The mechanism of stress-corrosion cracking in the brass-ammonia system

Corrosion Science, 1968, Vol. 8, Pp. 55 to 65. Pergamon Press. Printed in Great Britain

THE MECHANISM OF STRESS--CORROSION CRACKING IN THE BRASS-AMMONIA SYSTEM* MICHINORI TAKANO

and SABURO SHIMODAIRA

The Research Institute for Iron, Steel and Other Metals, Tohoku University, Sendai, Japan A b s t r a e t ~ T h e effect of corrosive environments, compositions of alloying elements and degree of pre-strain on the stress-corrosion cracking of the brass-ammonia system, has been investigated and the mechanism of intergranular cracking and of the transition from intergranular to transgranular cracking have been discussed. Intergranular cracking arises from the formation of thick and large cuprous oxide grains over slip steps and the resulting locking of the movement of dislocations. The transition from intergranular to transgranular cracking depends on the mechanical properties of the surface film produced in a corrosive environment. The mechanism of the stress-corrosion cracking has been explained on the basis of the film theory. R~sumA---On a ~tudi6 l'effet exercd sur la corrosion sous tension du syst~me laiton-ammoniaque par l'environment corrosif, la composition d'gl6ments d'alliages et le degrg de raise en tension; on a discut6 le mgcanisme de la corrosion intergranulaire et de la transition de la corrosion intergranulaire transgranulaire. Une corrosion intergranulaire rdsulte de la formation de grains dpais et grands d'oxyde cuivreux sur des bandes de glissement, et de l'obstacle qui en rgsulte sur le mouvement de dislocations. La transition de la corrosion de la forme intergranulaire ~ la forme transgranulaire dgpend des proprigtds mdcaniques du film superficiel produit darts un milieu corrodant. Le mdcanisme de la corrosion sous tension a gt6 expliqud sur la base de la th$orie du film. Zusammenfassung--Der Einflul3 der korrodierenden L6sung, der Zusammensetzung der Legierung und der angewendeten mechanischen Spannung auf die Spannungskorrosion yon Messing in Ammoniak wurde untersucht, und der Mechanismus des interkristallinen Bruchs und des Crbergangs zu transkristallinem Bruch wird diskutiert. Der interkristalline Bruch entsteht, wenn sich dicke und grobl3 Kupferoxydk6rner auf Gleitstufen bilden und damit die Bewegung von Versetzungen blockieren. Der CJbergang zu transkristallinem Bruch wird beeinflul3t durch die mechanischen Eigenschaften des gebildeten Oberfl~chenfilmes. Der Ablauf der Spannungsril3korrosion wird auf der Grundlage der Deckschicht-Theorie erkliirt. Pe~epaT - - I4cc~eaoBaHo B~HAH~e Koppos~OnHbIX cpe~, n p a p o a ~ ~erHpymm~x O~eMeHTOB n cTeneHH npe~BapHTe~bHoro HanpA~eaHa Ha HopposHOHHOe pacTpec~4BaH~e no~ Hanpah~enHeM CHCTeM JIaTyHL-aMMHaK; o 6 c y ~ e H MexaHH3M Meh~KpHcTa~THOrO pacTpecKHBaHHa nepexo~ OWMe~RpHcTa~J~HTHOrO K TpaHc~pHcTa~HTHOMy pacTpecKtmaHHm. Me~RpHcTa~HTHOe pacTpecK~naHHe np0Hcxo~HT I~3-3a oSpa3oBa~Ma n~0vH,-~x H 5 0 ~ m H X 3epeH ORHCH Me/In no nJIOCI~OCTI:IM CHOJIb;t~eHHH, qTO B~3HBaeT TOpMOmeHHe ~BHH~eHH~

IIepexo~ oT Me~i~p~tc~aJ~J~;~THOrO g TpaHc~pHcTa~THOMy p a c T p e c ~ B a ~ m HBJIneTCH ~yHK~He~ OV MexaHaqecsHx CBOI~CTBno~epxHOCTHOt~ n~eH~H, oSpaayeMo~ B goppos~OHHOfZ cpe~Ie. Mexa~HaM ~opp03~OHHOrO pacTpeCgHBaH~a ~0~ HanpameH~eM o6~ac~aeTc~ Ha OCHO~e n~eHOqHOi~ weopHm *Manuscript received 17 February 1967; received in revised form 12 July 1967. 55

IVhci-nNolu T A g ~ o and SABUROSHIMODAIRA

56

INTRODUCTION

THE MECHANISMof stress-corrosion cracking has been widely studied and discussed by many investigators, but many factors are involved in the cracking process and the influence of these factors is very complicated and no well-defined elucidation of the mechanism has been given heretofore. It is well known that in the brass-ammonia system the form of the cracking, intergranular or transgranular, depends on the corrosive environment,1-4 composition of alloying elements 5-7 and degree of work. s,9 Therefore, this system is of particular interest in relation to the mechanism of stress-corrosion cracking. It is our view that the existence of the characteristic black oxide layer, the tarnish, produced in a corrosive environment is essential to intergranular stress-corrosion cracking in this system. It has been emphasized by previous electrochemical studies that cupric complex ions Cu(NHs)~'+ play a major role in intergranular cracking.l.a.4,1°, u However, the fact that the micro-corrosion behaviour of alloys deformed plastically under stress cannot be directly determined makes it impossible to clearly define the mechanism of cracking. The film rupture theory which ascribes the mechanism to the formation of a surface film in a corrosive environment and its subsequent breakdown under stress has been developed by Logan, x2 Forty and Humble 1~ and McEvily and Bond, ~4 but none of them have fully explained the processes involved from the initiation of micro-stresscorrosion to the propagation of cracks. The present work was carried out to investigate the effects of corrosive environments, compositions of alloying elements and degree of pre-strain on stress-corrosion cracking in the brass-ammonia system, using thin foil or bulk specimens. The mechanism of intergranular cracking and of the transition from intergranular to transgranular cracking have been discussed from the standpoint of crystal structure. Following the previous work, x5 this paper is intended to describe at full length the mechanism of intergranular stress-corrosion cracking of a-brass, with emphasis on the film theory. EXPERIMENTAL

The chemical compositions of the brasses are listed in Table 1. The thin foil specimens (cold-rolled to 0.1 mm) were I00 × 5 × 0-1 mm and were wrapped in thin foils of the same composition and annealed at 600°C in vacuum. They were then electrolytically polished in 200 cm s of phosphoric acid + 80 g of chromic acid at 90°C. The deformation of the electropolished thin foil specimens in various corrosive environments was performed in the same way as that described previously.Is TABLE 1.

CHEMICAL COMPOSITIONS OF "IF.ST ALLOYS

Test alloys

Cu

Zn

Pb

F¢

10Zn-Cu brass 20Zn-Cu brass 30Zn-Cu brass 40Zn-Cu brass 0.86Al-brass 1.97Al-brass

90.51 79.58 70.02 60.18 79.04 78.54

R R R R R R

0-004 0.007 0.006 0-005

0.012 0.005 0.003 0.003

AI

0"86 1'97

The mechanismof stress--corrosioncrackingin the brass-ammoniasystem

57

Specimens for the stress--corrosion cracking tests, and for the observation of the stress--corrosion behaviour by the replica technique, were 60 × 7 × 1.2 ram. They were packed in carbon powder to prevent dezincification, annealed for 1 h at 600°C, and were then stressed by the three-point bend-beam system with a constant deformation. Tests were carried out in 0.05 M CuSO4 q- 0.5 M (NH4)zSO4, which is thesame composition as that used by Mattsson, t and also in ammonia vapour. The pH value of the solution was adjusted to 2.0, 7-4 or I0.0 by addition of sulphuric acid or ammonia. In order to give the pre-strain to the specimens an Instron tensile test machine (Type T'I'-CM-L) was used and each specimen was deformed by 2, 5, 10 or 15 per cent in tension, respectively. In order to investigate how the initial micro-attacked trenches extend at crystallographic defects, stress-corrosion tests of the annealed bulk specimens by the three-point bend-beam system were carried out, and the corrosion behaviour of the same specimens exposed for various times in the corrodent was examined by the replica technique. All the specimens used were annealed and electropolished prior to the experiment. The average grain size was about 5 i~m and the electropolished specimen surface prior to the experiment did not show any irregularity under various magnifications. At 5, 30, 60 and 120 min after the start of stresscorrosion tests, each specimen was removed temporarily from a given corrosive environment and the corrosion behaviour of the specimen surface was examined. RESULTS

Stress-corrosion cracking tests Tests for both annealed and pre-strained bulk specimens of each brass were carried out in Mattsson's solutions and in ammonia vapour, and the corrosion products on the specimen surface after the test were examined by electron diffraction. Generally, with increasing pre-strain, the failure time decreased and the number of cracks increased, and the specimens which suffered intergranular cracking were covered with a black oxide layer. The results of stress-corrosion tests are summarized in Table 2. 30% Zn-Cu alloy in Mattsson's solution at pH 2-0 and 10.0 did not fail within the test periods of 7 and 5 days, respectively. 10Yo Zn-Cu alloy and two kin:Is of Al-brass also did not fail within 12 days. With the 40% Zn-Cu alloys pre-strained above 5 per cent, both intergranular and transgranular cracks were observed. In the other cracked specimens, however, the difference in the form of cracking with the degree of pre-strain was not recognized, but only intergranular cracking was observed. All the specimens tested in ammonia vapour failed rapidly; in particular two kinds of Al-brass failed within 2 h. The difference in the form of cracking with the degree of pre-strain was not observed, but intergranular and transgranular cracks were observed in 10--30 % Zn-Cu alloys and in 40 Yo Zn-Cu alloy respectively, and both intergranular and transgranular cracks were observed in two kinds of Al-brass: intergranular cracking was predominant in 0"86~o Al-brass, and transgranular cracking in 1.97 % M-brass. The surface films after the test consisted mainly of cuprous oxide CuzO and were independent of the corrosive environments, susceptibility and form of cracking. Stress-corrosion behaviour of thin foil specimens From the results obtained by the stress-corrosion test, it is evident that the

58

MICHINORI TAKANO a n d SABURO SHIMODAIRA TABLE 2.

RESULTS OF STRESS--CORROSION CRACKING TESTS OF BULK SPECIMENS

Mattsson's Solution Specimens Annealed

Pre-strained "

Ammonia vapour Surface film

NC

Annealed

Prestrained

Surface film

I

I

Cu~O, Cu

10~. Zn---Cu

pH 7.4 NC

20~. Zn--Cu

pH 7"4 I

I

CutO, Cu

I

I

Cu20, Cu

30 ~. Zn-Cu

pH 2-0 NC pH 7"4 I pH 10"0 NC

I

Cu~O, Cu

I

I

Cu~O

40~0 Zn-Cu

pH 7"4 I

I for2~o I + T for 5~o

Cu20

T

T

Cu~O

pH 7.4 NC

NC

Cu20, Cu

I -t- T

I q- T

Cu20, Cu

pH 7.4 NC

NC

Cu~O, Cu

I + T

I + T

Cu20, Cu

0.86 AI-

brass 1.97 AI-

brass

NC: No cracking. I: Intergranular cracking. T: Transgranular cracking. susceptibility and the form of cracking depend on the corrosive environments, degree of pre-stain and composition. In order to examine these differences microscopically, the stress-corrosion behaviour of the thin foil specimens was observed by transmission electron microscopy. The annealed materials used as the thin foil specimens were deformed in various corrosive environments. The micro-stress-corrosion behaviour of 30 ~o Z n - C u alloy in Mattsson's solution at p H 2-0, 7.4 or I0.0 and of 40% Z n - C u alloy in the same solution at p H 7.4 is shown in Fig. 1. These micrographs show the micro-corrosion trenches nucleated along the slip steps which were produced by dislocation movements on the restricted slip planes under stress. In addition, preferential dissolution was also observed at the dislocation ends, twin boundaries and grain boundaries, but predominantly at slip steps. The above experiment was performed under the condition in which intergranular cracking occurred easily in the solution at p H 7.4 and also that in which the cracking occurred with difficulty in the solution at p H 2.0 or I0.0. Although there was no difference in the mode of initiation of micro-attacks, the nucleated micro-attacks in the former condition tended to extend immediately over the brass surface. The typical examples of the surface-extended corrosion traces in 30 ~o Z n - C u alloy and 40 ~o Z n - C u alloy at p H 7.4 are shown in Fig. 2. The selected area diffraction pattern obtained from the preferential corroded region in 30~o Z n - C u alloy is shown in Fig. 3. In this ease the corrosion products in the preferentially attacked region consisted of Cu and Cu20, and the oxide developed epitaxially on the dezincified copper metal. Zn and ZnO were not detected in the corroded thin foil specimen, and two straight lines which can be considered as the slip traces produced in the original brass specimen are visible in the extended films of Cu20 and Cu.

The mechanismof stress--corrosioncracking in the brass-ammoniasystem

59

As stated already, both intergranular cracking and transgranular cracking took place in 30 ~o Zn-Cu and 40 ~o Zn-Cu alloys in ammonia vapour, respectively. The stress--corrosion behaviour of these thin foil brass specimens in this corrodent are shown in Fig. 4. In the two specimens there is no difference in the initiation of microattacks as shown in Figs. 4 (a) and (b), and only the slip steps are preferentially corroded. In 30 ~o Zn-Cu alloy which cracked intergranularly in this corrodent, the formation of the surface-extended corrosion as seen in Mattsson's solution at pH 7.4 (see Fig. 2) was not observed in 3 min of stress-corrosion. This difference of microcorrosion behaviour in the Mattsson's solution at pH 7.4 and in ammonia vapour will be discussed later.

Stress-corrosion behaviour of bulk spechnens The characteristics of stress--corrosion cracking of the bulk specimens depended remarkably on the corrosive environments and compositions of alloying elements as seen in Table 2. The nucleation of the micro-stress-corrosion trenches of the thin foil specimens was independent of the susceptibility and the form of cracking, and took place predominantly at the disclocation ends, twin boundaries and grain boundaries, particularly at fresh slip steps. In the experiment of the thin foil specimens, however, only the initial stage of micro-attacks could be revealed. The earlier reported transmission micrographs (Figs. I and 2) were prepared from the annealed specimens which plastically deformed, and in this section the corrosion behaviour of the annealed bulk specimens which plastically deformed by the three-point bend-beam system was examined by the replica technique. Figure 5 shows the stress-corrosion behaviour of 30~o Zn-Cu alloy in Mattsson's solution at pH 2.0, 7.4 and I0.0. Figure 5 (a) and (b) are the micrographs obtained after 5 and 120 min, respectively. Figure 5 (a) shows the following characteristics: (a) at pH 2.0 corrosion is so slow that the corrosion trenches along the slip traces are seen, (b) at pH 7.4 the corrosion morphology extends over the bulk surface corresponding to the surface-extended micro-attacks as shown in Fig. 2 (a), and (c) at pH 10-0 the corrosion rate is very high, and the original surface topography of the bulk specimen has already disappeared, followed by the occurrence of general corrosion. The following characterislics can be seen from Fig. 5 (b): (a) at pH 2.0 the surface corrosion morphology in Fig. 5 (a) still remains, suggesting that the stresscorrosion cracking of brass in this solution is difficult to occur, (b) at pH 7.4 the corrosion morphology within the grains as shown in Fig. 5 (a) does not appear, but the grain boundaries are significantly corroded, indicating the appearance of intergranular cracking, (c) at pH 10.0 corrosion is much more severe than that after the 5 min period of exposure. The stress-corrosion behaviour of 30 ~o Zn- and 40 ~o Zn-Cu alloys and 1.97 YoAIbrass in ammonia vapour are shown in Figs. 6 (a), (b) and (c), respectively. Figure 6 represents the corrosion morphology of the specimen surface obtained after 5 and 120 min (upper 5 min, lower 120 min). After 5 min, corrosion along the slip traces is seen in 30 ~o Zn-Cu alloy, but the tendency toward general corrosion has already been seen in 40~o Zn-Cu alloy and in 1-97Yo Al-brass. After 120 min, corrosion proceeds significantly at grain boundaries in 30~o Zn-Cu alloy, suggesting the initiation of intergranular cracking. However, the morphology of corrosion in

60

Mxcr~mOR1TAKANOand SASUROSHIIdODAIRA

40 ~o Zn-Cu alloy and in 1"97~o Al-brass shows extremely irregular roughness, and the nucleation of micro-fracture in the grains of 40 ~o Zn-Cu alloy is observed. DISCUSSION In the present experiments the 30 ~o Zn-Cu alloy in Mattsson's solution at pH 2-0 and 10.0 did not fail. In particular, in the solution at pH 10.0 the specimens were severely corroded and their thickness was rapidly decreased by corrosion, so stressing could not be applied to the specimens. Mattsson1 has recently reported the electrochemical reactions of 37~o Zn-Cu alloy in the solution and found transgranular cracking in the solution having high pH values. The observed corrosion behaviour in the solution at each pH value in this experiment, generally, agree with his results. The micro-attacked trenches were very evident along the slip steps in the grains as shown in Figs. 1 and 4. The occurrence of these micro-trenches by a combined action of stress and corrosion suggests the nucleation of stress corrosion. Similar micro-corrosion behaviour, however, was observed even in the case of high stresscorrosion cracking resistance?9 The nucleation and propagation of stress-corrosion, therefore, cannot be elucidated according to the same mechanism. Previously, the authors reported x9that in austenitic stainless steels the mechanisms of nucleation amd propagation of stress-corrosion cracking were different from each other. Tromans and Nutting2° have reported that the preferential attacks of the thin foil 30~o Zn-Cu alloy in Mattsson's solution (pH 7.3) occur at grain boundaries and at individual dislocations adjacent to grain boundaries, and these micro-attacks are associated directly with intergranular micro-cracking. RSnnquist21 has also demonstrated that the micro-attacks occur at grain boundaries under a small stress and at slip steps in the grain under a large stress. In this experiment, however, the nucleation of stress-corrosion was not associated with the susceptibility and the form of cracking, and took place remarkably at the fresh slip steps. This fact agrees well with the results of the previous report? 9 The surface film consisted mainly of cuprous oxide in all cases, being independent of the susceptibility and the form of cracking. Whether the nucleated micro-corrosion trenches continuously proceed to cracking or cease to extend will be determined by the competition between passivation due to the action of H~O or dissolved oxygen and activation due to the breakdown of the film by the action of ammonia or other anions, rather than by the chemical compositions of surface films. Forty and Humble, x3 Pickering and Swarm, 1~and Hoar and Bookeris reported that the corrosion product on a brass surface in various ammonia environments consisted of small platelet cuprous oxide. In our work, under the condition which induced intergranular cracking, the cuprous oxide developed epitaxially on the dezincified copper metal and existed in the form of grains as large as ,,- 1 lain (Figs. 2 and 3). The oxide, which nucleated on the slip steps, would act as the cathode to the brass substrate and extend over the brass surface by the electrochemical reaction. For 30 ~o Zn-Cu alloy intergranularly cracked in ammonia vapour, the micro-attacked trenches which occurred on the slip steps at the initial stage did not extend so fast as that in Mattsson's solution, because at the initial stage the electrochemical reaction in vapour could not proceed as fast as that in aqueous solution. This agrees with the result of the stress--corrosion behaviour of the bulk specimens examined by the replica technique. Comparing Fig. 5 (Mattsson's solution, pH 7.4)

FIG. 1. Preferential micro-attacks obtained in deformed thin foil specinaens in Mattsson's solution for 3 min. (a) 30 Zn-Cu brass (pR 2.0). (b) 30 Zn-Cu brass (pFI 7.4). (c) 30 Zn--Cu brass (pH 10.0). (d) 40 Zn-Cu brass (pFI 7-4).

FIG. 2.

Typical examples of surface extended corrosion obtained in deformed thin foil specimens in Mattsson's solution (pH 7"4) for 3 min. (a) 30 Zn-Cu brass, (b) 40 Zn-Cu brass.

II (~i 3) Cu

. (ooz)C~ 20

"~o

-II .Cu

e#

lilt { )c.zo

4'

ms

10d (M

g o

FIG. 3. Selected area diffraction pattern from preferential corroded region obtained in deformed thin foil 30 Z n - C u brass in Mattsson's solution (pH 7.4) for 3 min.

FIG. 4.

Preferential micro-attacks at slip steps obtained in deformed thin foil specimens in ammonia vapour. (a) 30 Zn-Cu brass, (b) 40 Z n - C u brass.

FiG. 5.

Stress-corrosion behaviour of 30 Z n - C u brass in Mattsson's solution at pH values of 2.0, 7.4 and 100. Ca) Exposure time 5 min, (b) Exposure time 120 rain.

FIo. 6.

Stress~orrosion behaviour of (a) 30 Z n - C u brass, (b) 40 Z n - C u brass, and (c) 1.97 AI-brass in ammonia vapour. Upper 5 rain. Lower 120 min.

The mechanism of stress--corrosion cracking in the brass-ammonia system TABL~ 3.

RELATIONBzrw~E~ THE PROPERTY

OF SURFACE FILM AND THE FORM OF CRACKING

(~ 37 Zn-Cu Brass. CuSO4 + (NH4)~SO4 (0-05 g/l Cu + 1M NFIa) pH 4 4"7 5"2- 5"5 6.3- 6.6 6'7- 7'3 9.4-11.2

Surface Brown-red Brown-red Interference colour Black Black Dark stain

Crack T I, T T I, T I T

Reference 1. (ii) 30 Zn-Cu Brass. Ammonia Solution Solution IM or 5M NH4OH + CuSO4 1 or 5M NI-I4OH + CuCl~

I or 5M NH4OH + Cu(NO~)~

Surface

Crack

Visible film Visually bare Visible film Visually bare Visually bare Filmed surface

I T T T T N C

Reference 2. (iii) 30 Zn--Cu Brass. 15N Ammonia Solution + Cu, pH > 13 Cu Content (g/l)

Surface

Crack

0.5 0.5-1-5 1.5-3.25 3-25

No visible film Brown film, loosely No visible film Tarnish

N C I, T I, T I

Reference 3. (iv) 30 Zn--Cu Brass. 0.04M Cu + I'5M Total Ammonia pH

Surface

Crack

5 6--7.5 8

Copper c31our Blue-black tarnish Blue-black tarnish

T, I I I

Reference 4. (v) 30 Zn-Cu Brass. 5 %, 25 ~o Nt-140H + 0-5 y. (NH,),CrO, Surface

Crack

Grey translucent Interference colour No visible corrosion

T T T

Reference 22.

61

62

MICmNORI TAKANOand SABUROSmMODAIRA

and Fig. 6 (ammonia vapour) for 30 YoZn-Cu alloy, a remarkable difference in stresscorrosion behaviour can be seen after 5 min of the stress--corrosion test, i.e. the former showed the surface-extended corrosion and the later corrosion traces along the slip lines. However, the corrosion remarkably penetrated into grain boundaries in both cases after 120 min, suggesting the occurrence of intergranular cracking. When transgranular cracking occurred, the nucleated micro-trenches on the slip steps neither extended over the specimen surface nor formed the thick Cu20 oxide of large grain size as seen in Fig. 7 (Mattsson observed transgranular cracking at pH 10.0) and Fig. 4 (b), and the appearance of the specimen surface was not smooth, but showed a topography of general corrosion as seen in Figs. 5 and 6. Recently, there are many investigations1-4. 22 about the relation between the properties of the surface films and the form of cracking. Those results are summarized in Table 3. Excepting the results obtained in the solution containing CI- ions, Table 3 indicates the occurrence of transgranular cracking when the surface films in the ammonia environment are not thick. From the experiments and considerations mentioned above, it is apparent that the mechanical properties of surface films (grain size, thickness, porosity, hardness and strength, etc.) play an important role in the process of stress--corrosion cracking. From the viewpoint of the role of the surface films, the following hypotheses are proposed for the mechanism of intergranular stress-corrosion cracking of Cu-Zn alloys in the ammonia environment. (1) The fresh coarse slip steps produced by the movement of dislocations under stress in the corrodent are preferentially corroded resulting in the micro-attacked trenches; these micro-attacked regions consist of the surface films of Cu20-Cu layers (Fig. 7 (a)). (2) Cu~O-Cu films developed on the slip steps act as the cathode to the brass substrate and accelerate the electrochemical reaction. As the specimen surface is very rapidly corroded, the Cu20-Cu films extend immediately over the brass surface. These surface-extended films develop epitaxially to the brass substrate on the original slip steps, forming thick and large grains, so that the movement of dislocations on the slip plane below those films is blocked. On the other hand, the other slip steps could be newly formed in the uncorroded region, and these new slip steps would exhibit the same behaviour as shown in Fig. 7 (a) (Fig. 7 (b)). (3) The brass surface is covered with thick Cu20-Cu films of large grain size. At this stage the slip steps in the grains cannot develop, because the movement of the dislocations on the emergent slip planes are almost blocked by the surface films. On the other hand, dislocations, vacancies and impurities are present in grain boundaries, and unde stress the grain boundaries are slipped by the movement of dislocations and are disturbed by the interaction of the dislocations and vacancies with the impurities. Furthermore, the surface films on the grain boundaries involve many defects and the diffusion of the corrodent to the tip of micro-crevices, so that the chemical reaction in the grain boundaries are intensified and the local corrosion is accelerated. The path of stress--corrosion cracking thus changes from the slip steps in the grain to the grain boundaries (Fig. 7 (c)). As the surface film produced in the corrodents responsible for transgranular cracking was not so thick nor so large in grain size as that in the case of intergranular

The mechanism of stress--corrosion cracking in the brass-ammonia system

Grain boundary

/

Amonia aqueous

Slip step

63

Grain boundary

solution

(a) Brass

(b)

(c)

FIG. 7.

Schematic representation of intergranular stress-corrosion cracking of brass in ammonia aqueous solution. (a) Showing the preferential micro-attack at slip step.

(b) Micro-attack begins to extend over the brass surface and the movement of dislocations at coarse slip plane is locked by Cu20 and Cu layers. (e) After brass surface is covered with these layers, crack propagates along grain boundaries. cracking, such a film would not so strongly block the development of the slip steps created by the movement of the dislocations. Therefore, the nucleated micro-attacked trenches along the slip steps within the grains continuously grew up into transgranular cracking. In this case, the surface films are ruptured by the increased shear stress associated with the development of slip steps as shown by the surface topography of the micro-fracture of 40 ~o Z n - C u alloy (Fig. 6). Whether the alloys in which the dislocations are distributed in a planar arrangement fail by intergranular or transgranular cracking could be determined by the relative values of the amount of shear stress at the tip of the dislocation pile-up and the strength of the surface films formed in a corrosive environment; that is, if the n dislocations pile up against the brass surface, the amount of shear stress acting on the surface films is equal to n~', where ~ is the shear stress. Now denoting the strength of the surface films as F, the development of slip steps would be obstructed and intergranular cracking occurs for F > n~r; for F < nr, transgranular cracking may be preferred. Whether or not the surface films produced in the corrodent are ruptured by the development of steps would depend on the thickness, porosity, grain size, degree

64

MICI-ffNORITAKANO

and SABURO SHIMODAIRA

of pre-strain of brass substrate and the coherence of surface films to the brass substrate, etc. These factors would also depend on the corrodent, composition of alloying elements and amounts of stress. Transgranular cracking thus propagates by the corrosion of the fresh slip steps developed under plastic deformation, and the surface films produced are mechanically ruptured by the shear stress at the growing steps, facilitating the diffusion of the corrodent to the tip of microcracks. The mechanism of stress-corrosion cracking in other homogeneous alloys may be explained according to the above-mentioned film theory. Very few of the numerous micro-trenches nucleated on the slip steps in the early stage grow up into cracks. This fact can be explained as follows: the stress-corrosion cracking occurs in the corrosive environment where the influence of passivators is stronger than that of activators, and therefore repassivation in the slip steps usually prevails as corrosion progresses. Furthermore, the stress field in the alloys is microscopically non-uniform. Under these conditions, the corrosion rates of the microtrenches into the interior of the alloy would become different from one another with the lapse of time. At the tip of the most deeply corroded regions on the slip planes, a microscopic stress concentration occurs, and the movement of dislocations on these slip planes is further activated and the fresh slip planes are easily exposed to the corrodent. Passivators such as dissolved oxygen would also find difficulty in diffusing into the micro-crevices. The corrosion of these slip planes, therefore, is accelerated and cracks are formed. Then the corrosion would penetrate not only along the same slip planes but also along the adjacent slip planes, branching off at the planes of crossslip. The segregation of solute atoms (impurity atoms) should also play an important role in the penetration of the corrosion. The above explanation is concerned with the transgranular cracking when the applied stress is large enough to rupture the surface films. Even if the applied stress is not sufficient to rupture the surface film, a crack may develop at the base of a corrosion pit since here slip is presumably not hindered by the presence of a surface film. There appears to be a close relationship between pitting and stress-corrosion.

CONCLUSIONS From the results of these experiments, the following conclusions can be obtained. (1) The susceptibility and the form of stress-corrosion cracking in the brassammonia system depend on the nature of the corrosive environment, the composition of alloying elements and the degree of pre-strain. On the other hand, the nucleation of the micro-stress corrosion trenches is independent of the susceptibility and the form of cracking, and take place predominantly at the slip steps in all cases. (2) Under the condition of susceptibility to intergranular cracking, thick cuprous oxide of large grain size is formed over the slip steps. This surface film obstructs the development of steps in the grains. In the grain boundaries the dislocations and vacancies move considerably under stress, resulting in the formation of slipped regions, and therefore the grain boundaries become chemically active and corrosion penetrates along these boundaries. (3) When the strength of the surface films is not so high as to obstruct the developmerit of steps created by the movement of the dislocations, the films are ruptured by

The mechanism of stress-corrosion cl'acking in the brass-ammonia system

65

shear stress and the corrosion along the slip planes proceeds, resulting in the occurrence o f transgranular cracking. (4) The transition from intergranular to transgranular cracking depends on the mechanical properties o f the surface films produced in a corrosive environment, (5) W h e n the applied stress is not so large as to rupture the surface films, the pits nucleate and the node regions having no surface films are formed at the b o t t o m o f the pits, where slip occurs easily, forming the nucleus in the initiation o f stress-corrosion cracking. (6) Transgranular cracking propagates by the corrosion o f fresh slip steps produced at the tip o f the micro-fracture. The segregation o f solute atoms (impurity atoms) to the slip planes which form the chemically active path plays an important role. The mechanism o f stress--corrosion cracking o f ordinary homogeneous alloys can be explained according to the above-mentioned film theory. REFERENCES 1. E. MATrSSON,Electrochimica Acta 3, 279 (1961). 2. W. Lw~.s, Corrosion 21, 125 (1965). 3. E. N. Puon and A. R. C. W~STWOOD,Phil. Mag. 13, 167 (1966). 4. H. E. JOHNSONand J. L~A, Corrosion 22, 178 (1966). 5. M. E. WmTAr~R, Metallurgia Nov.-Dec., 21, 66 (1948). 6. A. K. LAnn~ and T. BAN~E, Corros. ScL 5, 731 (1965). 7. P. R. SwAn, Corrosion 19, 102 (1965). 8. W. D. ROBERTSONand A. S. TETELEMAN,Strengthening Mechanism in Solid, American Society of Metals, Novelty, Ohio, p. 217 (1960). 9. Y. MURAr,.AMX,H. YOSHIDAand Y. IKAI,J. Japan Inst. Metals 1215 (1965). 10. F. C. At.THOr,Z. Metallk. 36, 177 (1944). 11. T. A. READ, J. B. REED and H. ROSENTrr.~, Symposium on the Stress Corrosion Cracking o f Metals, p. 90, ASTM (1944). 12. H. L. LOGAN,J. Res. natn. Bur. Stand. 48, 99 (1952). 13. A. J. FORTYand P. HUMBLE,Phil. Mag. 8, 247 (1963). 14. A. J. McEVILYand A. P. BOND,J. electrochem. Soc. 112, 131 (1965). 15. M. TAKANOand S. SmMOD~RA, Trans. Japan bzst. Metals 7, 193 (1966). 16. M. TAKANOand S. SHIMOD~dRA,J. Japan Inst. Metals 29, 553 (1965). 17. H. W. PICIr.EmNGand P. R. SWANN,Corrosion 19, 373 (1963). 18. T. P. HOARand C. J. L. BOOKER,Corros. Sci. 5, 821 (1965). 19. M. TAKANOand S. ShUMOD~RA,Trans. Japan Inst. Metals 7, 186 (1966), 20. D. TRO~.NS and J. NtrrnNG, Corrosion 21, 143 (1965). 21. A. R6NNQtnST, 3rd Int. Congr. metallic Corros. p. 151, Moscow (1966). 22. W. LYNES, Corrosion 22, 113 (1966).