stress-corrosion susceptibility and the dislocation arrangements of austenitic stainless steels

Corrosion Science, 1965, Vol. 5, pp. 71 to 79. Pergamon Press Ltd. Printed in Great Britain

STRESS-CORROSION SUSCEPTIBILITY A N D THE

DISLOCATION ARRANGEMENTS OF AUSTENITIC

STAINLESS STEELS*

K . C. THOMAS,* R . STICKLER t a n d R . J. ALLIO*

*Atomic Power Division and tResearch Laboratories, Westinghouse Electric Corporation, Pittsburgh, Pa., U.S.A.

Abstract--The dislocation arrangements in 304 stainless steel and Incoloy 800 have been determined as a function of temperature and strain. Dislocation pile-ups were observed in both alloys in spito of the fact that the stacking fault energies were considerably different, viz. < 16 erg/cm 2 for the 304 stainless steel and > 30 erg/cm 2 for the Incoloy 800.

The stress-corrosion behaviour of the two alloys in magnesium chloride boiling at 142°C was determined. The Incoloy 800 was found to be considerably more resistant than the 304 stainless steel.

After critically reviewing the current theories of stress corrosion and structural and stress-corrosion data for a series of austenitic alloys, it was concluded that the data could best be correlated with the theory proposed by Swarm. According to this theory, the presence of short range order is expected to markedly influence dislocation arrangement and stress-corrosion behaviour. A review of the litera- ture revealed that there is some indirect evidence to indicate the presence of short range order in certain of the austenitic alloys, but the evidence is not conclusive.

R~sum6--La dislocation de la disposition dans l'acier inoxydable 304 et l 'Incoloy a 6t6 d6termin6e comme une fonction de la temp6rature et de la tension. Des dislocations d'empilements ont ~t6 observ6es dans ces deuz alliages rnalgr6 le fait que les 6nergies de d6faut d'empilement ~taient tr/~s diff6rentes ( < 16 erg/cm 2 pour l'acier inoxydable et > 30 erg/cm ~ pour l 'Incoloy 800).

Le comportement de la corrosion sous tension pour les deux alliages dans le chlorure de magn6sium bouiUant/t 142°C a 6t6 d6termin6. On a trouv6 que l'Incoloy 800 6tait beaucoup plus r~sistant que l'acier inoxydable 304.

Apr~s une revue critique des th6ories courantes sur la corrosion sous tension et des donn~es sur la corrosion structurelle et la corrosion sous tension pour une s6rie d'alliages aust6nitiques, on conclut que les donn~es correspondaient le mieux/~ la th6orie propos6e par Swarm. D'apr~s cette th6orie, l'existence d 'un arrangement b. courte distance doit influencer de fa~;on marquee la dislocation des arrangements et le comportement de la corrosion sous tension. Une revue de la litt~rature a r6v616 quelques preuves indirectes indiquant I'existence d'arrangements ~t courte distance dans certains des alliages aust6nitiques, mais les preuves ne sont pas concluantes.

Zusammenfassung--Die Versetzungsanordnung in rostfreiem Stahl Typ 304 und Incoloy 800 wurde in Abh~ngigkeit von Temperatur und Dehnung untersucht. Versetzungsanh~iufungen wurden in beiden Legierungen beobachtet, obwohl die Stapelfehler-Energie der beiden Legierungen deutlich unterschiedlich ist und 16 erg/cm 2 fiir den rostfreien Stahl und mehr als 30 erg/cm 2 fiir Incoloy 800 betr~gt.

Die Spannungskorrosions-Best~ndigkeit der beiden Legierungen wurde in Magnesiumchlorid- 16sung mit einem Siedepunkt yon 142°C untersucht. Incoloy 800 erwies sich als wesentlich besfiindiger als der rosffreie Stahl Typ 304.

Ein kritischer Oberblick fiber die derzeitigen Theorien der Spannungsrisskorrosion und der strukturellen und korrosionschemischen Eigenschaften einer Reihe austenitischer Legierungen ergibt, dass die eigenen Ergebnisse mit der Theorie yon Swann am besten iibereinstimmen. Nach dieser Theorie soll eine Nahordnung die Versetzungsanordnung und das Spannungskorrosions-Verhalten von Legierungen deutlich beeinflussen. Die Durchsicht der Literatur ergibt einige indirekte Hinweise auf das Vorliegen von Nahordnung in gewissen austenitisehe Legierungen. Diese Hinweise sind jedoch nicht ganz eindeutig.

*Manuscript received 11 May 1964.

71

INTRODUCTION

IN Tim last few years attempts have been made rather successfully to explain susceptibility to stress-corrosion cracking in terms of the dislocation arrangements in an alloy after small amounts of deformation. 1-4 The results from some of the earlier work suggest the following general principles. In a material with a low stacking fault energy (SFE) cross slip of extended dislocations is difficult and planar arrays of dislocations are formed; ~-~ alloys exhibiting this type of dislocation arrangement are highly susceptible to stress-corrosion failure. Conversely in an alloy with a high SFE cross slip is relatively easy and cellular dislocation arrays are obtained. Such materials are resistant to stress-corrosion failure. ~-4 The type of dislocation arrangement is markedly dependent on the SFE, which determines the equilibrium separation of partial dislocations in the face-centred-cubic lattice. Thus, before a dislocation can cross-slip, its partials must recombine. However, in an alloy of low SFE the widely extended dislocations make cross slip difficult; dislocations remain piled up on the slip planes to form a planar array. The ease of cross slip in an ahoy of high SFE prevents dislocation pile ups, resulting in a cellular arrangement of dislocations. 5

REVIEW OF PREVIOUS WORK

More recent work, however, indicates that the correlation between dislocation arrangement and stress corrosion behaviour is not as rigid as had been previously believed. In fact, it now appears that factors other than the SFE may markedly influence the dislocation arrangement and also the stress-corrosion behaviour. Some of the recent observations will briefly be reviewed. Douglass et aL e have shown that Nichrome, which is highly resistant to stress-corrosion cracking in chloride solutions, has a planar arrangement of dislocations. In addition, Incoloy 800, also a stress-corrosion resistant austenitic alloy, has a planar dislocation arrangement. In the case of Nichrome the SFE was stated to be ,-~ 16 erg/cm 2, whereas that of Incoloy 800* is high, > 30 erg/cm 2. Swarm 4 has shown that phosphorus additions to a 2 0 ~ Cr/20 ~ Ni austenitic steel results in planar array of dislocations and a high susceptibility to stress-corrosion failure, while the SFE remains at a value that is too high to measure, viz. > 30 erg/cm ~. Similarly, a quenched CuaAu alloy has a SFE in excess of 20 erg/cm ~, and a planar dislocation arrangement. Usikov et a lY showed that at low deformations ( < 8 per cent) binary nickel alloys contain planar groups of dislocations, whereas at higher deformations ( > 10 per cent) a cellular arrangement is observed. In one of the Ni-Cr alloys at low deformation ( ~ 3 per cent) the presence of planar groups of dislocation pairs .was noted. These dislocation pairs were attributed to the presence of long range order (LRO) in the solid solution. The effect of temperature of straining a N i - 4 ~ Ti alloy was also investigated by Usikov et al. ~ It was found that straining at 600°C produced no piled up groups of dislocations, but only narrow dislocation loops; the same alloy strained a similar amount at room temperature contained planar arrays of dislocations.

These observations suggest that the magnitude of the strain and the temperature at which the strain is applied may control the type of dislocation arrangement developed. It is thus evident that studies aimed at correlating stress-corrosion behaviour with dislocation arrangement should consider the effect of total strain and temperature

*See Table 3 for chemical composition.

Stress-corrosion susceptibility and dislocation arrangements of austenitic stainless steels 73

of strain. Stress-corrosion testing of austenitic stainless steels is generally carried out at temperatures of 140-155°C and at stresses corresponding to instantaneous strains of about 5 per cent and less. However, the dislocation arrangements in most studies have been determined after strains of 5 per cent or more at room temperature. In order to obtain a more valid correlation between dislocation arrangement and stress- corrosion behaviour, the arrangement should be determined under the conditions experienced during stress-corrosion testing.

Furthermore, the SFE may not be the sole factor that determines whether a planar or cellular arrangement of dislocations is obtained. Swann, 8 on the basis of work on CuaAu and austenitic stainless steels, suggests that the presence of short range order (SRO) markedly influences the type of dislocation arrangement developed in many alloys. SRO will confine slip in clusters on well-separated slip planes since the order destroyed by the motion of a dislocation across one plane will lower the stress to move others over the same plane. 8 He has further supposed that slip forms chemically reactive sites for crack initiation at the surface, and re-ordering destroys these sites. This is a dynamic process, dependent on stress and temperature; the sites attacked are those in which deformation is in progress and not old sites which may have been destroyed by re-ordering. The importance of the planar dislocation group is that it can provide a localized site for preferential electrochemical attack which possibly results in the formation of a nucleus for a brittle crack. In the case of a planar group resulting from a low SFE the high energy slip plane will be anodic with respect to the low energy matrix, and thus a crack nucleus may be formed. On the other hand a piled-up group resulting from the presence of SRO will have a local ctiemical composition change across the slip plane after slip and thereby cause an electrochemical effect. If re-ordering takes place at a rapid enough rate, then the net chemical composition change will be zero and the tendency to failure will be diminished.

In Table I are assembled dislocation arrangement and stress-corrosion data for various face-centred-cubic alloys. An interpretation of the stress-corrosion phenomenon in terms of dislocation arrangement must consist of two parts. First the dislocation arrangement must be described in terms of the factors that influence it, and secondly the arrangements have to be correlated with the stress-corrosion data.

According to present ideas the two important factors that influence the dislocation arrangement are (a) SFE and (b) SRO. It is assumed that SRO exists in the Fe-Ni-Cr system and that the degree of SRO and rate of ordering are dependent on alloy content. The effect of SFE and SRO on the dislocation arrangement is as follows. A low SFE is conducive to the formation of planar dislocation arrays as a result of a restriction of cross-slip. The presence of SRO supposedly tends to confine slip to well-separated slip planes, again promoting planar type arrangements. 4,s In the low- nickel alloy (18/8) the degree of SRO is expected to be low, the SFE is low, and therefore planar groups of dislocations form. In the 20 ~o Cr/20 ~o Ni stainless steel the SFE is high and the degree of SRO is still small, which means that there will be no tendency to form planar groups in this alloy. On addition of 0-3 Yo P the degree of SRO is expected to be increased, slip is restricted to well-separated planes, and a planar arrangement is formed. The SFE is still high which accounts for the absence of extended dislocations. The SFE may even be increased by the addition of phosphorus since it has been noted that the twinning frequency, which is related to the SFE, of a

74 K. C. THOMAS, R. STICr, Lt~ and R. J. ALtao

TABLE 1

Susceptibility Dislocation Stacking Short range Rate of to stress-

Alloy arrangement fault energy order re-ordering corrosion failure in MgCls

304 SS Planar Low Low Low High (18 Cr/8 Ni) (2, 3, 4)* (2, 3, 4)*

16/20 Cellular High Low Low Low (16 Cr/20 Ni) (2, 3) (2, 3) (2, 3)

20/20 Cellular High Low Low Low (20 Cr/20 Ni) (4) (4) (4)

20120/P Planar High High Low High (20 Cr/20 Ni]0'3 P) (4) (4) (4)

Incoloy 800 Planar High High High Low (20 Cr]30 Ni) * *

20/40 Planar High High High Low (20 Cr/40 Ni) (6) (6) (6)

Nichrome Planar Low High High Low (80 Ni/20 Cr) (6) (6) (6)

Nickel Cellular High Low Low Low (63 (6) (6)

*Present investigation.

316L stainless steel is decreased by the addition of phosphorus. 9 In the case of the 20 ~o Cr/30 Yo Ni alloy there is no tendency to form planar groups due to the high SFE, but the degree of SRO is now expected to be high enough to concentrate slip on a few widely separated planes, resulting in a planar dislocation array. A similar argument holds for the Ni--4 ~o Ti alloy.

The contention that alloys exhibiting planar dislocation arrays are susceptible to stress-corrosion failure and those exhibiting cellular arrays are not susceptible, must be modified to account for the resistance to chloride stress-corrosion failure of higher- nickel alloys which contain planar dislocation arrangements rather than cellular. According to Swann's argument, if planar groups form and re-ordering can take place at a fast enough rate to destroy the chemically reactive sites, then the net effect will be a high resistance to stress-corrosion failure. 4 Such a condition would hold in the higher-nickel alloys if it is assumed that increase in nickel favours an increase in the rate of re-ordering. In the low-nickel alloy, 304, the reactive sites are not destroyed since SRO does not exist in the first place; therefore, the susceptibility to failure will be high.

More recently Douglass, Thomas and Roser s have discussed the effect of SRO on the stress-corrosion phenomenon in austenitic alloys. They maintain that if SRO is strong, co-planar groups of dislocations form which enhance susceptibility to failure. This is in direct contrast to Swann who states that the presence of SRO promotes

Stress-corrosion susceptibility and dislocation arrangements of austonitic stainless steels 75

susceptibility only if re-ordering cannot take place at a rate in excess of disordering by plastic deformation. It should be added that the observed high resistance of a planar grouping at a high SFE (e.g. Incoloy 800) cannot be explained by the hypothesis due to Douglass et aL but can by Swann's model.

There is no direct experimental evidence for the existence of SRO in the Fe--Cr-Ni system; however, based on the following indirect evidence, mainly for the Fe-Ni system, the assumption is not unreasonable. Vidoz, Lazarevic and Cahn 1° cite the calorometric data of Iida 11,x~ and the creep data of Kornilov xa as indirect evidence for the development of SRO in Fe-Ni alloys. The maximum in the SRO occurred at a composition of about 75 ~o Ni/25 ~o Fe. It is possible that the addition of chromium would move this to lower nickel contents. In agreement with Bradley and Taylor 1~ and Wakelin and Yates, ~5 Vidoz et al. 1° noted that a ternary addition of aluminium gave a considerably greater rate of ordering. Substantial persistence of SRO well above the critical temperature has been detected by neutron diffraction by Collins, Jones and Lowde. xe These authors definitely showed that even after quenching from 1000°C there exists a considerable amount of SRO in a 60 ~ Fe/40 ~o Ni alloy.

In the case of the Ni-Cr system, Roberts and Swalin 17 did not detect any evidence for ordering by neutron diffraction. Recent Russian work, ~8 however, has shown that for a similar alloy, 30~o Cr, superlattice lines are obtained if single crystals are used instead of powder samples. Previously, resistivity and specific heat data have shown that some effect occurs on heating the Ni/30 ~o Cr alloy at about 500°C, resulting in the so-called K-state. The exact nature of this effect is not clear, but ' i t has been suggested as being due to either SRO, LRO or clustering.

Summarizing, therefore, there is strong evidence that SRO exists in the Fe-Ni system, and indications that it may be present in the Ni-Cr system, although the exact type of ordering cannot be specified. It is probable that ordering can take place in the Fe-Cr-Ni system, although this remains to be substantiated.

Although in the previous discussion the main emphasis has been placed on the effect of structure on the stress-corrosion susceptibility, other factors may be of equal importance. For example, the environment is extremely selective in determining stress- corrosion behaviour. In the case of stainless steels it has been definitely established that chloride ions induce failure, although the hydroxyl ion may also be effective, x9 In addition, it has been recently shown that the susceptibility to stress-corrosion failure, as measured by the time to failure, also varies with the type of cation associated with the chloride ion. z° In Table 2, stress-corrosion data are shown for 304 stainless steel in chloride solutions under conditions of constant stress, temperature, pH and chloride ion concentration; the only variable is the cation present in solution. It will be noted that the susceptibility to failure decreases in the order Mg ~+, Ca 2+, Li + under the given conditions.

The present investigation was undertaken to study the effect of strain and temperature on the dislocation arrangements at strains and temperatures close to those at which stress-corrosion tests are carried out. The two alloys investigated were 304 stainless steel, which is susceptible to stress-corrosion, and Incoloy 800, which is extremely resistant to failure by stress-corrosion in boiling aqueous magnesium chloride solution at 142°C.

76 K.C. THOMAS, R. S~CKLm~ and R. J. ALtaO

TABLE 2. Tttcm TO FAILURE OF 304 STAINLESS STEEL IN AQUEOUS CHLORIDE SOLUTIONS. TEST CONDITIONS: 60,000 lb/in s, 27 PER CENT CHLORIDE ION CONCENTRATION; 125°C; pH

(AT 80°C) "q- 3"0

Chloride solution Time to failure, h

MgCI2 6 CaCla 160 LiC1 300

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

The chemical compos i t ion o f the al loys used in the invest igat ion, in the fo rm of wire 0.035 in dia., is shown in Table 3 a long with the mechanica l proper t ies . Specimens were cut and tested "as received", fully softened but slightly cold worked due to coil ing and uncoil ing. The specimens were degreased by swabbing with acetone pr io r to test.

Samples were s t rained in an Ins t ron tensile machine for var ious amount s a t 30 and 142°C. Af ter s training at the high t empera tu re the load was removed and the specimen was quickly cooled to r o o m tempera tu re by means of an air blast . The wire samples were th inned by a technique which has been previously described, 21 and examined by t ransmiss ion in an electron microscope opera t ing at 100 kVA.

The s t ress-corrosion tests were carr ied out in reflux type glass cells, d i rect tensile stresses were appl ied to the sample by means o f a pul ley system. This equ ipment has previously been described. 2°.~2

TABLE 3

Alloy (wt. ~ ) 304 Incoloy 800

Cr 19.02 20'41 Ni 8"55 33-56 Mn 1 '23 0"82 Si 0-54 0'53 Nb 0.021 Me 0.30 0-08 Ti < 0.005 C 0.06 0"6 N 0.012 .0045 B 0.0049 S 0"011 0"007 P 0"029

0"2 ~o proof stress (lb/in 2) 41,600 42,800 Ultimate tensile

strength (lb/in 2) 93,750 104,600 ASTM grain size 7-8 7-8


RESULTS AND DISCUSSION

The dislocation arrangements of 304 stainless steel as a function of strain and temperature are shown in Fig. 1. Those for the Incoloy 800 are shown in Fig. 2.

The effect of increase in strain is to increase the dislocation density; for a given strain an increase in temperature results in an increase in dislo.cation density. In- creasing the temperature of straining from 30 to 142°C had little other effect, on the dislocation arrangement in either alloy.

In both alloys the dislocations have a planar arrangement as shown in Figs. 1 and 2. Extended and contracted dislocation nodes are visible in the 304 stainless steel as shown in Fig. 3. By comparison with previous measurements a on node radii the SFE is estimated to be low, viz. about 16 erg/cm 2. On the other hand no extended or contracted nodes are visible in the Incoloy 800 alloy, as illustrated in Fig. 3(b). Both these are indicative of a high SFE, certainly not less than about 30 erg/cmL

On the basis of the SFE estimates it must be concluded that the planar grouping observed in the Incoloy 800 alloy is caused by some effect other than the SFE. A similar conclusion was reached by Swann 4 to account for the planar grouping in a 20 ~o Cr/20 ~o Ni/0.3 ~o P stainless steel.

The stress-corrosion data for the two alloys are shown in Table 4. Although both alloys have planar dislocation arrangements, the times to failure in boiling magnesium chloride solution are widely different. Failures are obtained in 304 stainless steels at stresses of 5000 Ib/in ~ and above in times of 15 h and less, whereas failures were obtained in Incoloy 800 at a stress level of 60,000 lb/in ~ and above.

The dislocation arrangement and corrosion behaviour of both alloys may be explained by assuming that SRO exists in the Incoloy 800. In the case of the 304 stainless steel the planar dislocation grouping results from the low SFE. These dislocations remain piled up on their slip planes after plastic deformation and the emergence of these pile ups at a surface provides chemically reactive sites for crack initiation. In Incoloy 800 on the other hand, the planar dislocation arrangement is caused by the confinement of slip to well-separated planes by SRO in the lattice. 4,s,2z,24 However, in this case re-ordering is expected to take place rapidly at the test temperature resulting in the destruction of the chemically reactive sites and thus the alloy is resistant to failure.

As mentioned previously, it has not been established whether SRO does in fact exist in the Fe -Cr -Ni austenites. In electron diffraction studies no diffraction spots due to superstructure were detected, nor was domain contrast observed in bright

TABLE 4

Time to failure, in h, at stresses of* Alloy 60,000 40,000 30,000 5000

lb/in 2 lb/in" lb/in 2 lb/in 2

304SS

Incoloy 800

0-9 1.1 15

280 350

*Average of ton tests per stress level.

78 K.C. THOMAS, R. STICKLmt and R. J. ALLIO

field and dark field transmission electron microscopy. These would indicate an absence o f L R O ; however, SRO may be present.

Ordering is the location o f certain elements on preferred lattice sites. In the case o f the ternary alloys o f Fe -Cr -Ni , which are further complicated by the presence o f other elements such as carbon and nitrogen, it is not known which element or~upies these preferential sites. Considerably more work will have to be undertaken to establish whether SRO does in fact exist, and to detei'mine the atomic species responsible for the effect. Other effects, such as clustering or precipitation, should also be considered since they create regions of inhomogeneity which are a prerequisite for crack initiation. However, since it is fairly certain that the required inhomogeneities are on an atomic scale this would rule out gross precipitation effects.

CONCLUSIONS

1. Increase in strain at a given temperature, and increase in temperature at a give strain, increase the dislocation density o f 304 stainless steel and Incoloy 800.

2. There is no appreciable difference in dislocation arrangement, determined at rt temperature, on increasing the temperature of strain for both alloys.

3. Both alloys exhibit planar arrays of dislocations, but the SFE varies f rom a . ." value of N 16 erg/cm 2 for the 304 stainless steel to a high value of ~ 30 er~_ for the Incoloy 800. The planar array o f dislocations in the Incoloy 800 is expl,~ ._-, on the basis of SRO in the lattice.

4. The time to failure o f 304 stainless steel, in boiling magnesium chloride solution, at 40,000 lb/in ~ is less than one hour, whereas failure does not occur in Incoloy 800 in less than 350 h under similar conditions. These results are significant since both alloys exhibit planar dislocation arrangements and both would be expected to 1-- highly susceptible to failure on the basis o f previous theories. It is suggested, as previously proposed by Swann, a that the resistance to failure o f Incoloy 800is due to the re-ordering reaction taking place at a sufficiently rapid rate at the test temperature to destroy the chemically reactive sites which are required for crack initiation.

R E F E R E N C E S 1. P. R. SWANN and J. NUTTING, J. Inst. Metals 88, 478 (1960). 2. R. STICKLER and S. BARNARTT, J. Electrochem. Soc. 109, 343 (1962). 3. S. BARNARTT, R. STICKLER and D. VAN ROOYEN, Corros. Sci. 3, 9 (1963). 4. P. R. SWANN, Corrosion 19, 102t (1963). 5. P. R. SWANN and J. NtrrriNG, J. Inst. Metals 90, 133 (1961). 6. D. L. DOUGLASS, G. THOMAS and W. R. ROSER, Corrosion 20, 15t (1964). 7. M. P. USIKOV and L. M. UTEVSKn, Phys. Metals Metallogr. 13, 61 (1953). 8. A. H. COTTRELL, Relation of Properties to Microstructure. American Society of Metals, Cleveland

(1954). 9. K. C. THOMAS and E. B. WE~LEIN, Trans. Quart. Araer. Soc. Metals 56, 566 (1963).

10. A. E. Vrooz, D. P. LAZAREVIC and R. W. CAHN, Acta Met. 11, 17 (1963). II. S. IIDA, 3". Phys. Soc. Japan 7, 373 (1952). 12. S. IIDA, J. Phys. Soc. Japan 9, 346 (1954). 13. I. I. KORNILOV, Mechanical Properties oflntermetallic Compounds. Wiloy, New York (1960). 14. A. J. BRADLEY and A. TAYLOR, Proc. Roy. Soc. A166, 353 (1938). 15. R. J. WAKEUN and E. L. YATES, Proe. Phys. Soc. Lond. B66, 221 (1953). 16. M. F. COLLINS, R. V. JONES and R. D. LOWDE, J. Phys. Soc. Japan 17, Suppl. B-III, 19 (1962). 17. B. W. ROBERTS and R. S. SWALt•, Trans. Amer. lnst. Mech. Engrs 209, 845 (1957).

- 1

,

(b)

Cc) (4

FIG. 1. (see over)

(4

FIG. 1 (CONT’D). 304 stainless steel.(a) 03 % strain. (b) 2.9 % strain.(c) 4.9% strain. (d) 7.4 % strain.(e) 0.5 strain. ( f ) 26 % strain.

(g) 5.6 % strain.

(4 Cd)

FIG. 2. (see over)

W

FIG. 2 (CONT'D). Incoloy 800.(a) 0.5 % strain. (b) 2.0 y0 strain.(c) 5.0 % strain. (d) 8 % strain.(e) 0.5 % strain. ( f ) 3.0% strain.(g) 4.6 % strain. (h) 8.2% strain.

FIG. 3.(a) 304 stainless steel, strained 1.0% at 30°C. Extended dislocation nodes; low

stacking fault energy.(b) Incoloy 800, strained 0.5% at 30°C. Unextended dislocation nodes; high

stacking fault energy.


18. V. I. GOMANKOV, D. F. LrrvIN, A. A. LOSHMANOV and B. G. LYASHCHENKO, Phys. Metals Metallogr. 14, 133 (1962).

19. P. P. SNOWDEN, J. Iron Steel Inst. 194, Pt. 2, 181 (1960). 20. K. C. THOMAS, H. M. F ~ A ~ a and R. J. ALLIO, Corrosion 20, 88 (1964). 21. R. STICKLER and R. J. ENGLE, Microjet Method of Preparation of Wire-Samples for Transmission

Electron Microscopy. Westinghouse Electric Corporation Scientific Paper 63~148-546-p4 (1963). 22. S. BARNARTT and D. VAN ROOYEN, 3". Electrochem. Soc. 108, 222 (1961). 23. J. B. COHEN and M. E. Fn~rE, J. Phys. Radium 23, 794 (1962). 24. V. I. Sx'trno~A and E. S. YAKOVLEVA, Phys. Metals Metallogr. 14, 92 (1962).

stress-corrosion susceptibility and the dislocation arrangements of austenitic stainless steels

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