rea

0.0281 sq.in. 1.17 ton 0.0280 sq.in. 0.0001 sq.in.

0.0286 1.14 0.0274 0.0012

0.0286 0.78 0.0187 0.0099

0.0282 0.73 0.0175 0.0107

0.0282 0.87 0.0208 0.0074

0.0281 0.70 0.0167 0.0114

0.0322 0.60 0.0144 0.0178

--- ~--~ ----- -------- ----

1 CQ.

......;) 1

Test Time of Width at No. Exposure Neck

E2 24 hr. 0.338 in.

E7 46 0.337

E5 48 0.337

E9 72 0.338

G1 102~ 0.334

TABLE 11

Materia1 - B, Thickness = 0.063 in., U. T.S. = 97,300 p.s.i.

Width at Area of Cross- Maximum Area of Cross- Effective Decrease Fracture section at 1oad sec t i on of Un- in Cross-sectiona1

Fracture affected Mat eria1 Are a

0.338 in. 0.0213 sq.in. 0.915 ton 0.0210 sq.in. 0.0003 s q.in.

0.351 0.0221 0.90 0.0207 0.0014

0.337 0.0212 0.91 0.0209 0.0003

0.338· 0.0213 0.90 0.0207 0.0006

0.352 0.0222 0.87 0.0200 0.0022

·-- -- - - -- --- ------- ---- - - --

1

1

1

i

1 00. 00. 1

Mate rial - C,

!Test Time of \'lidth at vlidth at No. Exposure Neck Fracture

J5 12 hr. 0. 33 7 in. 0.337 i n .

E3 24 0.336 0.336

J6 36 0.340 0.340

ES 46 0.332 0.395

E6 48 0.334 0.414

Fl 48 0.404 0.404

ElO 72 0.337 0.402

G2 102~ 0.335 0.337

TABLE 12

Thickness = 0.026 in.,

Area of Cross- Maximum section at load Fracture

0. 0088 sq. in. O. 42 ton

0.0087 0.41

0.0088 0.42

0.0103 0.35

0.0108 0.37

0.0105 0.485

0.0104 0 .365

0.0088 0.205

-

U.T.S. = 105,600 p .s. i.

Area of Cross- Effective Decrease 1

section of Un- i n Cross-sectional affected Material Are a

0.0088 s q.in. nil

0.0087 nil

0.0088 nil 1

0.0074 0.0029 sq.in.

0.0078 0.0030

0.0103 0. 0002

0. 0077 0. 0027

0. 0043 0. 0045

--· - ---- -- - - -- -

1 00. \.() 1

N s:: '" ~

Id cv .;: M

~ 0

'" 4J 0 cv (Il

1 (Il

Ill 0 1-1 u s:: '" cv Ill Id cv ~ cv ~

cv > '" 4J

·•

• 0200

.CH 50

.0100 ';

~ .0050 4-1 4-1 f;l::l

... 0

0

• • •

20 40

--~

Fig,,l3

••

60

Material A •

Material B.

Material C.

Relationships between effective àecrease in cross-sectional area aad time of exposure.

)(

80 100

Time of Exposure, hrs.

1 '-[) 0 1

-91-

explanations of the variation in the length of the induction

period are possible. If the localized corrosion is stress­

assisted then it would seem axiomatic that the material

with the highest internal stress should be attacked most

quickly. Alternatively, if the sites of anodic attack prior

to transgranular cracking are slip bands, a possibility

discussed in Section 3.3, then corrosion will proceed at the

fastest rate when slip-baads are small i~ number. A more

heavily deformed material with many slip bands will suffer a

decrease in the ratio of cathode area to anode area, and

sites of less intense attack will be generally distributed.

In vie w of the small difference in the levels

of residual stress in materials B and C and the large dif­

ference i n the plastic deformation to which they have been

subjected, it seems that the latter of the abovementioned

effectsis the most important.

A further feature of the curves for materials A

and C, being most noticeable in the case of the former,

is that in the early stages the effective cross-section

decreases at a steady rate, which however, become s smaller

in ti me. This is bel ieved t o be a functi on of the direction

of propagation of the cracks as discussed in Section 5.6.

In the early stages a crack propagates in a plane perpendicular

to the tensile stress axis and is r epresented by the straight

-92-

line portions of the graph. Eventually, the crack branches,

each branch advancing at an angle to the original crack-

ing plane. This means that even if the crack advances at

the same rate it ~d.ll have less effect in decreasing the

remaining cross-sectional area and the slope of the curve

becomes smal ler. The branches change direction progressively

until they are propagating in a plane that is parallel to

the surface of the strip, at which stage they have virtually

no effect on the remaining cross-sectional area of the

sample. Hence the curves in E'ie;.l3, at sufficiently high

values of time of exposure should become straight lines

parallel to the time axis.

As far as can be judged the curve representing

the behaviour of material B, after the induction period, is

a straight line representing a steady rate of decrease of

effective cross-sectional area. This, no doubt, is due to

the fact that even at the maximum time of exposure cracks in

the material were not suff iciently àeveloped to branch.

The curves reveal that the initial rate of crack­

ing is great est in material A and least in material B. The

high initial rate of material A is at least partially

ex~lained by the comparativel y high r es idual stress l evel

existing therein. The reason why material C, which was

shown in Sec ti on 5. 4 to have a lovrer residual stress th an B,

should exhibit a higher rate of crac king t han this mat erial

is more difficult to understand.

-93-

The first possibility is that the method of stress

measurement used, one that is admittedly not very accurate,

gave incorrect values of residual stress for these two

materials. Alternatively, an explanation is possible in

terms of the theory of Forty (30) as discussed in Section

3.4. According to this theory susceptibility to cracking

is progressively reduced by increased cold-work until an

optimum value is reached. Beyond this value susceptibility

may inc reas e a gain due to the hardening of slip bands arrl

the consequent favourable conditions for the propagation

of brittle cracks.

-94-

5.8 Results of Loop Tests

These tests were performed in accordance with

the procedure outlined in Section 4.4. The only materia1

subjected to this test was C, the specimens being taken

perpendicu1ar to the direction of rolling.

Many of the specimens tested c racked in two or

more stages, each stage of cracking being accompanied by

a distinct metallic click. The total times to reach the

various stages of cracking for a number of specimens are

summarized in Table 13. In each case the last time recorded

represents complete failure of the sample.

TABLE 13

Time to Failure Tes t No.

lst Stage 2nd Stage

A2 6 min. A7 9.!_ 15 min • . AS ti A9 10 A10 All 21 29 Al2 g 12

It may be seen from the tabulated results that the

times for cracking to be completed were by no means con­

sistent. For this t est t o be used f or comparative purpo s es

it would be neces sary t o use a large numbe r of specimens

-95-

and apply statistical methods.

However, it was found that improved agreement

between cracking times could be obtained if the position

of the test loop in the corrosion vessel was standardized.

Consequently, a number of hooks of standard length were

prepared so that the specimens were always suspended at the

same depth in the battery jar. Results obtai ned when

position had been standardized are recorded in Table 14:

TABLE 14

Time to Failure Test No.

lst Stage 2nd Stage 3rd Stage

Al5 il! min. 14 min. Al6 12 16i Al7 10 18; 21~ min. Al8 il 25

The results indicate that tmre is greater agree-

ment between the times for cracking to complete the first

stage than was apparent with the results of Table 13. It

seems that by standardizing the position of the test speci­

mens and recording the time at which the first stage is

reached, then a measure of comparative susceptibility may

be obtained with fewer tests.

-96-

Presumably, the reason that position is so critical

is that ~ atmosphere composition, particularly with respect

to oxygen and carbon dioxide, will vary from the bot tom of

the jar to the top.

-97-

6. SUMMARY

The literature relating to the stress-corrosion

cracking of brasses has been surveyed arrl the theory of

su ch cracking dis cussed. A study has be ru ma de of the

cracking behaviour of three different gauges of cartridge

brass strip.

The ultimate tensile strengths of all three

materials were d etermined both parallel and perpendicular

to the direction of rolling. It was fourrl that in all

three materials the tensile strength varied with the

direction in which it was measured, being greatest perpen­

dicular to the rolling direction. The differential increased

with increase of cold-rolling. This variation of strength

"'ith direction has been reported for sheet ma.terials by

oth er wor kers.

Longitudinal sections and sections parallel to

the plane of rolling were prepared from each material and

the microstructure examined. As exp e cted, e ach structure

was markedly defdrmed and traces of pronounced slip bands

were visible. Evid ence of the e xistenc e of annealing twins,

resulting from the previous history of the strip, was

present in each material.

-98-

Specimens of the test materials, taken both

parallel and perpendicular to the direction of rolling

were subjected to pre-corrosion tests. For each material,

the specimen taken parallel to tre direction of rolling

suffered the greater loss in strength. This was found to

be due to the fact tha t cracking was trans verse in these

specimens, thus indi ca ting tha t tre operati ve residu al

tensile stresses in the stri p were orient ed in the rolling

direction.

Values of approximated residual stress at the

surface -..vere àetermined for each gauge of strip. The

highest value of residual stress was fo und t o exist in

the least heavily reduced material, whilst of tte other

two materials, tre more he a vi ly re duc ed was indic ated as

having a slightly lower stress level. The tensile stresses

at the surface were to be expected in thi n strip reduced

by large dianeter rolls. Hcw ever, the r elative values of

stress exi sti ng in the respective mat erials c an. only be

explained in terms of th e rolling schedule to which they

had bee n subjected , since it has be en shov.rn t hat th e

ultima te residual stress pattern depends largely upon the

final pass.

The effect of stress-relief annealing at 250°C

for 30 minutes on the ultimate tensile strength of the1east

-99-

heavily reduced material was determined. The strength was

found to increase both parallel and perpendicular to the

rolling dire ct ion, the effect bei ng most pronounced in

the latter case. The increase in strength of cold-worked

brass upon law-temperature annealing has 'œen widely reported

but never satisfactorily explained.

Specimens of the same ma teri al, stress-relief

annealed at the same temperature and for the same t ime, were

subjected to pre-corrosion tests. Tpe stress-relieved

material exhibited a decided decrease in susceptibility

to stress-corrosion cracking.

Longitudinal sections of specimens of all three

materials which had been exposed to corrosive conditions

and had cracked, were polished and etched. Also, the

rolled surface of a sample of the least heavily reduced

material was polished, etched and exposed to the corrosive

environment, after which it was re-polis hed and re-etched.

Examination of these sections revealed that cracking vlas

basically transgranular, although occasionally a crack

followed a suitably oriented grain boundary for a short

distance. This confirmed the view held by others that

plastic deformation of alpha brasses res ults in a tendency

for cracking t o be corœ transgranular.

-100-

Inspection of the sections also indicated that,

in its early stages, a crack tended to propagate i::1 a

plane that was perpendicular to the stress axis. After

proceeding sorne distance in this manner the crack branched.

The branches so f ormed progressi vely changed direction

until they were advancing in a plane that was parallel to

the specimen surface. It was presumed that changes in

direction of crack propagation resulted from alteraUons

iQ residual stress distribution caused oy the cracking.

The sections also shov1ed tha t what appeared to

be areas of localized corrosion \vere associ at ed wi th the

cracks. Although in many cases no visible connection

existed between the se areas and the crack, there is a possi­

bility that they were connected in planes other than the one

examined.

The susceptibility of the strip materials was

compared by subjecting specimens, all eut parallel to the

direction of rolling, to pre-corrosion tests for varying

periods. Graphs relati ~ effective decrease in cross­

sectional area to time of exposure were prepared for each

type af strip. The curves revealed that an induction

period, the length of which varied for each material, was

necessary before the onset of cracking. A possible explana­

tion of the variation in length of the i nduction peri od,

-101-

in terms of residual stress and deforma ti on structure, was

proposed.

The shape of the curves was explained as being

a function of tre direction of propagation of the cracks,

a matter which has been discussed previously.

The graphs also revealed that tœ initial rate

of crack propagation, as indicated by the original slopes

of the curves, was different for each material. An explana­

tion was proposed in terms of the theory discussed in

Section 3. 4.

The thinnest gauge cartridge brass strip was

employed to evalua te the comrnonly us ed loop test as a means

of comparing stress-corrosion cracking susceptibility of

copper alloys. It was shawn th at pro vi ding the position of

the loop in the corrosion vessel was standardized the number

of tests needed to determine the tendency tD f'ail was com­

paratively small.

-102-

APPENDIX I

The Hounsfield Tensometer

The Hounsfield Tensometer is a portable, laboratory­

scale machine for determining the mechanical properties of

materials, both metallic and non-metallic. It rœ.y be used

for tensile, compression, not ched-bar, hardness and bend

tes ting. For the pur poses of the programme undertaken the

Tensometer was used only as a tensile machine.

The tensile test pie ce is supported horizont ail y

by attachment at one end to a spring b eam and at the otrer

to a cross-head. The load is applied t o the cross-head

by hand via a worm gear and causes the spring bearn to deflect.

The deflection is proportional to the load and is measured

by the movement of a mercury c olumn. A recorder dru rn with

graph paper attached is driven through sui table geari ng

so that the circ~~erential movement of the drum is propor­

tional to the strain imparted to the specimen. Thus by

following the meniscus of the mercury column with a pricker

as the load is applied, a replica of the stress-strain

diagram may be obtained on the graph paper.

A frequent source of ina cc uracy, when t ensile­

testing strip or wire, is the tendency of the specimen to

slip within or break at the chucks. By waisting the

-103-

test pieces and using the special "Quick Grip Chucks"

these troubles were avoided. However, care must be

taken when using these accessories as under full load there

is a possibility th at halls may escape from the moving ball

race which is an integral part of each chuck.

A comparison has be en made between data obtai ned

with a Hounsfield Tensorœ ter and wi th a standard Riehle

Screw Power Uni versal Tes ting Machine ( 69). Values of

ultimate tensile strength for SAE 1035 and SAE 2330 steels

were measured using both machines. There was little

difference between the two values of ultimate tensile strength

obtained for either material. In the course of the sane

investigation the properties of alumin~m alloy sheet were

determined using the 1'ensometer. The ultimate tensile

strengths obtained for Alcan 2-SO and Alcan 57-SH differed

by less than 1.25 percent from the values specifi ed by the

manufacturer.

-104-

APPENDIX II

Residual Stress Measurement

Residual stresses cannot be measured directly

in the manner that applied stresses are measured. Residual

stresses are, in fact, calculated indirectly from the strains

that exist within internally stressed material {70). These

strains are usually measured by mechanical or X-ray methods

and the corresponding stresses calculated by applying

elastic theory formulae.

Exact analysis of the residual stress in a part

is a long and tedious procedure. Consequently, several time

saving shorter methods of analysis have been developed.

These are widely used even though the results obtained may

be somewhat distorted {67). The stresses calculated by

these methods are frequently named "approximated residual

stressesn, a term which will be used hereon.

Many cornmonly used approximation methods of

stress measurement are incl uded L1 the category known as

ndeflection techniques". These methods may be applied

where stresses are believed to vary linearly through the

thickness of a plate or tube wall, but are constant along

the lengt h, across the width of plate or around the circum­

ference of the tube, due to the conti nuous nature of the

-105-

fabrication process. The techniques involve mechanical

slitting of the material followed by measurement of the

deflection of the slit length.

One such approximation methcxi may be applied to

the measurement of residual stresses in rolled strip and

is known as the "Modified Anderson and Fahlman" technique.

The test assumes that the distribution of stress through

the thickness of the strip is linear, and is performed

as follm"ls. A sanple of the strip is partially slit along

its central plane. The two halves of tm strip curl back

releasing whatever bending moment existed in them prior

to splitting. This bending moment, M, according to the

elastic theory of the bending of beams, is given by the

following expression:

M = E1I ••••••••••••••••••••• (4) (>

where El = E

l -)A- 2

E = Young's modulus of elast i city

/- = Poisson's ratio

I = Moment of inertia of split section

f = Radius of curvature.

Making the assumption th at distribution of stress

through the t hickness of the s trip is linear, the maximum

longitudinal stress at the surface, s1 , is given by the

equation:

MC r

-106-

•••••••••••••••••••••••••••• ( 5)

where C = distance from the neutral axis to the outer fibre.

The distance, C, in this case is gi ven as follows:

c = t •••••••••••••••••••••••••••• ( 6) 4

where t = thickness of strip before slitting.

Thus, substituting this value for C in Eq.(5):

= Mt 4I

•••••••••••••••••••••••••••• ( 7)

From Eqs. ( 4) and ( 7):

•••••••••••••••••••••••••••• ( g)

The radius of curvature rray be expressed

in terms of tœ deflection and the length of tœ slit sec-

tion of the strip, where the deflection is very small in

comparison to the radius of curvature, by the following

expression:

= •••••••••••••••••••••••••••• ( 9)

where f = deflection, i.e. dista~ce between the curled-back

ends of the strip.

l = length of slit section.

-107-

From Eqs. (8) and (9):

= • • • • • • • • • • • • • • • • • • • • • • ( 10)

Due to the small gau ge of the cartridge brass

strip used far this work the Modified Anderson and Fahlman

technique could not be used. However, a method of measur­

ing approximated residual stresses in thin strip, based on

the Anderson and Fahlman technique, was developed.

The method involved machining a portion of each

strip dawn to half thickness and measuring the deflection

at the centre of the remaining part. The maximwn longi­

tudinal stress at the surface, sl, is given by Eq.{8):

= ••••••••••••••••••••••••• ( 8)

Now as the def1ection at t he centre is very

sma1l compared with the radius of c urvature, the follol."'ing

expression applies :

••••••••••••••••••••••••• (11)

where y = def1ection at centre of ma chined strip

1 = 1ength of machined strip

From Eqs. (8) and (11):

= ••••••••••••••••••••••• ( 12)

-108-

A value of approxina ted residual stress was obtained

for each material by measuring y with a micrometer gauge and

applying Eq.(l2). As a check on the accuracy of measure-

ment the radius of curvature of each machined piece of strip

was obtained by construction, and the. value inserted in

Eq. (8). Thus two values of s1 were obtained for each material.

The method may be criticised on several gro unds,

such criticism applying also to the modified Anderson and

Fahlman technique. The stress distribution through the strip

is assumed to be linear and it is unlikely that this is so • .

Also, in the course of machining or splitting the sanple,

beat will be generated and this will have a stress-relief

effect. Sorne distortion of the material in close proximity

to the eut or machined surface will occur, and this, undoubt­

edly, will affect the final result. Hovvever, the technique

is regarded as givine sufficiently accurate results so that

sorne comparison between the residual stresses of strip

mat erials may be made.

-109-

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-110-

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-111-

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-112-

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