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The fractography of stress corrosion cracking in carbon steels
Corrosion Science, 1975. Vol. 15. pp. 469 to 477. Pergamon Press. Printed in Great Britain
THE FRACTOGRAPHY OF STRESS CORROSION CRACKING IN CARBON STEELS* B. POULSON Materials Research, A.T.D., Clarke Chapman Ltd., Gateshead, England Abstract--The stress corrosion cracking of carbon steels in a number of environments (nitrates, hydroxides, carbonates, chlorides and liquid ammonia) has been studied by scanning electron microscopy. The crack path depends on both heat treatment and environmental conditions. In annealed steels it appears that solutions which produce thick corrosion products induce intergranular cracking, while in solutions where adsorption induced inhibition occurs or thin films form, cracking is transgranular. Although the available fractographic evidence is ambiguous, it appears to rule out either hydrogen embrittlement or brittle films as being important in crack propagation. However, the fracture surface morphologies are consistent with the cracks propagating by anodic dissolution, possibly aided by mechanical tearing. INTRODUCTION FRACTOGRAPHY is an important experimental technique in b o t h failure analysis and mechanistic studies. This is because the unambiguous interpretation o f fracture surface m o r p h o l o g y would allow how the crack propagated to be determined. N u m e r o u s papers have dealt with fractographic aspects o f stress corrosion cracking (SCC) in: stainless steels, 1.2 aluminium alloys, a,4 titanium alloys, 5, 6 magnesium alloys, 7 high strength steels, a,9 and copper base alloys. 1°,xl However, carbon steels have been somewhat neglected. 9 This is possibly due to the mistaken belief that mild steels only crack intergranularly. 12,1a Furthermore, m a n y o f the solutions which cause cracking produce thick corrosion products and it is assumed 9 that such layers will obscure any fractographic detail. The purpose of this paper is to summarize some fractographic observations obtained in recent investigations 14,~5,~a o f SCC in carbon steels. EXPERIMENTAL The SCC testing technique has been described in detail elsewhere. 14 Briefly constant strain rate (2.2 × 10-6/s) tests were performed on unnotched cylindrical tensile specimens; either electropolished or finished to grade 600 emery paper. The compositions and heat treatments of the steels examined are shown in Tables 1 and 2. The SCC tests were carried out in a number of solutions over a wide range of potentials (Table 3). Immediately failure occurred specimens were removed from the cell, cleaned in methanol and stored until they were examined in the SEM. RESULTS The potential ranges over which SCC occurred for certain steels in the different solutions are shown in Fig. 1. The results for the liquid a m m o n i a having been converted to the normal hydrogen scale (assuming p N H ~ = 16 and p H + ~ p N H ) + they are therefore somewhat approximate. Elsewhere in this paper they are quoted as they were measured, i.e. against a [~P d - H electrode in liquid a m m o n i a . These potential ranges p r o m o t i n g SCC are in general agreement with previous work. M o s t o f the fracture surface o f specimens removed f r o m the environment directly they failed was rather clean. A n exception was the failures in nitrate solutions at potentials ( - - 200 to *Manuscript received 8 July 1974. 469
]], POULSON
470 TABLE 1. C
Si
Mn
S
P
Ni
AI
0.09 0.09 0"075 0.08 0.08 0' 1 0" 14
0.06 0.09 0.06 0.09 0.04 0.22 0.36
0.82 0.48 0.19 0.75 0-49 1 1"73
0"01 0.02 0.04 0.03 0.04 0-05 0.013
0.02 0.01 0.03 0.02 0.04 0.06 0.01
0.04 1.05 2.15 4'57 6.07 0.02 --
0.02 0'03 0.06 0'05 0.02 0.05
Alloy NiO Nil Ni2 Ni3 Ni4 Mo4 SIS 2103
COMPOSITION OF STEELS USED
TABLE 2. As received Annealed Quenched H i g h T e m p . Anneal.
TABLE 3.
Others
5 /o,o M o
--
0-I ,%Cu
HEAT TREATMENTS USED (N) (A) (Q) (HTA)
i.e. normalized 950°C for 1 h, furnace cooled 950°C for 1 h, water quenched 1150°C for 3 h, furnace cooled
ENVIRONMENTAL CONDITIONS INVESTIGATED
Temperature
Solution
Nitrate, 4N N H 4 NO3 Chloride, approx. 45 ~ M g C I 2 Hydroxide, 35 ,%NaOH Carbonate, 1 : 1 I N Na~ COs: 1N N a H C O s A m m o n i a , liquid a m m o n i a
pH
Potential range m V
Boiling Boiling at 154°C Boiling
4
-- 600 to d 1200 S.C.E. - 1000 to - 150 S.C.E. -- 1300 to + 450 S.C.E.
95°C Ambient
9 (at 25°C)
-- 1100 to -- 500 S.C.E. -- 1000 to + 2000 P d H
1500 I
OOO
I000 500
Liquid
Nitrate
IiJ m
ammonia
u')
0
._o
E
E
-500
13_
Data for
-
NIO(A) in nitrate in chloride
-tO00 --Ni4(A) NiO(A)
-
FIG. 1.
150C
t
~
i
500 n
I
in c a r b o n o t e
SIS2103(Q} in hquid ammonia ~ ~ NiO(A) in hydroxide ~ydrogen ernbritl'lernent
f
-
-
I000
Schematic d i a g r a m s h o w i n g potential ranges over which S C C o f carbon steels occurs in various solutions.
q- 100 mV S.C.E.) where the general corrosion rate was high, and failures (Fig. 2) in hydroxide at potentials where oxygen evolution occurred. Because of certain similarities the results of tests in nitrate, carbonate and hydroxide solutions are presented first. This is followed by the results of tests in liquid ammonia and chloride
, '
,4
~
~,~.~,~7.."~
3
"
~
~:?!
FIG. 2.
NiO(A) tested in hydroxide at -' 400 m V SCE. Crack initiation a n d thick corrosion products on electropolished surface. Ft6. 3. NiO(A) which failed in carbonate at - 675 mV SCE. Typical intergranular cracking. Fro. 4. NiO(Q) tested in c a r b o n a t e at - 675 mV SCE. (a) l n t e r g r a n u l a r fracture surface with secondary cracks along martensite laths. (b) Metallographic n a t u r e o f cracking.
[facing p. 4701
m 5(b)
I
FIG. 5. Si iron tested in hydroxide at - 950 mV SCE. (a) Intergranular fracture surface with transgranular cracks initiating at slip steps. (b) Resultant stepped transgranular fracture surface. FIG. 6. Ni2 tested under freely corroding conditions in chloride. (a) Annealed specimen showing typical transgranular cracking together with some intergranular cracking. (6) Quenched specimen, some intergranular cracking and transgranular cracking which appears unrelated to any structural features.
Jt
FIG. 7. S1S 2103(HTA) tested in liquid ammonia at -t- 700 mV vs ~PdH. (a) Wide range of features. (b) Stepped region near initiation cf. Fig. 5b. (c) Flatter region crossing what is thought to be a grain boundary.
FIG. 8. SIS 2103(Q) which failed in liquid ammonia at ÷ 700 mV [3PdH. Intergranular and transgranular cracking together with some mechanical tearing prior to final mechanical failure. FIG. 9. (a) SIS 2103(HTA) tested in liquid ammonia at ~ 700 mV f~PdH showing crack initiation at slip lines. (b) SIS 2103(Q) tested in liquid ammonia at -t 700 mV ~PdH. Crack initiation apparently unrelated to structure.
FIG. 10. Fracture surfaces of specimens unambigouslyhydrogen embrittled. (a) NiO(A). (b) S1S 2103(N). (c) S1S 2103(Q).
"~
" "t o
0
( 7 "~'~d~.
r , , * ' . . , o- ; ~ , "
..-.~,,~. 't:.~
,,,~
"
~
,J
E ~i " .',,~
10~
c>"
• • :,~. ,,'~_~ ~,v, ~ . , ~ r ~ , ' ""
0
,L~,.,~, ~ ,
FIG. 11. FIG. 12. FIG. 13.
o .,.,_vp
•
~
~'~6
1""
",
Spheroidized eutectoid steel which failed in hydroxide at -- 950 mV SCE. Carbides remaining in cracks. Mo4(A) tested in hydroxide at - 750 mV SCE. Pits in intergranular fracture surface resulting from carbide dissolution. Typical region prior to final mechanical failure showing so-called serpentine glide.
The fractography of stres.scorrosion cracking in carbon steels
471
solutions, and then some general observations pertaining to carbide reactivity and the transition region between SCC and final mechanical failure. a Nitrate, carbonate and hydroxide solutions In these solutions annealed steels cracked predominantly intergranularly (Fig. 3). Intergranular cracking in this ease along the prior austenite grain boundaries, was also predominant in the quenched steel, together with some attack along the martensite lath boundaries (Fig. 4a, b). An exception to this generalization was the varying amounts of transgranular cracking which occurred in annealed steels. This was potential dependent tending to occur outside the range of maximum susceptibility and was promoted by the addition of certain alloying elements. 17 Thus Mo and Ti promote transgranular fissures, rather than cracking, in nitrates, Ti transgranular cracking in carbonates and Si transgranular cracking in hydroxide. This transgranular cracking appeared related--initially at least--to attack on slip lines (Fig. 5a). The resulting fracture surface exhibiting stepped and coalescence features (Fig. 5b). b Chloride sohltions Only steels containing approximately 1% or more Ni failed in the chloride solution. The fractography of the annealed steels has been described in detail elsewhere, a5 Briefly all the failures were predominantly transgranular except for the most susceptible steel, containing approximately 2%Ni, where mixed inter/transgranular failure occurred (Fig. 6a). These transgranular features were similar to those reported for austenitic stainless steels tested under similar conditions. The two predominant features were the so-called fan or fern-like regions and areas having a striated appearance. The susceptible quenched steels failed by a mixture of intergranular cracking along prior austenite boundaries and transgranular cracking across the martensite laths (Fig. 6b). c Liquid ammonia Annealed steels failed predominantly transgranularly when tested in liquid ammonia. A wide variety of fracture surface features were observed (Fig. 7a) ranging from steps (Fig. 7b) to flatter regions (Fig. 7c). The quenched specimens failed by a mixed mode (Fig. 8) intergranularly along prior austenite boundaries, transgranular cracking which did not appear related to martensite laths and a proportion of dimple formation prior to final mechanical failure. Crack initiation studies indicated attack was not dependent on pitting (Figs. 9a, 9b) but was related to slip line attack in the annealed steel (Fig. 9a). d Hydrogen embrittled specimens For comparative purposes some specimens were unambiguously hydrogen embrittled by holding at very negative potentials in the respective solutions (See Fig. 1) or by cathodic charging at I0 mA/cm in IN H2SO4. Annealed steels failed by a mixture of quasicleavage and dimple formation (Fig. 10a). With the fine grained steel SIS2103(N) this was not so apparent as Fig. 10b shows. The quenched steels also failed in a complex manner: intergranular, quasicleavage and dimple formation all being observed, the latter two being shown in Fig. 10c.
472
B. POULSON
e General observation
Although the dissolution characteristics of ferrite, cementite and their various interfaces have been extensively studied under bulk environmental conditions, 18,1D there have been few reported observation relating to crack propagation. In this study little--if any---carbide dissolution appeared to occur during crack propagation in any solution except hydroxide. Even in hydroxides carbide dissolution was seldom apparent with cracks containing undissolved carbides (Fig. I I). An exception to this was the steel containing 5%Mo where carbide dissolution appeared complete (Fig. 12). Small regions were sometimes observed, between the typical SCC region and the .final mechanical failure, which have been termed serpentine glide (Fig. 13). It has been suggested 2° such features are the result of the initial opening or blunting of a sharp crack and are therefore evident that the SCC was indeed sharp; but no figure was indicated as to what sharp meant. However, Beachem21 has shown similar features at the base of a filed notch which was subsequently opened. DISCUSSION First some fractographic observations are considered which suggest that certain proposed modes of crack propagation are unlikely to be operative. Next the idea that crack propagation occurs by anodic dissolution is examined to see if it is consistent with the fracture surface morphology. And finally the practical, that is the failure analysis, implications of the work are discussed. 1 Mechanistic aspects
In recent years the concept that hydrogen embrittlement may be responsible for SCC of not only high strength steels but also lower strength steels has been receiving increasing attention?2 Thus it has been suggested that the cracking of carbon steels in hydroxides,23 nitrates, u liquid ammonia~5 and Ni containing steels in chlorides2e is due to hydrogen. While there is substantial evidence that hydrogen embrittlement contributes to crack growth in high strength steels the evidence is much less favourable for the lower strength alloys. There is no dispute to the contention that they can be made to fail by severe cathodic charging, what is disputed is the relevance of these observations to conditions normally thought to exist during crack propagation. If hydrogen is involved in crack propagation it would appear reasonable to expect that: (a) The fracture surface of an unambiguously hydrogen embrittled specimen and that produced during SCC would at least be similar. In the present investigation this was not found to be so. The one exception being the varying amounts of cracking along prior austenite grain boundaries which occurred in quenched steels either hydrogen embrittled or stress corroded. (b) Although it has been shown by others 27 that hydrogen entry can occur at anodic potentials, a pre-requisite would seem to be the formation of pits in which the conditions for hydrogen evolution exist. This means that if hydrogen is responsible for cracking at anodic potentials, crack nucleation must be associated with pitting. Although some cracks were associated with pits this was not the usual situation. Transgranular cracks initiated at emergent slip steps while intergranular cracks started as grain boundary attack. Furthermore, the suggestion that hydrogen aids martensite formation by lowering
The fractography of stress corrosion cracking in carbon steels
473
the stacking fault energy and the martensite subsequently rapidly dissolves 2s or fails in a brittle manner 29 would seem most unlikely. Firstly in the annealed (ferritic) steels the formation of martensite during the conditions existing during a SCC test would appear most improbable, no matter how much hydrogen was adsorbed. For the quenched steels no brittle failure of the pre-existing martensite laths occurred, failure being by void coalescence in mechanical tests. Also localized dissolution of the martensite did not appear to occur either, suggesting such processes are unimportant in crack propagation. The stress adsorption model advocated by Uhlig s° is, as others have indicated, rather difficult to test. In particular it is difficult to design experiments which could lead to its refutation. Although the fractographic features observed in this investigation could possible have resulted from such a process the available evidence a7 is strongly in favour of a dissolution model. The brittle film model 1° proposes that crack propagation occurs discontinuously on a fine scale by the repeated formation and rupture of a brittle surface layer; dissolution not contributing to crack growth. Much of the early evidence for this model derives from the brass/tarnishing ammonical solution, system and was fractographic. Thus striations at 90 ° to the crack propagation direction and supposedly representing the stepwise advance of the crack were noted; but only on approximately 30% of the fracture surface and only when the brass was cold worked. More recently this interpretation has been questioned, al Certainly where striations are shown adjacent to the final mechanical failure--to identify the crack direction--they probably resulted from serpentine glide rather than SCC. In the present investigation quite detailed features were observed on intergranular fracture surfaces, suggesting that the observed lack of striations was not due to them being obscured by subsequent corrosion. Thus there appears no fractographic evidence which can support the idea of the brittle film model operating in any of the systems investigated in this study. It has been suggested that the fractographic detail observed on specimens of austenitic steel which had failed in chloride solutions was strong evidence against a dissolution model operating. Thus Nielsen 1 commented that in the absence of any known environmental effects such failures would be classified as cleavage or quasicleavage. Similarly Pugh 32 recently remarked that he couldn't visualize such features resulting from a dissolution process. However, neither of these are sound reasons for rejecting the importance of dissolution. In fact, similar morphologies have been produced on unstressed specimens by dissolution alone. ~s Furthermore, Scullyz reasoned that cleavage over the distances required to produce such features would be detected by elongation or acoustic measurements. He then developed a model which could explain the wide range of features commonly observed. This is shown schematically in Fig. 14. It is based on cracks propagating by anodic dissolution (possibly as tunnels) along different planes. Ductile tearing then causes these cracks to coalesce and the resultant surface depends on the uniformity of tearing and the ratio of tearing to dissolution.
2 Dependence of crack path on electrochemicalfactors As summarized in Table 4 there appears to be marked similarities between failures in, on the one hand nitrates, carbonates and hydroxides, and on the other in liquid
474
B. POOl.SON Four crocks on parallel planes The small ligaments of metal between crocks in a single groin will undergo considerable distortion as o result of the external stress These ligaments will break at some time resulting in o stepped crock profile
Four cracks coalesce "~o form
Fzc. 14.
a
single crock
During the development of these configurations severe distortion of some of cracks may occur due "to the action of shearing stresses resulting in curved fan like features Schematic diagram showing development o f fracture surface topography (after Scully~).
TABLE 4.
Steel
SUMMARY OF CRACK MORPHOLOGIES
Heat treatment
Solution
Crack morphologies
NiONiO Ni, O, Mo4
Annealed Annealed Annealed
Nitrates Carbonates Hydroxides
f Predominantly intergranular, d[ attack on unstressed specimens is structurally sensitive
NiO NiO NiO
Quenched Quenched Quenched
Nitrates Carbonates Hydroxides
r Intergranular along prior ~ austenite boundaries plus some cracking along martensite laths
SiS 2103 Nil, Ni2, Ni3 and Ni4 *Mild steel
Liquid ammonia Annealed Annealed Chlorides Normalized CO-COz-H20
SIS 2103 Nil, Ni2, Ni3 and Ni4
Quenched Quenched
Liquid ammonia Chlorides
f Predominantly transgranular ~ cracks which initiate at slip steps r Along prior austenite boundaries ~ plus some transgranular cracking not along martensite laths
*From M. Kowaka and S. Nagura, Corrosion 24, 427, 1968. ammonia, concentrated chloride and from other work CO inhibited solutions. Mechanistically this is important because in the former solutions dissolution is thought to be restricted by an oxide, typical polarization curves exhibiting an active to passive transition (Fig. 15a). However, in liquid a m m o n i a le concentrated chloride is and CO containing solutions s8 dissolution is thought to be hindered--at leastinitially--bya thin adsorbed layer whose formation is reversible. In this case polarization curves do not
The fractography of stress corrosion cracking in carbon steels
475
exhibit an active to passive transition. However, in chlorides and CO inhibited solutions the anodic polarization curves exhibit an adsorption induced inflection as shown schematically in Fig. 15(b). Furthermore, the crack path in the oxide forming solutions is structurally preexisting, attack on unstressed specimens is structurally dependent and in nitrates intergranular corrosion can occur which is quantitatively related to susceptibility. However, in the solutions where adsorption effects are thought to be important cracks initiate at slip steps. Thus the path is in a sense strain generated. The obvious suggestion is that in the oxide forming solutions thick corrosion products form whose formation is structurally dependent. ~s That is they form last at grain boundaries which also dissolve more rapidly. Subsequently the oxide either never manages to form at the crack tip or is continually ruptured. In the adsorbing solutions a thin layer forms rather quickly and its formation is possibly less structurally sensitive. Subsequently when slip ruptures this layer localized dissolution can occur at the emergent slip step.aS Such a sharp division between oxide-forming solutions and solutions where adsorption inhibits dissolution is made for the sake of convenience of discussion and more correctly represents extreme cases. This is readily apparent when we consider some transgranular cracking does occur in hydroxides and some intergranular cracking was found in chlorides. Furthermore there is still considerable discussion as to whether an adsorbed layer precedes the formation of a passive film. a4
2 Practical flnplications As others have indicated a~ the failure of annealed mild steels by SCC can no longer be regarded as occurring exclusively intergranularly, thus the existence of intergranular cracking cannot be used as a diagnostic criteria for SCC. This means that transgranular SCC must be distinguished from transgranular fatigue or corrosion fatigue. Where fractographic detail is observed by corrosion products and even second order characteristics, such as "clamshell markings" on fatigue failures cannot be observed, reliance on metallographic examination and the failed components history must be used. This is not always decisive. Failures of quenched mild steels involve similar
5CC
SCC
~
CC
(b) Log C.D.
Log G.D
FIo. 15. (a) Schematic anodic polarization curve exhibiting an active to passive transition and the potential ranges which are most likely to promote SCC, (b) Schematic anodic polarization curve exhibiting adsorption induced inhibition and the potential range most likely to promote SCC.
476
B. POULSOr~
fractographic diagnostic problems. Thus intergranular cracking, along prior austenite grain boundaries, can result from SCC, hydrogen embrittlement, quench cracking and temper embrittlement. Another important point is that the crack path depends markedly on the metallurgical structure and differences in the parent metal, the heat affected zone and the weld metal are to be expected. Furthermore this effect of structure on the crack path depends on the solution and the relative susceptibilities of each zone is unlikely to be invariant. Thus in liquid ammonia practical experience 16 is that cracks tend to occur in the weld metal, while in hydroxides the parent metal or the H.A.Z. is preferred. If a failure of a mild steel component is intergranular and SCC is diagnosed but environmental conditions responsible are unknown, extensive carbide attack is reason to suspect hydroxides. The potential dependence of cracking is another diagnostic method which has been used in service failure analysis. 36 It must be pointed out that the number of solutions known to cause SCC is increasing. It must not become an obsession to prove a SCC failure resulted from the presence of nitrates or hydroxides.37 FUTURE WORK Much of the fractographicwork on SCC has been ambiguous. That is, the features observed could be interpreted in terms of a number of mechanistic ideas. A noteworthy exception to this is Scully's work on SCC in Ti alloys. 2,~ Real progress in this field will probably result from two separate approaches. The first, the application of a wide range of other techniques to examine the fracture surface. These would include dislocation etching, channeling patterns, X-ray broadening, electron diffraction, ion spluttering or laser probe-mass spectroscopy, auger spectroscopy, ESCA, transmission electron microscopy (after thinning from one side) and electrochemical techniques designed to assess the protectiveness of any film formation on the fracture surface --which the physical techniques cannot do. The second approach involves changing a variable during crack propagation and observing the effect on crack propagation and the fracture surface. From a practical aspect much work needs to be done on the fractography of real materials under service conditions. For example, welded or coldworked alloy steels which after failure have been left for a period prior to examination. CONCLUSIONS 1. In nearly all cases the fracture surfaces of carbon steels which had failed by SCC were clean enough for fractographic details to be observed. 2. The crack path was found to depend on--heat treatment, environment, potential and alloying element. 3. For annealed steels in solutions (nitrate, hydroxide and carbonate) which promoted thick oxide formation, cracking was predominantly intergranular. In chloride and liquid ammonia where thin films or an adsorbed layer are thought to be important cracking was predominantly transgranular. 4. For quenched steels cracking was transgranular and intergranular along the prior austenite grain boundaries. In oxide-forming solutions the transgranular cracking was along martensite lath boundaries, while in chloride and liquid ammonia it was not apparently related to any structural feature. 5. The available fractographic evidence appears inconsistent with either hydrogen
The fractography of stress corrosion cracking in carbon steels
477
e m b r i t t l e m e n t o r brittle films p l a y i n g a role in c r a c k p r o p a g a t i o n . It is, h o w e v e r , c o n s i s t e n t w i t h c r a c k p r o p a g a t i o n by a n o d i c d i s s o l u t i o n p o s s i b l y a i d e d by t e a r i n g . Acknowledgements The author wishes to thank: Prof. R. N. Parkins and Civ.-Eng. Hans Arup for helpful discussions and valuable criticisms, and the staffof the S.E.M.s in the Department of Chemistry, University of Newcastle upon Tyne, and the Danish Technical Institute at Tastrup, for their help. REFERENCES 1. N. A. NIELSEN,Fundamental Aspects of Ctress Corrosion Cracking (Eds. R. W. STAEHLEel al.), p. 308. NACE, Houston, Texas (1967). 2. J. C. SCULLY, The Theory of Stress Corrosion Cracking in Alloys, p. 127. NATO, Brussels (1971). 3. F. E. WATKINSONand J. C. SCULLY, Corros. Sci. 12, 905 (1972). 4. W. R. MIDDLE'tON and R. N. PARKINS, Corrosion 28, 881 (1972). 5. D. T. POWELLand J. C. SCULLV, Corrosion 24, 151 (1968). 6. R. J. H. WANHILL, Corrosion 29, 435 (1973). 7. G. DEAr% Ph.D. Thesis, University of Newcastle upon Tyne (1971). 8. V. J. CALANOELOand M. S. FERC,trsoN, Corrosion 25, 509 (1969). 9. J. C. SCULLY, Paper presented at the International Conference on Stress Corrosion Cracking and Hydrogen Embrittlement of Iron Base Alloys. Firminy, France, June 1973, to be published by NACE. 10. E. N. PUGH, J. V. CRAIG and A. J. SEDRIKS,Fundamental Aspects of Stress Corrosion Cracking (Eds. R. W. S'rAErtLEet al.), p. 118. NACE, Houston, Texas (1967). 11. G. M. SPARKSand J. C. SCtrLLY, Corros. Sci. 11, 641 (1971). 12. R. D. BARERand B. F. PE'rEF,S, Why Metals Fail. Gordon and Breach, New York (1970). 13. J. E. TRU'MAN,Paper presented at the International Conference on Stress Corrosion Cracking and Hydrogen Embrittlement of Iron Base Alloys. Firminy, France, June 1973, to be published by NACE. 14. B. Pout.sos, Ph.D. Thesis, University of Newcastle upon Tyne (1972). 15. B. PouLsos and R. N. PARKINS, Corrosion 29, 414 (1973). 16. B. POULSONand H. ARtrp, In preparation. 17. R. N. PAR.KINS,B. PouLSos and P. SLATTERY,In preparation. 18. R. N. PARKINS, The Theory of Stress Corrosion Cracking in Alloys (Ed. J. C. SCULLY),p. 167. NATO, Brussels (1971). 19. C. J. CROS, J. H. PAYERand R. W. STAEHLE,Corrosion 27, l (1971). 20. J. C. WILLIAMS,Discussion to ref. I. 21. C. D. BEACHEM,Electron Fractography, STP436, p. 59. ASTM, Philadelphia (1968). 22. N. A. NIELSES, Corrosion 27, 173 (1971). 23. F. P. A. ROBIr'~SONand L. G. NEE,Proc. 2ndbzternational Congress on Metallic Corrosion, p. 171. NACE, Houston 0969). 24. B. E. WILDE, Corrosion 27, 326 (1971). 25. F. J. RADO and D. H. OERTLE,N A C E Annual Conference. Preprint, Chicago, March 1971. 26. H. H. UHLlC~,Physical Metallurgy of Stress Corrosion Fracture (Ed. R. RHODIN), p. 258. Interscience, New York (1959). 27. C. F. BARa'H, E. A. STEIGERWALDand A. R. TPaASO, Corrosion 25, 353 (1969). 28. S. S. BIRLEYand D. TROMANS,Corrosion 27, 63 (1971). 29. C. Et~ELEANU,JlSI 143, 140 (1953). 30. H. H. UHLIG, Fundamental Aspects of Stress Corrosion Crackhtg (Ed. R. W. S'rAEHLEet al.), p. 86. NACE, Houston (1969). 31. D. J. LEEs, Corros. Sci. 12, 811 (1972). 32. E. N. PtJGH, Discussion at International Conference on Stress Corrosion Cracking and Hydrogen Embrittlement of Iron Base Alloys. Firminy, France, June 1973, to be published by NACE. 33. K. E. HEUSLERand G. H. CARTLEDGE,J. electrochem. Soc. 108, 733 (1961). 34. J. O'M. BOCKRISand P. K. StraRAMASVAS, Corros. Sci. 10, 435 (1970). 35. M. KOWAKAand S. NAGA'rA, Corrosiot, 24, 422 (1968). 36. B. Potrt.sos, L. C. HENRIKSENand H. ARUP, Br. Corros. J. In press. 37. R. N. PARrdNS, Paper presented at the International Conference on Stress Corrosion Cracking and Hydrogen.Embrittlement of Iron Base Alloys. Firminy, France, June 1973, to be published by NACE. 38. A. BRows, J. T. HAmUSOS and R. WILrdNS, ibid.
THE FRACTOGRAPHY OF STRESS CORROSION CRACKING IN CARBON STEELS* B. POULSON Materials Research, A.T.D., Clarke Chapman Ltd., Gateshead, England Abstract--The stress corrosion cracking of carbon steels in a number of environments (nitrates, hydroxides, carbonates, chlorides and liquid ammonia) has been studied by scanning electron microscopy. The crack path depends on both heat treatment and environmental conditions. In annealed steels it appears that solutions which produce thick corrosion products induce intergranular cracking, while in solutions where adsorption induced inhibition occurs or thin films form, cracking is transgranular. Although the available fractographic evidence is ambiguous, it appears to rule out either hydrogen embrittlement or brittle films as being important in crack propagation. However, the fracture surface morphologies are consistent with the cracks propagating by anodic dissolution, possibly aided by mechanical tearing. INTRODUCTION FRACTOGRAPHY is an important experimental technique in b o t h failure analysis and mechanistic studies. This is because the unambiguous interpretation o f fracture surface m o r p h o l o g y would allow how the crack propagated to be determined. N u m e r o u s papers have dealt with fractographic aspects o f stress corrosion cracking (SCC) in: stainless steels, 1.2 aluminium alloys, a,4 titanium alloys, 5, 6 magnesium alloys, 7 high strength steels, a,9 and copper base alloys. 1°,xl However, carbon steels have been somewhat neglected. 9 This is possibly due to the mistaken belief that mild steels only crack intergranularly. 12,1a Furthermore, m a n y o f the solutions which cause cracking produce thick corrosion products and it is assumed 9 that such layers will obscure any fractographic detail. The purpose of this paper is to summarize some fractographic observations obtained in recent investigations 14,~5,~a o f SCC in carbon steels. EXPERIMENTAL The SCC testing technique has been described in detail elsewhere. 14 Briefly constant strain rate (2.2 × 10-6/s) tests were performed on unnotched cylindrical tensile specimens; either electropolished or finished to grade 600 emery paper. The compositions and heat treatments of the steels examined are shown in Tables 1 and 2. The SCC tests were carried out in a number of solutions over a wide range of potentials (Table 3). Immediately failure occurred specimens were removed from the cell, cleaned in methanol and stored until they were examined in the SEM. RESULTS The potential ranges over which SCC occurred for certain steels in the different solutions are shown in Fig. 1. The results for the liquid a m m o n i a having been converted to the normal hydrogen scale (assuming p N H ~ = 16 and p H + ~ p N H ) + they are therefore somewhat approximate. Elsewhere in this paper they are quoted as they were measured, i.e. against a [~P d - H electrode in liquid a m m o n i a . These potential ranges p r o m o t i n g SCC are in general agreement with previous work. M o s t o f the fracture surface o f specimens removed f r o m the environment directly they failed was rather clean. A n exception was the failures in nitrate solutions at potentials ( - - 200 to *Manuscript received 8 July 1974. 469
]], POULSON
470 TABLE 1. C
Si
Mn
S
P
Ni
AI
0.09 0.09 0"075 0.08 0.08 0' 1 0" 14
0.06 0.09 0.06 0.09 0.04 0.22 0.36
0.82 0.48 0.19 0.75 0-49 1 1"73
0"01 0.02 0.04 0.03 0.04 0-05 0.013
0.02 0.01 0.03 0.02 0.04 0.06 0.01
0.04 1.05 2.15 4'57 6.07 0.02 --
0.02 0'03 0.06 0'05 0.02 0.05
Alloy NiO Nil Ni2 Ni3 Ni4 Mo4 SIS 2103
COMPOSITION OF STEELS USED
TABLE 2. As received Annealed Quenched H i g h T e m p . Anneal.
TABLE 3.
Others
5 /o,o M o
--
0-I ,%Cu
HEAT TREATMENTS USED (N) (A) (Q) (HTA)
i.e. normalized 950°C for 1 h, furnace cooled 950°C for 1 h, water quenched 1150°C for 3 h, furnace cooled
ENVIRONMENTAL CONDITIONS INVESTIGATED
Temperature
Solution
Nitrate, 4N N H 4 NO3 Chloride, approx. 45 ~ M g C I 2 Hydroxide, 35 ,%NaOH Carbonate, 1 : 1 I N Na~ COs: 1N N a H C O s A m m o n i a , liquid a m m o n i a
pH
Potential range m V
Boiling Boiling at 154°C Boiling
4
-- 600 to d 1200 S.C.E. - 1000 to - 150 S.C.E. -- 1300 to + 450 S.C.E.
95°C Ambient
9 (at 25°C)
-- 1100 to -- 500 S.C.E. -- 1000 to + 2000 P d H
1500 I
OOO
I000 500
Liquid
Nitrate
IiJ m
ammonia
u')
0
._o
E
E
-500
13_
Data for
-
NIO(A) in nitrate in chloride
-tO00 --Ni4(A) NiO(A)
-
FIG. 1.
150C
t
~
i
500 n
I
in c a r b o n o t e
SIS2103(Q} in hquid ammonia ~ ~ NiO(A) in hydroxide ~ydrogen ernbritl'lernent
f
-
-
I000
Schematic d i a g r a m s h o w i n g potential ranges over which S C C o f carbon steels occurs in various solutions.
q- 100 mV S.C.E.) where the general corrosion rate was high, and failures (Fig. 2) in hydroxide at potentials where oxygen evolution occurred. Because of certain similarities the results of tests in nitrate, carbonate and hydroxide solutions are presented first. This is followed by the results of tests in liquid ammonia and chloride
, '
,4
~
~,~.~,~7.."~
3
"
~
~:?!
FIG. 2.
NiO(A) tested in hydroxide at -' 400 m V SCE. Crack initiation a n d thick corrosion products on electropolished surface. Ft6. 3. NiO(A) which failed in carbonate at - 675 mV SCE. Typical intergranular cracking. Fro. 4. NiO(Q) tested in c a r b o n a t e at - 675 mV SCE. (a) l n t e r g r a n u l a r fracture surface with secondary cracks along martensite laths. (b) Metallographic n a t u r e o f cracking.
[facing p. 4701
m 5(b)
I
FIG. 5. Si iron tested in hydroxide at - 950 mV SCE. (a) Intergranular fracture surface with transgranular cracks initiating at slip steps. (b) Resultant stepped transgranular fracture surface. FIG. 6. Ni2 tested under freely corroding conditions in chloride. (a) Annealed specimen showing typical transgranular cracking together with some intergranular cracking. (6) Quenched specimen, some intergranular cracking and transgranular cracking which appears unrelated to any structural features.
Jt
FIG. 7. S1S 2103(HTA) tested in liquid ammonia at -t- 700 mV vs ~PdH. (a) Wide range of features. (b) Stepped region near initiation cf. Fig. 5b. (c) Flatter region crossing what is thought to be a grain boundary.
FIG. 8. SIS 2103(Q) which failed in liquid ammonia at ÷ 700 mV [3PdH. Intergranular and transgranular cracking together with some mechanical tearing prior to final mechanical failure. FIG. 9. (a) SIS 2103(HTA) tested in liquid ammonia at ~ 700 mV f~PdH showing crack initiation at slip lines. (b) SIS 2103(Q) tested in liquid ammonia at -t 700 mV ~PdH. Crack initiation apparently unrelated to structure.
FIG. 10. Fracture surfaces of specimens unambigouslyhydrogen embrittled. (a) NiO(A). (b) S1S 2103(N). (c) S1S 2103(Q).
"~
" "t o
0
( 7 "~'~d~.
r , , * ' . . , o- ; ~ , "
..-.~,,~. 't:.~
,,,~
"
~
,J
E ~i " .',,~
10~
c>"
• • :,~. ,,'~_~ ~,v, ~ . , ~ r ~ , ' ""
0
,L~,.,~, ~ ,
FIG. 11. FIG. 12. FIG. 13.
o .,.,_vp
•
~
~'~6
1""
",
Spheroidized eutectoid steel which failed in hydroxide at -- 950 mV SCE. Carbides remaining in cracks. Mo4(A) tested in hydroxide at - 750 mV SCE. Pits in intergranular fracture surface resulting from carbide dissolution. Typical region prior to final mechanical failure showing so-called serpentine glide.
The fractography of stres.scorrosion cracking in carbon steels
471
solutions, and then some general observations pertaining to carbide reactivity and the transition region between SCC and final mechanical failure. a Nitrate, carbonate and hydroxide solutions In these solutions annealed steels cracked predominantly intergranularly (Fig. 3). Intergranular cracking in this ease along the prior austenite grain boundaries, was also predominant in the quenched steel, together with some attack along the martensite lath boundaries (Fig. 4a, b). An exception to this generalization was the varying amounts of transgranular cracking which occurred in annealed steels. This was potential dependent tending to occur outside the range of maximum susceptibility and was promoted by the addition of certain alloying elements. 17 Thus Mo and Ti promote transgranular fissures, rather than cracking, in nitrates, Ti transgranular cracking in carbonates and Si transgranular cracking in hydroxide. This transgranular cracking appeared related--initially at least--to attack on slip lines (Fig. 5a). The resulting fracture surface exhibiting stepped and coalescence features (Fig. 5b). b Chloride sohltions Only steels containing approximately 1% or more Ni failed in the chloride solution. The fractography of the annealed steels has been described in detail elsewhere, a5 Briefly all the failures were predominantly transgranular except for the most susceptible steel, containing approximately 2%Ni, where mixed inter/transgranular failure occurred (Fig. 6a). These transgranular features were similar to those reported for austenitic stainless steels tested under similar conditions. The two predominant features were the so-called fan or fern-like regions and areas having a striated appearance. The susceptible quenched steels failed by a mixture of intergranular cracking along prior austenite boundaries and transgranular cracking across the martensite laths (Fig. 6b). c Liquid ammonia Annealed steels failed predominantly transgranularly when tested in liquid ammonia. A wide variety of fracture surface features were observed (Fig. 7a) ranging from steps (Fig. 7b) to flatter regions (Fig. 7c). The quenched specimens failed by a mixed mode (Fig. 8) intergranularly along prior austenite boundaries, transgranular cracking which did not appear related to martensite laths and a proportion of dimple formation prior to final mechanical failure. Crack initiation studies indicated attack was not dependent on pitting (Figs. 9a, 9b) but was related to slip line attack in the annealed steel (Fig. 9a). d Hydrogen embrittled specimens For comparative purposes some specimens were unambiguously hydrogen embrittled by holding at very negative potentials in the respective solutions (See Fig. 1) or by cathodic charging at I0 mA/cm in IN H2SO4. Annealed steels failed by a mixture of quasicleavage and dimple formation (Fig. 10a). With the fine grained steel SIS2103(N) this was not so apparent as Fig. 10b shows. The quenched steels also failed in a complex manner: intergranular, quasicleavage and dimple formation all being observed, the latter two being shown in Fig. 10c.
472
B. POULSON
e General observation
Although the dissolution characteristics of ferrite, cementite and their various interfaces have been extensively studied under bulk environmental conditions, 18,1D there have been few reported observation relating to crack propagation. In this study little--if any---carbide dissolution appeared to occur during crack propagation in any solution except hydroxide. Even in hydroxides carbide dissolution was seldom apparent with cracks containing undissolved carbides (Fig. I I). An exception to this was the steel containing 5%Mo where carbide dissolution appeared complete (Fig. 12). Small regions were sometimes observed, between the typical SCC region and the .final mechanical failure, which have been termed serpentine glide (Fig. 13). It has been suggested 2° such features are the result of the initial opening or blunting of a sharp crack and are therefore evident that the SCC was indeed sharp; but no figure was indicated as to what sharp meant. However, Beachem21 has shown similar features at the base of a filed notch which was subsequently opened. DISCUSSION First some fractographic observations are considered which suggest that certain proposed modes of crack propagation are unlikely to be operative. Next the idea that crack propagation occurs by anodic dissolution is examined to see if it is consistent with the fracture surface morphology. And finally the practical, that is the failure analysis, implications of the work are discussed. 1 Mechanistic aspects
In recent years the concept that hydrogen embrittlement may be responsible for SCC of not only high strength steels but also lower strength steels has been receiving increasing attention?2 Thus it has been suggested that the cracking of carbon steels in hydroxides,23 nitrates, u liquid ammonia~5 and Ni containing steels in chlorides2e is due to hydrogen. While there is substantial evidence that hydrogen embrittlement contributes to crack growth in high strength steels the evidence is much less favourable for the lower strength alloys. There is no dispute to the contention that they can be made to fail by severe cathodic charging, what is disputed is the relevance of these observations to conditions normally thought to exist during crack propagation. If hydrogen is involved in crack propagation it would appear reasonable to expect that: (a) The fracture surface of an unambiguously hydrogen embrittled specimen and that produced during SCC would at least be similar. In the present investigation this was not found to be so. The one exception being the varying amounts of cracking along prior austenite grain boundaries which occurred in quenched steels either hydrogen embrittled or stress corroded. (b) Although it has been shown by others 27 that hydrogen entry can occur at anodic potentials, a pre-requisite would seem to be the formation of pits in which the conditions for hydrogen evolution exist. This means that if hydrogen is responsible for cracking at anodic potentials, crack nucleation must be associated with pitting. Although some cracks were associated with pits this was not the usual situation. Transgranular cracks initiated at emergent slip steps while intergranular cracks started as grain boundary attack. Furthermore, the suggestion that hydrogen aids martensite formation by lowering
The fractography of stress corrosion cracking in carbon steels
473
the stacking fault energy and the martensite subsequently rapidly dissolves 2s or fails in a brittle manner 29 would seem most unlikely. Firstly in the annealed (ferritic) steels the formation of martensite during the conditions existing during a SCC test would appear most improbable, no matter how much hydrogen was adsorbed. For the quenched steels no brittle failure of the pre-existing martensite laths occurred, failure being by void coalescence in mechanical tests. Also localized dissolution of the martensite did not appear to occur either, suggesting such processes are unimportant in crack propagation. The stress adsorption model advocated by Uhlig s° is, as others have indicated, rather difficult to test. In particular it is difficult to design experiments which could lead to its refutation. Although the fractographic features observed in this investigation could possible have resulted from such a process the available evidence a7 is strongly in favour of a dissolution model. The brittle film model 1° proposes that crack propagation occurs discontinuously on a fine scale by the repeated formation and rupture of a brittle surface layer; dissolution not contributing to crack growth. Much of the early evidence for this model derives from the brass/tarnishing ammonical solution, system and was fractographic. Thus striations at 90 ° to the crack propagation direction and supposedly representing the stepwise advance of the crack were noted; but only on approximately 30% of the fracture surface and only when the brass was cold worked. More recently this interpretation has been questioned, al Certainly where striations are shown adjacent to the final mechanical failure--to identify the crack direction--they probably resulted from serpentine glide rather than SCC. In the present investigation quite detailed features were observed on intergranular fracture surfaces, suggesting that the observed lack of striations was not due to them being obscured by subsequent corrosion. Thus there appears no fractographic evidence which can support the idea of the brittle film model operating in any of the systems investigated in this study. It has been suggested that the fractographic detail observed on specimens of austenitic steel which had failed in chloride solutions was strong evidence against a dissolution model operating. Thus Nielsen 1 commented that in the absence of any known environmental effects such failures would be classified as cleavage or quasicleavage. Similarly Pugh 32 recently remarked that he couldn't visualize such features resulting from a dissolution process. However, neither of these are sound reasons for rejecting the importance of dissolution. In fact, similar morphologies have been produced on unstressed specimens by dissolution alone. ~s Furthermore, Scullyz reasoned that cleavage over the distances required to produce such features would be detected by elongation or acoustic measurements. He then developed a model which could explain the wide range of features commonly observed. This is shown schematically in Fig. 14. It is based on cracks propagating by anodic dissolution (possibly as tunnels) along different planes. Ductile tearing then causes these cracks to coalesce and the resultant surface depends on the uniformity of tearing and the ratio of tearing to dissolution.
2 Dependence of crack path on electrochemicalfactors As summarized in Table 4 there appears to be marked similarities between failures in, on the one hand nitrates, carbonates and hydroxides, and on the other in liquid
474
B. POOl.SON Four crocks on parallel planes The small ligaments of metal between crocks in a single groin will undergo considerable distortion as o result of the external stress These ligaments will break at some time resulting in o stepped crock profile
Four cracks coalesce "~o form
Fzc. 14.
a
single crock
During the development of these configurations severe distortion of some of cracks may occur due "to the action of shearing stresses resulting in curved fan like features Schematic diagram showing development o f fracture surface topography (after Scully~).
TABLE 4.
Steel
SUMMARY OF CRACK MORPHOLOGIES
Heat treatment
Solution
Crack morphologies
NiONiO Ni, O, Mo4
Annealed Annealed Annealed
Nitrates Carbonates Hydroxides
f Predominantly intergranular, d[ attack on unstressed specimens is structurally sensitive
NiO NiO NiO
Quenched Quenched Quenched
Nitrates Carbonates Hydroxides
r Intergranular along prior ~ austenite boundaries plus some cracking along martensite laths
SiS 2103 Nil, Ni2, Ni3 and Ni4 *Mild steel
Liquid ammonia Annealed Annealed Chlorides Normalized CO-COz-H20
SIS 2103 Nil, Ni2, Ni3 and Ni4
Quenched Quenched
Liquid ammonia Chlorides
f Predominantly transgranular ~ cracks which initiate at slip steps r Along prior austenite boundaries ~ plus some transgranular cracking not along martensite laths
*From M. Kowaka and S. Nagura, Corrosion 24, 427, 1968. ammonia, concentrated chloride and from other work CO inhibited solutions. Mechanistically this is important because in the former solutions dissolution is thought to be restricted by an oxide, typical polarization curves exhibiting an active to passive transition (Fig. 15a). However, in liquid a m m o n i a le concentrated chloride is and CO containing solutions s8 dissolution is thought to be hindered--at leastinitially--bya thin adsorbed layer whose formation is reversible. In this case polarization curves do not
The fractography of stress corrosion cracking in carbon steels
475
exhibit an active to passive transition. However, in chlorides and CO inhibited solutions the anodic polarization curves exhibit an adsorption induced inflection as shown schematically in Fig. 15(b). Furthermore, the crack path in the oxide forming solutions is structurally preexisting, attack on unstressed specimens is structurally dependent and in nitrates intergranular corrosion can occur which is quantitatively related to susceptibility. However, in the solutions where adsorption effects are thought to be important cracks initiate at slip steps. Thus the path is in a sense strain generated. The obvious suggestion is that in the oxide forming solutions thick corrosion products form whose formation is structurally dependent. ~s That is they form last at grain boundaries which also dissolve more rapidly. Subsequently the oxide either never manages to form at the crack tip or is continually ruptured. In the adsorbing solutions a thin layer forms rather quickly and its formation is possibly less structurally sensitive. Subsequently when slip ruptures this layer localized dissolution can occur at the emergent slip step.aS Such a sharp division between oxide-forming solutions and solutions where adsorption inhibits dissolution is made for the sake of convenience of discussion and more correctly represents extreme cases. This is readily apparent when we consider some transgranular cracking does occur in hydroxides and some intergranular cracking was found in chlorides. Furthermore there is still considerable discussion as to whether an adsorbed layer precedes the formation of a passive film. a4
2 Practical flnplications As others have indicated a~ the failure of annealed mild steels by SCC can no longer be regarded as occurring exclusively intergranularly, thus the existence of intergranular cracking cannot be used as a diagnostic criteria for SCC. This means that transgranular SCC must be distinguished from transgranular fatigue or corrosion fatigue. Where fractographic detail is observed by corrosion products and even second order characteristics, such as "clamshell markings" on fatigue failures cannot be observed, reliance on metallographic examination and the failed components history must be used. This is not always decisive. Failures of quenched mild steels involve similar
5CC
SCC
~
CC
(b) Log C.D.
Log G.D
FIo. 15. (a) Schematic anodic polarization curve exhibiting an active to passive transition and the potential ranges which are most likely to promote SCC, (b) Schematic anodic polarization curve exhibiting adsorption induced inhibition and the potential range most likely to promote SCC.
476
B. POULSOr~
fractographic diagnostic problems. Thus intergranular cracking, along prior austenite grain boundaries, can result from SCC, hydrogen embrittlement, quench cracking and temper embrittlement. Another important point is that the crack path depends markedly on the metallurgical structure and differences in the parent metal, the heat affected zone and the weld metal are to be expected. Furthermore this effect of structure on the crack path depends on the solution and the relative susceptibilities of each zone is unlikely to be invariant. Thus in liquid ammonia practical experience 16 is that cracks tend to occur in the weld metal, while in hydroxides the parent metal or the H.A.Z. is preferred. If a failure of a mild steel component is intergranular and SCC is diagnosed but environmental conditions responsible are unknown, extensive carbide attack is reason to suspect hydroxides. The potential dependence of cracking is another diagnostic method which has been used in service failure analysis. 36 It must be pointed out that the number of solutions known to cause SCC is increasing. It must not become an obsession to prove a SCC failure resulted from the presence of nitrates or hydroxides.37 FUTURE WORK Much of the fractographicwork on SCC has been ambiguous. That is, the features observed could be interpreted in terms of a number of mechanistic ideas. A noteworthy exception to this is Scully's work on SCC in Ti alloys. 2,~ Real progress in this field will probably result from two separate approaches. The first, the application of a wide range of other techniques to examine the fracture surface. These would include dislocation etching, channeling patterns, X-ray broadening, electron diffraction, ion spluttering or laser probe-mass spectroscopy, auger spectroscopy, ESCA, transmission electron microscopy (after thinning from one side) and electrochemical techniques designed to assess the protectiveness of any film formation on the fracture surface --which the physical techniques cannot do. The second approach involves changing a variable during crack propagation and observing the effect on crack propagation and the fracture surface. From a practical aspect much work needs to be done on the fractography of real materials under service conditions. For example, welded or coldworked alloy steels which after failure have been left for a period prior to examination. CONCLUSIONS 1. In nearly all cases the fracture surfaces of carbon steels which had failed by SCC were clean enough for fractographic details to be observed. 2. The crack path was found to depend on--heat treatment, environment, potential and alloying element. 3. For annealed steels in solutions (nitrate, hydroxide and carbonate) which promoted thick oxide formation, cracking was predominantly intergranular. In chloride and liquid ammonia where thin films or an adsorbed layer are thought to be important cracking was predominantly transgranular. 4. For quenched steels cracking was transgranular and intergranular along the prior austenite grain boundaries. In oxide-forming solutions the transgranular cracking was along martensite lath boundaries, while in chloride and liquid ammonia it was not apparently related to any structural feature. 5. The available fractographic evidence appears inconsistent with either hydrogen
The fractography of stress corrosion cracking in carbon steels
477
e m b r i t t l e m e n t o r brittle films p l a y i n g a role in c r a c k p r o p a g a t i o n . It is, h o w e v e r , c o n s i s t e n t w i t h c r a c k p r o p a g a t i o n by a n o d i c d i s s o l u t i o n p o s s i b l y a i d e d by t e a r i n g . Acknowledgements The author wishes to thank: Prof. R. N. Parkins and Civ.-Eng. Hans Arup for helpful discussions and valuable criticisms, and the staffof the S.E.M.s in the Department of Chemistry, University of Newcastle upon Tyne, and the Danish Technical Institute at Tastrup, for their help. REFERENCES 1. N. A. NIELSEN,Fundamental Aspects of Ctress Corrosion Cracking (Eds. R. W. STAEHLEel al.), p. 308. NACE, Houston, Texas (1967). 2. J. C. SCULLY, The Theory of Stress Corrosion Cracking in Alloys, p. 127. NATO, Brussels (1971). 3. F. E. WATKINSONand J. C. SCULLY, Corros. Sci. 12, 905 (1972). 4. W. R. MIDDLE'tON and R. N. PARKINS, Corrosion 28, 881 (1972). 5. D. T. POWELLand J. C. SCULLV, Corrosion 24, 151 (1968). 6. R. J. H. WANHILL, Corrosion 29, 435 (1973). 7. G. DEAr% Ph.D. Thesis, University of Newcastle upon Tyne (1971). 8. V. J. CALANOELOand M. S. FERC,trsoN, Corrosion 25, 509 (1969). 9. J. C. SCULLY, Paper presented at the International Conference on Stress Corrosion Cracking and Hydrogen Embrittlement of Iron Base Alloys. Firminy, France, June 1973, to be published by NACE. 10. E. N. PUGH, J. V. CRAIG and A. J. SEDRIKS,Fundamental Aspects of Stress Corrosion Cracking (Eds. R. W. S'rAErtLEet al.), p. 118. NACE, Houston, Texas (1967). 11. G. M. SPARKSand J. C. SCtrLLY, Corros. Sci. 11, 641 (1971). 12. R. D. BARERand B. F. PE'rEF,S, Why Metals Fail. Gordon and Breach, New York (1970). 13. J. E. TRU'MAN,Paper presented at the International Conference on Stress Corrosion Cracking and Hydrogen Embrittlement of Iron Base Alloys. Firminy, France, June 1973, to be published by NACE. 14. B. Pout.sos, Ph.D. Thesis, University of Newcastle upon Tyne (1972). 15. B. PouLsos and R. N. PARKINS, Corrosion 29, 414 (1973). 16. B. POULSONand H. ARtrp, In preparation. 17. R. N. PAR.KINS,B. PouLSos and P. SLATTERY,In preparation. 18. R. N. PARKINS, The Theory of Stress Corrosion Cracking in Alloys (Ed. J. C. SCULLY),p. 167. NATO, Brussels (1971). 19. C. J. CROS, J. H. PAYERand R. W. STAEHLE,Corrosion 27, l (1971). 20. J. C. WILLIAMS,Discussion to ref. I. 21. C. D. BEACHEM,Electron Fractography, STP436, p. 59. ASTM, Philadelphia (1968). 22. N. A. NIELSES, Corrosion 27, 173 (1971). 23. F. P. A. ROBIr'~SONand L. G. NEE,Proc. 2ndbzternational Congress on Metallic Corrosion, p. 171. NACE, Houston 0969). 24. B. E. WILDE, Corrosion 27, 326 (1971). 25. F. J. RADO and D. H. OERTLE,N A C E Annual Conference. Preprint, Chicago, March 1971. 26. H. H. UHLlC~,Physical Metallurgy of Stress Corrosion Fracture (Ed. R. RHODIN), p. 258. Interscience, New York (1959). 27. C. F. BARa'H, E. A. STEIGERWALDand A. R. TPaASO, Corrosion 25, 353 (1969). 28. S. S. BIRLEYand D. TROMANS,Corrosion 27, 63 (1971). 29. C. Et~ELEANU,JlSI 143, 140 (1953). 30. H. H. UHLIG, Fundamental Aspects of Stress Corrosion Crackhtg (Ed. R. W. S'rAEHLEet al.), p. 86. NACE, Houston (1969). 31. D. J. LEEs, Corros. Sci. 12, 811 (1972). 32. E. N. PtJGH, Discussion at International Conference on Stress Corrosion Cracking and Hydrogen Embrittlement of Iron Base Alloys. Firminy, France, June 1973, to be published by NACE. 33. K. E. HEUSLERand G. H. CARTLEDGE,J. electrochem. Soc. 108, 733 (1961). 34. J. O'M. BOCKRISand P. K. StraRAMASVAS, Corros. Sci. 10, 435 (1970). 35. M. KOWAKAand S. NAGA'rA, Corrosiot, 24, 422 (1968). 36. B. Potrt.sos, L. C. HENRIKSENand H. ARUP, Br. Corros. J. In press. 37. R. N. PARrdNS, Paper presented at the International Conference on Stress Corrosion Cracking and Hydrogen.Embrittlement of Iron Base Alloys. Firminy, France, June 1973, to be published by NACE. 38. A. BRows, J. T. HAmUSOS and R. WILrdNS, ibid.