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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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-111-
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-112-
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