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Kinetic and fractographic study of the stress corrosion cracking of Austenitic stainless steels Russell, Alan James 1977

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C P . < A KINETIC AND FRACTOGRAPHIC STUDY OF THE STRESS CORROSION CRACKING OF AUSTENITIC STAINLESS STEELS by ALAN JAMES RUSSELL B.Sc., University of Glasgow, 1972 A THESIS SUBMITTED IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF MASTER OF SCIENCE We accept this thesis as conforming to the required standard THE UNIVERSITY OF BRITISH COLUMBIA January, 1977 f c ) Alan James Russell, 1977 in THE DEPARTMENT OF METALLURGY In p r e s e n t i n g t h i s t h e s i s i n p a r t i a l f u l f i l m e n t o f the r e q u i r e m e n t s f o r an advanced d e g r e e a t t he U n i v e r s i t y o f B r i t i s h C o l u m b i a , I a g r e e t h a t t he L i b r a r y s h a l l make i t f r e e l y a v a i l a b l e f o r r e f e r e n c e and s t u d y . I f u r t h e r a g r e e t h a t p e r m i s s i o n f o r e x t e n s i v e c o p y i n g o f t h i s t h e s i s f o r s c h o l a r l y p u r p o s e s may be g r a n t e d by the Head o f my Depar tment o r by h i s r e p r e s e n t a t i v e s . I t i s u n d e r s t o o d t h a t c o p y i n g o r p u b l i c a t i o n o f t h i s t h e s i s f o r f i n a n c i a l g a i n s h a l l no t be a l l o w e d w i t h o u t my w r i t t e n p e r m i s s i o n . Depar tment o f M e t a l l u r g y The U n i v e r s i t y o f B r i t i s h C o l u m b i a V a n c o u v e r 8, Canada Date F e b r u a r y 21, 1977 ABSTRACT A variation of the double cantilever beam specimen has been calibrated and used to study the propagation of stress corrosion cracks as a function of stress intensity ln Jl6 and 310 stainless steels, and a TRIP steel exposed to hot aqueous magnesium chloride solutions. The effects of cold work, tem-perature and applied potential on both the fractography and cracking rates have been examined. The effects of cold work and crack path on crack branching were also investigated. Bpth stress Intensity dependent (Region I) and stress intensity independent (Region II) cracking were observed. Region II having apparent activation energies from 15.1 kcal/g.mole to 18.1 kcal/g.mole. The crack velocities of Z% cold rolled 316 were found to be independent of applied potential pver a range of more than 5° m V in Region I and 75mV in Region II. In the same material the crack path changed from solely transgranular at low stress intensities and noble potentials to more than 86% intergranular at high stress intensities and active potentials. The topography of the transgranular fracture was similar to that observed by others except in the case of the TRIP steel where nodular features were observed. These observations have been discussed with respect to mechanisms involving the follpwingj (i) electrochemical diss-olution, (ii) absorption of hydrogen and ( i i i ) adsorption of a damaging species. Of these, an adsorption assisted process is most compatible with the observations. Qualitatively the adsorbed species are envisioned as modifying the behaviour of the surface atoms at the crack tip. -11-ACKNOWLEDGEMENT I would like to thank Dr. D. Tromans for his readily available help and encouragement. I am also grateful to the faculty, staff and fellow students of the Department of Metallurgy for a l l the useful advice given to me regarding this research. Thanks are also due to Mrs. Anne Russell for financial support, patience and understanding. The financial assistance provided by the National Research Council (grant number A-^ 534) is also acknowledged. - i i i -TABLE OF CONTENTS PAGE 1. INTRODUCTION ,,, 1 1.1 General Description , 1 1.2 Fracture Mechanics Approach to 3.C.C.. 2 1.3 The Chloride Cracking o f Fe-Cr-Ni Alloys... 7 1.4 ahbice of Materials and Scope of the Present Work 10 2 • EXPERIMENTAL. ,•»•»,•••••,•*••••••••»•..«••.•»••••••.••••»•••*•»••. 12 2.1 Specimen Geometry, 12 2.2 Stress Intensity Calibration. ,..., 12 2.3 Specimen Thickness 18 2.4 Advantages and Limitations of the Specimen Geometry......... 19 2.5 Materials and Specimen Preparation 20 2.6 Apparatus and Testing Procedure............................. 22 .3. • RESULTS.., . 28 3.1 Stress Corrosion Crack Morphology 28 3.1.1 316 Stainless Steel..... 28 3.1.2 TRIP Steel.. 31 3.1.3 310 Stainless Steel 31 3.2 Considerations of Errors............. 31 3.2.1 Errors in Crack Length Measurements 31 3.2.2 Errors due to the Cracks Growing Downwards 33 3.3 Kinetic Data...... 33 3.3.1 316 Stainless Steel 33 3.3.1.1 Effect of Stress Intensity on Crack Velocity 33 3.3.1.2 Effect of Plastic Deformation, . 3 4 -iv-PAGfi 3.3.1 .3 Effect of Temperature. 34 3.3.1 .4 Effect of Applied Potential 40 3.3.2 TRIP Steel 43 3 .3 .2 .1 Effect of Martensite on the Region II Velocity.. 43 3 .3 .2 .2 Effect of Temperature on the Region II Velocities.... . . . . . 4 3 3.3/3 310 Stainless Steel..... .' 43 3 .3 .3 .1 Effect of Stress Intensity on the Crack Velocity..... 43 3 .3 .3 .2 Effect of Temperature on the Crack Velocity..... 50 3.4 Corrosion Potentials..,. 50 3.4.1 Effect of Stress Intensity on the Corrosion Potential. 50 3.4.2 Effect of Composition and Processing Treatment on the Corrosion Potential .. 54 3 .4 .3 Effect of Temperature on the Corrosion Potential..... 54 3.5 Frac tography 54 3.5.1 316 Stainless Steel 54 3.5.1.1 Effect of Plastic Deformation 54 3.5.1.2 Effects of Electrochemical Potential and Stress Intensity 57 3.5.1.3 Effect of Temperature 68 3.5.2 TRIP Steel 74 3 .5 .2 . 1 Comparison with the Fractography of the 316 Stainless Steel..., 74 3.5.2.2 Effects of Cold Working and the Presence of Martensite 74 - V -PAQE 3.5.2.3 Effect of Stress Intensity 78 3.5.2.4 Effect of Temperature 78 3.5.3 310 Stainless Steel 78 3.5.3.1 General Description 78 4. DISCUSSION 83 4.1 Crack Branching. 83 4.2 Kinetics 85 4.2.1 Region I 85 4.2.2 Region II 86 4.2.3 Crack Mechanisms. 87 4.2.3.1 General Considerations 88 4.2.3.2 Dissolution 91 4.2.3.3 Adsorption 93 4.2.3.4 Absorption of Hydrogen... 96 4.3 Corrosion Potentials 97 4.4 Fractography » •• 99 4.4.1 General 99 4.4.2 Effect of Applied Potential 100 4.4.3 Effect of Stress Intensity..... 101 4.4.4 Effect of Composition 103 4.4.5 Effect of Martensite ...104 4.5 Assessment of S.C.C. Models 105 5. CONCLUSION........ 108 5.1 Conclusions 108 5.2 Suggestions for Future Work 109 BIBLIOGRAPHY HO APPENDIX 1.... .....114 -vi-LIST OF TABLES TABLE PAGE I Alloy Compositions 21 II Mechanical Properties of the Various Alloys and Processing Treatments (Tensile data measured perpendicular to the rolling direction at 150 C.) 23 III Variation with Temperature of Region II Velocities for 2% and 50% Cold Rolled 316 Stainless Steel............. 39 IV Variation with Temperature of Region II Velocities for 25% Cold Rolled Austenitic and Partially Martensitlc TRIP Steel 47 •V Variation with Temperature of Initial and Subsequent Constant Velocities for 50% Cold Rolled 310 Stainless Steel 51 VI Corrosion Potentials of the Various Alloys and Processing Treatments ... 56 Al Errors due to the Downward Growth of the Cracks 11$ - v i i -LIST OF FIGURES FIGURE PAGE 1 Typical relationship between stress corrosion cracking velocity and stress intensity showing Region I, II and III and K^gGC 5 2 Specimens used in the present work and by others to measure crack velocities as a function of stress intensity...... , • 13 3 Double cantilever beam specimen showing side cracking. 8x. Fatigue crack at tip of notch. Other cracks are S.C.G 14 4 Compliance of the specimen used in the present work as a function of crack length 16 5 Stress intensity calibrations for the specimen used in the present work and for a three point bend specimen of similar size... 17 6 Experimental arrangement showing S.C.C. cell, reflux condenser and load stabilizing spring mounted in tensile testing machine ..24 7 Crack morphology for cracking parallel to the rolling direction in Z% cold rolled 316 stainless steel. (a.) Side view. 7x. (b) Fracture surface. 4,8x. (c) End view at 15 and 30MPa.m¥. 23x 29 8 Crack morphologies in JlS stainless steel. 12x. (a) Cracking perpendicular to the rolling direction. (b) Cracking in the annealed alloy. (The crack has been opened to facilitate observation of cracking.) 30 - v i i i -FIGURE PAGE 9 Crack morphology in the 310 and TRIP steels. (a) TRIP - 25% cold rolled at 200°C. Cracking parallel to the rolling direction. 7x, (b) 310 - 50% cold rolled. Cracking parallel to the rolling direction. 17x. , 32 10 Variation of crack length and load with time for a specimen of 25% cold rolled Jl6 stainless steel having an i n i t i a l load of 500N. MgClg solution boiling at 154°C 35 11 Variation of crack length and load with time for a specimen of 25% cold rolled 31° stainless steel at 154°C. having an i n i t i a l load of 1200N. MgCl2 solution boiling at 154°C 36 12 Effect of stress intensity on crack velocity for 25% cold rolled JlS stainless steel exposed to boiling MgCl2 solution at 154°C 37 13 Effect of stress intensity on crack velocity for 50% cold rolled 316 stainless steel exposed to boiling MgCl2 solution at 154°C 38 14 Arrhenius plot of the Region II velocities of cold rolled Jl6 stainless steel cracking in a h5% MgCl2 solution 41 15 Effect of applied potential and temperature on the Region I cracking of 25% cold rolled 316 stainless steel exposed to a ^ 5% MgCl2 solution 42 16 Effect of cathodic polarization on the Region II velocity of 25% cold rolled 316 stainless steel cracking in a MgCl? solution boiling at 154°C. 44 -ix-FIGURE PAGE 17 Effect of anodic polarization on the Region II velocity of 2% cold rolled Jlh stainless steel cracking in a MgClg solution boiling at 154°G 45 18 Effect of stress Intensity on crack velocity for 25% cold rolled TRIP steel exposed to MgCl2 solution boiling at 154°C 46 19 Arrhenius plot of the Region II velocities of 25% cold rolled TRIP steel cracking in a l±5% MgCl2 solution 48 20 Variation of crack length, load and stress intensity with time for a specimen of 50% cold rolled 310 stainless steel exposed to boiling MgCl^ solution at 154°C 40 21 Arrhenius plot of the crack velocities of 50% cold rolled 310 stainless steel cracking in a k$% MgCl2 solution 52 22 Variation of corrosion potential with time as a function of applied load. 25% cold rolled Jl£ stainless steel in boiling MgCl2 solution at 154°G 53 23 Effect of stress intensity on the corrosion potential of 25% cold rolled 316 stainless steel cracking in boiling MgCl2 solution at 154°G... 55 24 Effect of temperature on the corrosion potential of 25% cold rolled Jl6 stainless steel cracking in ^5% MgCig s° 1- u't l o n a"t a stress intensity of 25MPa.m2. 55 25 Effect of cold work on the fractography of 316 o stainless steel at 154 G. 400x. A (a) Annealed, K 1 = lOMPa.m2 (b) 25% cold rolled K± = 20MPa.m2 (c) 50% cold rolled Kn = 20MPa.mJ 58 - X -FIGURE PAGE 26 Matching fractographs from opposite sides of the crack in 2% cold, rolled 316 stainless steel. 400x. 59 27 Matching fractographs from opposite sides of the crack in 2$% cold rolled 316 stainless steel. lOOOx 60 28 Effects of applied potential and stress intensity on the amount and distribution of intergranular cracking. 2% cold rolled 316 stainless steel exposed to boiling MgClg solution at 154°C 61 29 Examples of fractographs of 2% cold rolled^ 316 stainless steel at 154°G. = llMPa.m2, E = -0.25V (S.C.E.). 400x. (a) Edge of crack (b) Centre of crack 62 30 Examples of fractographs of 23% cold rolled^ 316 stainless steel at 154°C. Kx = 22MPa.m% E = -0.25V (S.C.E.). 400x. (a) Edge of crack (b) Centre of crack 63 31 Examples of fractographs of 25$ cold rolled^ 316 stainless steel at 154°C. = llMPa.m^, E - -0.3V (S.C.E.). 400x. (a) Edge of crack (b) Centre of crack 64 32 Examples of fractographs of 25$ cold rolled^ 316 stainless steel at 154°C. K 1 = 22MPa.m% E = -0.3V (S.C.E.). 400x. (a) Edge of crack (b) Centre of crack 65 -xi-FIGURE PAGE 33 Examples of fractographs of 2% cold rolled 316 stainless steel at 154°C. Kn = 33MPa.mT. Centre of crack. 400x. (a) E = -O.325V (S.C.E.) (b) E - -0.35V (S.C.E.) 66 34 Example of fractograph of 25$ cold rolled 316 stainless steel cracking without an applied load. 800x 6? 35 Example of fractograph of 50% cold rolled 316 stainless steel showing small area of ductile failure. = 70MPa.m*. 400x 69 36 Distribution of ductile fracture in a specimen of 50% cold rolled Jl6 stainless steel exposed to boiling MgClg solution at 154°C. (a) As a function of position across the specimen. (b) As a function of stress intensity 70 37 Examples of precipitates formed on the fracture surface during cathodic polarization. (a) E = -0.325V (S.C.E.). 800x. (b) E = -0.35V (S.C.E.). 400x... 71 38 Effect of temperature on the fractography of 50% cold rolled 316 stainless steel. lOOOx. (a) T - 154°C, Kx - 27MPa.m2 (b) T = 116°C, Kx - 27MPa.m2 72 39 Effect of temperature on the fractography of 25$ cold rolled 316 stainless^steel. 400x. (a) T - 154°C, Kx = 27MPa.m2 (b) T - 116°C, K. = 27MPa.m2 73 - x l i -FIGURE PAGE 40 Effect of cold work on the fractography of the TRIP steel. 400x. (a) Annealed TRIP, K± - lOMPa.m^, T - 154°C. (b) 25% cold rolled (at 200°C.), - 9MPa.m^ , T - 154°C 75 41 Effect of martensite on the fractography of 25% cold rolled TRIP steel. lOOOx. (a) Fully austenitic, 1^ - 22MPa.m^ , T - l l6°C. (b) Containing martensite, * 22MPa.m^ , T - l l 6°C 76 42 Matching fractographs from opposite sides of the crack in 25% cold rolled (at 200GC.) TRIP steel. 800x 77 43 Effect of stress intensity on the fractography of 25% cold rolled (at 200°C.) TRIP steel at open c i r c u i t potentials, lOOOx. (a) No applied load, T - 154°C. (b) K x - 8.5MPa.m^, T - 154°C. (c) K x •- 40MPa.m2, T - 154°C 79 44 Effect of temperature on the fractography of the 25% cold rolled (at 200°C.) TRIP steel. lOOOx. (a) T ~ 154°C., K, - 8.2MPa.m^ " _ 1 i (b) T - 119 C , K x - 6.6MPa.mT 80 45 Effects of cold work and stress intensity on the fractography of 310 stainless steel. 400x. (a) Annealed, - 20MPa.m^ , T - 154°C. (b) 50% cold rolled, ^  - 25MPa.m^ , T - 154°C. (c) 50% cold rolled, «= 50MPa.m2\ T - 154°C 81 46 Effects of temperature and applied potential on the fractography of 50% cold rolled 310 stainless steel. 400x. (a) T - 120°C, - 25MPa.m^ , E » -0.27V(S.C.E.) 82 x 4 (b) T - 154°C., K x - 25MPa.m% E - -0.325V(S.C.E.) - x i i i -FIGURE PAGE 47 V-K1 plot for 7039-T64 aluminum alloy in 5M aqueous solution showing two stress intensity independent regions and the effect of tem-perature (after Hyatt and Speidel [5]) 90 48 Relationship between the heat of adsorption and the activation energy for chemisorption (after Glasstone, Laidler and Eyring [56]) 95 49 Schematic Evens diagram illustrating the effect of stress intensity on the corrosion potential 98 50 Schematic representation of the effect of adsorption at the crack tip.. 107 Al Effect of the downward growth of the cracks on the measurements of V and K. 115 -1-1. INTRODUCTION 1.1 General Description Stress corrosion cracking (S.C.C.) is a time dependent, low energy fracture process which may occur in materials subject to the conjoint action of a tensile stress and a specific corrosive environment. Some common examples of susceptible systems are» (i) Copper alloys In the presence of ammonia and water. (ii) Aluminum alloys in aqueous solutions containing chloride ions, ( i i i ) Mild steel in caustic solutions. (iv) Austenitic stainless steels in aqueous solutions containing chloride ions. S.C.C. is a particularly annoying form of corrosion for several reasonst (i) S.C.C. reduces the usefulness of materials having especially desirable properties such as high strength to weight ratios, e.g., 7000 series aluminum alloys, Ti-Al-Mo alloys and 4340 type steels, or excellent corrosion resistance (i.e., to general dissolution), e.g.» 300 series stainless steels. (ii) Cracking can take place with l i t t l e or no applied stress (the residual stress from machining, welding or quenching being sufficient), and with only a few parts per million of aggressive species present. ( i i i ) Cracks may propagate rapidly once they have been initiated, -2-making them difficult to detect prior to failure. S.C.C. has been responsible for a large number of service failures during the last 35 years and has been estimated to have cost the chemical industry in the U.S.A. alone over $30 million annually in maintenance and replacement of failed equipment [l]. The knowledge gleaned from failure analysis, as well as from experimental work, has enabled designers, manufacturers and corrosion engineers to reduce considerably the incidence of failure due to S.C.C. However, there is as yet no theory that can predict which materials will be susceptible in a particular environment, or even how minor changes in an environment will affect the resistance of a given material to S.C.C. 1.2 Fracture Mechanics Approach to S.C.C. Prior to 1965, most tests of susceptibility to S.C.C. were performed on smooth specimens (e.g., U-bend and tensile specimens), the time to failure being the measure of susceptibility. While these tests have provided a great deal of useful, practical information, from the point of view of improving understanding of the mechanisims involved in S.C.C. they have several dis-advantages, viz: (i) The time to failure usually includes the crack initiation time, which, especially in the case of thin specimens, may be a large fraction of the total time. (ii) The stress is not well defined (usually given as the net section stress prior to cracking), and varies in an unknown manner as the specimen cracks. -3-( i i i ) In the case of U-bend specimens and constant load tensile specimens, overload failure will occur, with less S.C.C. having taken place, i f either the i n i t i a l stress is increased or the ultimate tensile strength is reduced. The fracture mechanics approach to S.C.C. overcomes these disadvantages by studying the propagation of single planar cracks. The stress around the crack tip is measured by the stress intensity factor, designated 'K*. When the crack opens normal to the crack plane, as is the case in S.C.C, i t is called opening mode I, and the stress intensity factor is written as K^ . For a straight through centre crack of length 2a, or an edge crack of length a, with a straight crack front in an infinite plate having a tensile stress c normal to and distant from the crack, the stress intensity factor is given by \_2~\t 1. = O ( -na ) 2 Thus, has the dimensions of STRESS x (LENGTH)2. In practice, for small test specimens, i t is given by • % -1 . Y(a) where P is the applied load B is the specimen thickness and Y(a) is a function of the crack length and the specimen geometry. The measure of susceptibility to cracking is the crack propagation rate or crack velocity (V), and, since i t can vary over several orders of magnitude, data is normally displayed graphically as log V against K,. The fracture mechanics -4-approach has been used to study the S.C.C. of the following materialsi soda lime and other glasses £3»*0> various aluminum alloys [5 »6 ] , titanium alloys [7 ] , high strength steels [8], 70-30 brass [9] and an Fe-Mn-Cr austenitic steel [lO]. All these materials have V-K^ plots similar to that shown in Fig. 1, although not a l l exhibited every feature. With reference to Fig. 1, in Region I the velocity increases with stress intensity, the variation often being approximated by a straight line on the log V vs.K^ graph. At low velocities the curve may turn to be parallel to the ordinate, in which case a threshold stress intensity, K-J^QQ t can be defined. If cracking continues to lower and lower stress intensities an arbitrary K^ CJQQ i s sometimes defined as the stress Intensity resulting in some arbitrarily small velocity. In Region II the crack velocity is constant and thus independent of stress intensity. In Region III the crack velocity increases rapidly with stress intensity and is independent of the environment. The transition from region to region may not be as sharp as shown by the solid curve in Fig. 1, but instead may be quite gradual as shown by the broken line. The reason for this may be that any local variation in stress intensity along the crack front will smooth out any sudden changes in dV/dK^ . Possible causes of variations could be changes in stress state at the edges of the specimen, local changes in crack length and variations in the residual stress from grain to grain in cold worked material. -5-C9 O ^ I S C C S T R E S S I N T E N S I T Y FIGURE 1 Typical relationship between stress corrosion cracking velocity and stress intensity showing Regions I, II and III and K. g c c. - 6 -Most of this type of research has been carried out by Industry or other organisations concerned with the practical problems of S.C.C. Consequently much of the work has dealt with alloy development and measurement of K - ^ Q values, e.g., Boeing Airplane Co. on aluminum and titanium alloys, U.S. Naval Research Laboratory on high strength steels, and Brown Boveri Research Centre on Fe-Mn-Cr alloys. From a mechanistic point of view, however, the real merit of the fracture mechanics approach to S.C.C. is in its ability to distinguish stress dependent mechanisms (Region I) from stress independent phenomena (Region I i ) . This is best illustrated by an example. Speidel obtained activation energies of 27 kcal/g.mole and 3.8 kcal/g.mole for Regions I and II respectively, for a 7079-T6 aluminum alloy in 3M aqueous potassium iodide solution £6]. Further, by determining the effects of iodide concentration and solution viscosity on the Region II velocity, he was able to show that the stress independent cracking rate was controlled by mass transport in the solution f i l l i n g the crack. Speidel did not come to any conclusion about the mechanism for Region I. On the other hand, Helfrich obtained activation energies from time to failure tests of a 7039-T6 alloy in a IN sodium chloride solution [ l l ] . His values ranged from 20.4 kcal/g.mole at 15 ksi to 16.9 kcal/g.mole at 55 ksi. These values together with a calculated activation volume of 28.5 c.c./g.mole led him to the conclusion that the rate controlling process for inter--7-granular stress corrosion of the 7039 aluminum alloy might he based on stress and thermally activated dissolution of MgZn2 at the grain boundaries. It is interesting to note that Helfrich's activation energies l i e between those of Speidel for Regions I and II and that the decrease with increasing stress could be explained by the fact that at higher stresses more time would be spent in Region II (lower activation energy). 1.3 The Chloride Cracking of Fe-Cr-Nl Alloys The f i r s t reported industrial failure of austenitic stainless steel, that was attributed to S.C.C, took place in 19^ 0 [12]. Since then i t has been well established that most failures occur in the presence of chloride ions at temperatures above about 60°C £l3rl4]. However, other environments such as sulphuric acid [15] 1 caustic solutions [l6] and low pH chloride solutions at room temperature [17] have also been shown to cause cracking. During the last 30 years much effort has gone into trying to understand the mechanisms involved in the chloride cracking of austenitic Fe-Cr-Ni alloys. Many time to failure tests have been carried out, as well as numerous phenomenological experiments, and, on the basis of these, several theories have been proposed. Rather than trying to summarise a l l of this work, only two of the more important aspects will be considered at this stage. Full details of the previously published studies are presented in excellemt reviews of the subject by Latanision and Staehle -8-[18], Nielsen [ l ] and Staehle [19]. Nickel extends the time to failure [20,2l] but as yet its precise role has not been satisfactorily explained. Nickel is more noble than either iron or chromium and according to Latanision and Staehle [ l 8 ] the crack sides might become enriched in nickel. Freshly exposed metal at the crack tip is then anodic to the crack sides and dissolves. At higher nickel contents the driving force for this dissolution is reduced and so cracking is slower. There is, however, no good evidence for nickel enrichment on the crack sides and i t has now been established that the crack sides are covered with a passive oxide film [22,23]. Nickel has been shown to raise the stacking fault energy (S.F.E.) of austenitic Fe-Cr-Ni alloys [24], Swann and Nutting [25] demonstrated preferential dissolution along stacking faults in copper-zinc and copper-aluminum alloys. It was later proposed that the low S.F.E. led to co-planar dislocation arrangements which, in turn, result in coarse slip [26]. This coarse slip has a greater chance of breaking the oxide film on the surface than i t would i f the slip was finer. However, co-planar dis-locations have been observed in the higher nickel alloys in spite of their higher S.F.E. [27]. Thus, while nickel will increase the S.F.E., i t is not clear exactly how the S.F.E. affects the susceptibility to S.C.C. It is possible that nickel might affect the properties of the oxide film formed on the stress corrosion fracture surface. Properties such as the conductivity and ductility -9-of the film are important ln film rupture models of S.C.C. [28]. The structure and composition of the films formed on the stress corrosion fracture surface of 304, 316 and 310 alloys in MgCl2 solution have been studied by Nikiforuk [29]. In a l l cases he found the structure to be that of a chromium enriched spinel and that the compositions were similar. Thus, i f film rupture does occur, then the effect of nickel is more likely to be on some other property such as the slip step height or the dissolution rate and not on its effect on the passive film. The role of hydrogen in the S.C.C. of austenitic stainless steels is uncertain. The pH within the crack has been measured to be 1.0 or less in the case of 304 alloy cracking in boiling MgCl2 solution a,t 154°C. [ 3 0 ] . This low pH results frpm hydrolysis of the corrosion products [ 3 l ] and means that proton reduction, H + + e~5=? H , ads is possible within the crack. That this is in fact the cathodic reaction has been confirmed by the observation of hydrogen bubbles emanating from stress corrosion cracks [ l , 3 2 ] . Holzworth and Louthan [33] and Vaughn et al. [34] showed that cathodic charging with hydrogen resulted in transformation of 304L alloy to € and a'phases. Birley and Tromans [35] detected a'martensite on the stress corrosion fracture surface of 304L stainless steel and showed that the topography of the fracture surface was consistent with the presence of a'martensite laths. Since a does not normally form under tensile stresses at the temperature of the test, they suggested that hydrogen absorbed -10-into the lattice might be responsible for the transformation. It is not clear, however, whether the crack propagates through the a' phase by a hydrogen embrittlement mechanism or by dis-solution. On the other hand, there is no evidence for martensite involvement in the S.C.C, of the higher nickel alloys such as 310 and Inconel 600 [36]. Other mechanisms by which absorbed hydrogen might promote cracking have been proposed [l9»37]i but as yet there is no consensus as to whether or not absorbed hydrogen is a necessary feature of the S.C.C. of austenitic Fe-Cr-Ni alloys. 1.4 Choice of Materials and Scope of the Present Work Three alloys were selected for this investigation, namely AISI 316, AISI 310 and a TRIP* steel v The 316 and 310 alloys are both known to crack in hot chloride solutions. However, the 310 with 9# more nickel takes approximately ten times as long to f a i l . The TRIP steel was chosen because of its greater tendency to form martensite on deformation. The cracking of the TRIP steel was examined in the fully austenitic condition as well as partially transformed to martensite, but was not studied in the TRIP condition (i.e., both work and precipitation hardened by deforming 80% at400°C.) As far as is known there is no TRIP is an acronym for Transformation Induced Plasticity and was proposed by Zackay [38]. TRIP steels make use of the strain induced martensite transformation to delay the onset of necking and thereby to increase the ductility. -11-published data on the susceptibility to S.C.C. of the TRIP steel, although there is some information available on the hydrogen embrittlement of these steels [39]. Cracking of a l l the steels was examined in both the fully annealed and cold worked conditions, however, quantitative results were only obtained from the cold worked specimens. The environment used for cracking was a concentrated solution of magnesium chloride in water having a boiling point of 154°C, Although such a medium is not common industrially, the cracking susceptibilities of alloys determined in magnesium chloride compare well with those obtained from long term tests and from industrial experience [4o], This environment also has the advantage that reasonably fast cracking rates can be achieved without the need to go to higher than atmospheric pressure. The other experimental variables studied were the tem-perature, the electrochemical potential and the amount of prior deformation to the alloys. The effects of these variables on the topography of the stress corrosion fracture surface were also examined. It was hoped that by applying the fracture mechanics approach to S.C.C., to the chloride cracking of Fe-Cr-Ni alloys, that new insight could be gained into the cracking mechanism. -12-2. EXPERIMENTAL 2.1 Specimen Geometry The type of specimen most often used for measuring stress corrosion crack velocities as a function of stress intensity is the double cantilever beam specimen (D.C.B.) \_5~}. See Fig. 2(a). When wedge loaded (i.e.* constant deflection) i t has several important advantages: (i) It is compact while allowing long cracks to be studied, (ii) It gives a large range of stress intensity, decreasing from the i n i t i a l value to that at which the crack stops. (Kigrjrj) • ( i i i ) Initiation times are short. Thus, accurate data can be obtained reasonably quickly without requiring elaborate apparatus or a great deal of material. For these reasons D.C.B. specimens were used for preliminary tests in the present study. Cracking to one side, however, was always favoured over cracking along the centre plane (see Fig. 3)» in spite of changing several different experimental parameters such as: the degree of cold work, the specimen thickness, the i n i t i a l crack length tp specimen width ratio, the surface finish and the manner in which the pre-crack was introduced. In an attempt to benefit from the observation of these side cracks, the specimen shown in Fig. 2(b) was tried and proved to be more useful. 2.2 Stress Intensity Calibration The stress intensity calibration for this specimen was obtained by the method originally used by Irwin and Kies [4l], -13-L= 6-4 cm W = 3-2cm BENDING M O M E N T = Pl P P 2P V//////////////////777 " 2P BENDING M O M E N T P P Specimens used in the present work and by others to measure crack velocities as a function of stress intensity. (a) Double Cantilever Beam Specimen (b) Specimen used in the present work (c) Two 3 point bend specimens face to face to show the similarity with (b) -14-PIGURE 3 Double cantilever beam specimen showing side cracking. Fatigue crack at tip of notch. Other cracks are S.C.G. -15-in which the specimen compliance is measured as a function of the crack length. Use is then made of the following relationships [42,43]; Kx - ( G M ) 2 * (1) and G dC 2B ' da where «= stress intensity factor G = energy release rate M = elastic modulus P «? load G = compliance a e crack length Equations (l) and (2) can be combined to give; (2) Kx . B 3/2 Bj~M dC 2 * da (3) The crack was made with a fine jewellers saw and the load and displacement measurements were made with an Instron tensile machine equipped with a D cel l and a 0,2" crack opening dis-placement gauge. The compliance curve is shown in Fig. 4, A polynomial of degree six was fitted to the data points and used to calculate the expression on the right hand side of Eqn. 3-This is plotted as a function of crack length to specimen width ratio (a/fy) in Fig. 5. It is interesting to compare this relationship, between 2 — Irwin [43] included a factor ( 1-1/ )" 2 in the right hand side of this equation, (y is Poisson's ratio). -16-0 0-5 10 1-5 C R A C K L E N G T H ( c m ) 2 0 FIGURE 4 Compliance of the specimen used ln the present work (see Fig. 2fb)) as a function of crack length. -17-14 0 01 0-2 0-3 0-4 0-5 0 6 Q W FIGURE 5 Stress intensity calibrations f o r the specimen used in the present work (Fig. 2(b)) and for a three point bend specimen of similar size (Fig. 2(c)). See Eqns. 3 and 4. -18-stress Intensity and crack length, with the case of three point bending. The latter has been calibrated by means of boundary collocation [44] which gives, for the case of S/W = 4 (see Fig. 2(c)), the relationship} h • ^ /l = 12B2a2Y(a/W ) (4) P 2 r s i r where Y ( a/W ) = 1.93-3.07(a/V)+l4.53(a/M)2-25.1l(a/V)3+25.8(a/i0IJ' B = thickness W - width S = span The right hand side of Eqn. 4 is also plotted agianst a/W in Fig. 5» for the case of S * 2L. There is good agreement between equations (3) and (4) which is perhaps not too surprising considering the similarities in the geometry of the two specimens. The boundary collocation calibration for three point bending is, therefore, applicable to the specimen used in the present study. 2.3 Specimen Thickness A criterion for the minimum thickness of specimens used for fracture mechanics type stress corrosion tests has not yet been established, as i t has in the case of fracture mechanics. Since neither pop-in nor slant cracks are found in stress corrosion 2 i t does not seem reasonable to simply adopt the value of 2.5(K1/oy) used in fracture mechanics. Some work has been done on the effect of specimen thickness on K^gGC values in a titanium alloy, a steel and an aluminum alloy, and the findings have been reviewed by Smith and Piper [45]. The results from these -19-investigations are conflicting. No thickness effect was observed with the titanium and aluminum alloys, however, in the case of the steel, a minimum K^gGG was found i f the thickness was greater than or equal to 2.5(K^ /o ) . There is no data available on the effect of thickness on crack velocities. In view of this, a specimen thickness of 3.2mm. was selected as a compromise among four opposing factors: (i) This thickness resulted in an almost straight crack front, meaning that the crack lengths measured at the surface were representative of the crack throughout the thickness. This is not necessarily the case with thicker samples [46], (ii) This thickness prevented the specimen from twisting appreciably out of the plane of the specimen on loading (i.e., from introducing KJJJ opening mode). (i i i ) It made the final cold rolling and stress relieving relatively straightforward. It might also be noted that the thickness of 3.2mm. does satisfy the condition, B > 2.5(K../c ) , for a l l the materials y used for quantitative measurements, up to a stress intensity of K j ^ 19MPa.m2. (Yield strength data is given in Table II). 2.4 Advantages and Limitations of the Specimen Geometry There are several advantages to the specimen geometry chosenj (i) The specimen is quite compact, easily loaded in the corrosive environment and has a useful crack length of approxi-mately l.?cm. (i.e., from 0.3cm. to 2,0cm.). -20 -(ii) The stress intensity changes by approximately a factor of five, over the useful length of the crack, i f the load is kept constant, see Fig. 5. ( i i i ) The specimen will not be as prone to corrosion product * wedging as the self loading D.C.B, specimen since the stress intensity increases rather than decreasing as the crack grows. There are however, two drawbacks to this type of specimen} (i) The initiation time for cracking may be long since cracking proceeds from low to high stress intensities, (ii) ^2SCC v a^- u e s w i l l be time consuming to determine since each test will only give an upper or lower limit for K-^g^ 2.5 Materials and Specimen Preparation The compositions of the three alloys studied are shown in Table 1. The Jl6 and 310 stainless steels were commercial alloys purchased as 0.64 cm, thick, hot rolled, annealed, and pickled plate. The TRIP steel was produced by vacuum melting a 2-^ kg. heat. The resulting ingot was then electro-slag refined and hot rolled to a thickness of 0.64cm. Further reduction of a l l three alloys was by cold rolling, with or without an intermediate anneal, in order to arrive at the final thickness with the desired amount of cold work. A l l annealing was carried out in sealed Sen-Pak stainless steel containers at 1050°C. for one hour and was followed by quenching in water. The TRIP steel was cold Accumulation of corrosion product in the crack can produce sufficient hydrostatic force to invalidate the calculated stress intensities in wedge loaded D.C.B. specimens, especially in the case of stainless steels, uranium and aluminum alloys [47], -21-TABLE I Alloy Compositions 316 310 TRIP ELEMENT wt % wt % wt % Fe 66.35 53.67 73.20 Cr 16.85 21.8 11.0 Ni 12.08 21.8 8.8 Mo 2.13 3.66 Mn 1.58 1.64 2.01 Si 0.89 1.10 1.10 C 0.0? 0.05 0.22 S 0.011 0.019 0.005 P 0.036 0.02 0.006 -22-rolled at 200°C. when a fully austenitic structure was required, and at room temperature when a partially martensitic structure was desired. The crystal structure of a l l the alloys was checked "by X-ray diffraction, with only the TRIP steel rolled at room temperature displaying an a peak (i.e., b.c.c. phase). Also, only the TRIP steel rolled at room temperature was found to be magnetic. In order to reduce the residual stresses left after rolling, the rolled strip was strained 2% by pulling in tension. This "stress relieving" treatment was only carried out on material used to measure Region I (i.e., stress dependent) velocities. Table II lists the mechanical properties of the different alloys and processing treatments used. Specimens were cut from the rolled strip, so that the cracks would, be either parallel or normal to the rolling direction, and were then machined to a final size of 11 x 6.5cm, The central slot and loading pin holes were then machined and the side notches cut, either by spark machining or with a jewellers saw, the normal length being about Jmm. Occasionally, fatigued pre-cracks were used. The specimens were then polished to a 600 grit finish and a 2cm. long millimetre scale lightly scribed on each side. The entire specimen, except for the ends of the side notches, was then coated in transparent lacquer in order to prevent secondary cracks from initiating ahead of the primary ones. 2.6 Apparatus and Testing Procedure A photograph:of the experimental arrangement is shown in Fig. 6. The tests were carried out in a 500ml. Pyrex glass cell TABLE II Mechanical Properties of the Various Alloys and Processing Treatments (Tensile data measured perpendicular to the-rolling direction at 150°C.) ALLOY % REDUCTION DEFORMATION TEMPERATURE 0.2% OFFSET YIELD STRENGTH ULTIMATE TENSILE STRENGTH % ELONGATION (G.L. - 2.5cm.) HARDNESS °C * MPa MPa Rc 316 25 R.T. 534 903 21.0 29.0 316 50 R.T. 903 1230 9.0 37 310 50 R.T. 696 883 10,5 26.5 TRIP 25 200°C. 580 952 18.0 36.5 TRIP 25 R.T. 655 1080 9.5 42.0 IMPa = 10 Newtons/metre = 0.145 ksi -24-FIGURE 6 Experimental arrangement showing S.G.C. c e l l , reflux condenser and load stabilizing spring mounted in tensile testing machine. -25-fitted with a Teflon cap through which two stainless steel grips were free to slide. This cell was mounted in a Hounsfield tensile testing machine in series with an elastically 'soft* spring, which was used to reduce the drop in load which occurred as the cracks lengthened. The cell was heated by a hot.plate and an immersion heater, the latter being used to control the temperature to within + -|0C. by means of a temperature controller.(Model ?1A Yellow Springs Industrial Go.). The thermistor probe was immersed in the solution close to the cracks. The solution was stirred by a magnetic stirrer in order to reduce temperature and concentration gradients. A reflux condenser was used to minimise the water loss from the cell. Electrochemical potentials were measured with respect to a saturated calomel electrode at room temperature. A luggin capilliary was fixed in position about 1mm. from the specimen, midway between the cracks, and was connected to the reference electrode by means of a magnesium chloride salt bridge. A platinum counter electrode was used. A potentiostat (Princeton Applied Research model 173) w a s used to measure and control the potential. Potentials quoted are with respect to the saturated calomel eleptrode and are not corrected for either liquid junction or thermal gradient potentials. To start a. test, a specimen was mounted in the cel l and 350ml. of solution was added. The solution was prepared by adding distilled water to reagent grade magnesium chloride hexahydrate and heating until the boiling point reached 154°C. This cor-responds to a concentration of MgCl_ [48], The cel l was -26-heated and time allowed for temperature equilibrium to be attained. The mercury column in the tensile machine was zeroed and the desired load applied, or in some cases, to reduce initiation times, a greater load was applied and then reduced once the cracks had initiated. If a temperature of less than 154°C. was required, the solution was allowed to cool to the desired temperature and the temperature controller set. In spite of the reflux condenser some water vapour was lost from the cel l . For the boiling solution, sufficient distilled water to restore the temperature to 154°G, was periodically added. At the lower temperatures the level of the solution was marked on the side of the ce l l and distilled water added periodically to maintain the level. Typically, a few c c . of water were re-quired each day. Measurements of the crack lengths were read, by eye, from the scales on both sides of the specimen, the cylindrical half ce l l acting as an adequate magnifying glass. The cracks were referred to as •north* and 'south* in order to distinguish one from the other. On completion of the experiment, the specimen was removed from the solution, pulled carefully apart in the tensile machine, washed ultrasonically in hot distilled water, dried and stored in a dessicator. The specimens were subsequently mounted for ovservation of the fracture surfaces in a Scanning Electron Microscope (S.E.M,), (E T E C Corporation Auotscan). The scribed scale could be seen when the specimen was tilted appropriately in -27-th e S.E.M. and so this scale was used to locate particular portions of the fracture surface. Quantitative measurements of the relative amounts of intergranular S.C.C, transgranular S.C.C. and ductile overload fracture were made by measuring linear intercepts of each directly from the viewing screen. Crack velocities were obtained from the crack length versus time data by fitting curves to the data points using the method of least squares and differentiating these with respect to time. For Region I a polynomial of degree less than or equal to six was used while for Region II a straight line f i t was obtained. -28-3. RESULTS 3.1 Stress Corrosion Crack Morphology 3.1.1 316 Stainless Steel The nature of the cracks obtained in 316 stainless steel depended on the degree of cold work and on whether the crack was parallel or perpendicular to the rolling direction. In the case of the 2$% and 50% cold rolled material with the cracks propa-gating in the rolling direction, the cracks were usually quite straight and unbranched, although there was some microbranching, especially at high stress intensities (K^ > ^ OMPa.m2 ), An example is shown in Fig. ?(a). The crack front was usually quite straight and perpendicular to the propagating direction, Fig. 7(b). The roughness of the fracture surface increased with stress, as shown In Fig. 7(c). There was a tendency, particularly at low 1. stress intensities (K^ < 15MPa.mi;), for the cracks to grow down-wards at an angle to the normal to the edge of the sample. This angle was usually less than 10°. However, when the rolling direction was normal to the intended cracking direction, the cracks grew downwards at an angle of about ^5° and tended to branch. An example of this is shown in Fig. 8(a). Consequently, quantitative data could only be obtained from.specimens in which the crack was parallel to the rolling direction. Cracks obtained in annealed 316 tended to branch, lead on one side of the specimen and to twist and "bend out of the crack * i i IMPa.m2 - 0.909 ksl(in.) 2 -29 -(a) NOTCH S.C.C. OVERLOAD FRACTURE (c) Kx = 15MPa.m2 K± - 30MPa.m2 Crack morphology for cracking parallel to the rolling direction in 2% cold rolled 316 stainless steel. (a) Side view. 7x. (b) Fracture surface. 4.8x. 4 (c) End view at 15 and 30MPa.m2. 23x . -30-FIGURE 8 Crack morphologies in Jib stainless steel. 12x. (a) Cracking perpendicular to the r o l l i n g direction. (b) Cracking in the annealed alloy. (The crack has been opened to f a c i l i t a t e observation of cracking.) -31-plane. An example of the branching is shown in Fig. 8(b). Tic quantitative measurements were made of cracking in annealed 316 since the branching would have affected both the crack velocity and stress intensity. 3-1.2 TRIP Steel Cracking in the TRIP steel was very similar to that found in the 31& stainless steel. An example of a crack propagating parallel to the r o l l i n g direction in TRIP deformed Z$% at 200°C. is shown in Fig, 9(a). Similar straight and unbranched cracks were found in the TRIP steel rolled at 20°C. Cracks in the annealed TRIP were branched and unsuitable for quantitative measurements. 3.1.3 310 Stainless Steel An example of cracking in 50% cold rolled 3 1 0 is shown in Fig. 9(b). The crack is not straight and displays multiple branching. In spite of this, some measurements were made of the crack velocity at high in the hope that the cracking was in Region II. Fatigued pre~cracks were used since these were help-f u l in i n i t i a t i n g single cracks. Cracking in annealed 310 was also branched. 3.2 Considerations of Errors 3.2.1 Errors in Crack Length Measurements Usually the crack t i p could be clearly seen and the position determined to within ±0,02cm. At very high stress This error was estimated by comparing readings taken just before and after unloading and removing the specimens from the solution. - 3 2 -FIGURE 9 Crack morphology in the 310 and TRIP steels. (a) TRIP - Z% cold rolled at 200°C. Cracking parallel to the rolling direction. 7x. (b) 310 - 50% cold rolled. Cracking parallel to the rolling direction. 17x. - 3 3 -intensities, especially in the softer materials, the size of the plastic zone at the crack tip made determination of the crack length more difficult resulting in an error of approximately ±0.C4cm. 3.2,2 Errors due to the Cracks Growing Downwards If the crack grows downwards, then i t crosses the scale at an angle which will result in an underestimation of the velocity. It can be shown that, i f the crack makes an angle of 10° with the horizontal, then the measured velocity will be in error by about 1.5% and for an angle of 20° the error will be about 5%> See Appendix 1. Also, as the crack grows downwards, the length of the moment arm (L, see Fig. 2(b)) Increases causing the actual stress intensity to be greater than that calculated using the in i t i a l value of L. It can be shown that for a 2cm. long crack at an angle of 10°, is underestimated by about 5%- The downward growth of the crack also results in a! measured ligament length which is less than the actual length. This will result in a calculated stress intensity greater than the true value. The magnitude of this error can be estimated, and for a 2cm. long crack at 10° to the horizontal, i t is found to be 1.5$. See Appendix 1. 3.3 Kinetic Data 3.3.1 316 Stainless Steel 3.3.1.1 Effect of Stress Intensity on Crack Velocity It was observed that both stress dependent (Regions I and III) and stress independent (Region II) stress corrosion cracking were present ln cold worked Jl6 exposed to concentrated magnesium chloride solution. Figs. 10 and 11 show the variation with time of crack length and load for two specimens of 25% cold rolled 316 cracking at 15^0. with i n i t i a l loads of 500N and 1200N respectively. This same data i s plotted as log^V against i n Fig. 12. At low stress intensities the velocity increases rapidly for small increases in K^ , At intermediate Values of K-jV V i s independent of K^t while at higher stress intensities the velocity increases again. Thus, Region I and II and the start of Region III are present. (Region III was not studied i n the present work). If the specimen was unloaded after cracking had initiated, i t was found that the crack continued to grow slowly. Thus, i t was not possible to determine a K 1 S G C value. 3.3.1.2 Effect of Plastic Deformation Fig. 13 shows the variation of crack velocity with stress intensity for two specimens of 50% cold rolled 316 at 15^°G. The variation i s quite similar to that found in the 25% reduced material, except that Region II extends to higher stress inten-s i t i e s . A comparision of the Region II velocities of the 25% and 50% reduced alloy was made at three temperatures by cracking several specimens. The mean values obtained are shown i n Table III. At a l l temperatures the 50% reduced material cracked about 30% faster than the 25% reduced material. 3.3.1.3 Effect of Temperature The data in Table III also show the effect of temperature TIME (hours) FIGURE 10 Variation of crack length and load with time for a specimen of 25% cold rolled 3 i 6 stainless - steel - having an i n i t i a l load "of 500N, HgCl solution bailing- ' at 154°c. 0 • L O A D • C R A C K L E N G T H JL 1 1200 8 0 0 Q < O 4 4 0 0 6 8 TIME (hours) /0 0 FIGURE 11 Var i a t i o n of crack length and load with time f o r a specimen of ?.%• cold r o l l e d 316 s t a i n l e s s s t e e l at 15k°C. having an i n i t i a l lo*.d o f 1200N, Mg^l.^ s o l u t i o n b o l l i n s at 1C4°C. 'TGURE 12 Effect of stress intensity on crack velocity f.or-7.% cold r o l l e d 316 stainless steel'tfxponed to boiling KgOL, solution at 1S^°0. -39 -TABLE III Variation with Temperature of Region II Velocities for 25% and 50% Cold Rolled 316 Stainless Steel 154°C. 136°C 116°C. 50% cold rolled V m/s S.D. m/s >.7pxl0~7 0.4lxl0~ 7 • 9 2.20xl0~7 0.20xl0~7 3 0.8Q5xlO-7 0 .034xl0" 7 3 154°C. 135°C 116°C. 25% cold rolled V m/s S.D. m/s 3.80xlC~7 0.26xl.(f 7 5 1.56xl0~7 0.1?xlO"? 3 0 .59xl0" 7 0.048xl0"7 V is the average Region II velocity S.D. is the standard deviation of V n is the number of velocity measurements -to-on the Region II velocities for "both the cold rolled 316 alloys. The data was plotted as log 1 QV against l/T in Fig. Ik to see i f the crack velocity followed an Arrhenius rate law of the form: V - V e ^ T o where V is the velocity V is a constant o Q is the activation energy R is the gas constant and T is the temperature (°K.) As can he seen, both sets of points lie close to straight lines which yield apparent activation energies of l6 . 0 i 0 . 5 kcal/g.mole for the 25% reduced 316 and 15.1±0.5 kcal/g.mole for the 50% reduced Jl6. Fig. 15 shows the effect of temperature on Region I. At o A 116 G. and < 15MPa.m2, the cracking rate was extremely slow. Because of this and the steepness of the log V - slope i t was not felt that the effect of temperature on the Region I crack velocity could be determined with sufficient accuracy within a reasonable length of time. 3.3.1A Effect of Applied Potential The effect of applied potential (E) on Region I velocities of 25% cold, rolled 316 alloy at 15^°C. is also shown in Fig. 15. There does not appear to be a dependence of V on E between -0.25V(S.G.E.) and -0.3V(S.C.E.). The scatter in the data, however, would mask a l l but a large effect. For E = -0.325V(S.C.E.) 1 there was no cracking at or below 12MPa.m?. -41-FIGURE 14 Arrhenius plot of the Region II velocities of cold rolled 316 stainless steel cracking in a 45^ MgCl, solution. -42-5 0 2 0 O x >-t 0 5 o o _J LU > o <t O 0-2 01 T(°C) E (SCE) • 154 -0-25V • it -0 26V A it -0-275V • II -0-30V O 116 0/C 5 10 S T R E S S 15 2 0 N T E N S I T Y ( M P a / m ) FIGURE 15 Effect of applied potential and temperature on the Region I cracking of 25$ cold rolled Jl6 stainless steel exposed to a h% MgCl2 solution. ( 0/C • open circuit potential). -43-The effect of applied potential on Region II velocities of 25/6 cold rolled 316 at 15^°C. is shown in Figs. 16 and 17. Between -0.25V(S.G.E.) and -0.325V(S.C.E.) there is no detectable dependence on E. On reducing the potential to -0.35V(S.G.E.), the crack slowed down and appeared to stop after a distance of approximately 1mm. 3.3.2 TRIP Steel 3.3.2.1 Effect of Martensite on the Region II Velocity Fig. 18 shows the relationship between V and for 2% cold rolled TRIP steel at 15^°C. in both the fully austenitic and partially martensltlc conditions. The steel containing no martensite has a Region II which extends to higher values. Also, the martensitic TRIP cracks faster In Region II. This was found to be the case at lower temperatures as well, as shown in Table IV. 3.3.2.2 Effect of Temperature on the Region II Velocities The data in Table IV also show the increase in velocity with temperature of the TRIP. An Arrhenius plot of the data is shown in Fig. 19. Both sets of data li e close to straight lines having gradients which yield apparent activation energies of 15.5±0.6 kcal/g.mole for the TRIP steel deformed at 200°C. and 15.6±0.6 kcal/g.mole for the TRIP steel deformed at 20°C. 3.3.3 310 Stainless Steel 3.3.3.1 Effect of Stress Intensity on the Crack Velocity Fig. 20 shows the variation of crack length with time, at 154°G., for a specimen of 5°$ cold rolled 310 stainless steel. -44-V=3 77x|0"7m/sll V = 3-76x|0"7rTVS 0 • NORTH CRACK • SOUTH CRACK 30<K<:35 MPa/m 0 - 3 5 0 V -. 3 4 5 T I M E ( h o u r s ) 7 Effect of cathodic polarization on the Region II velocity of 2% cold rolled 316 stainless steel cracking in a MgCL solution boiling at 154°C. Potentials are w.r.t. a saturated calomel electrode. (o/C «= open circuit, juk ~ 10"^  amps). - 4 5 -FIGURE 17 Effect of anodic polarization on the Region II velocity of 25% cold rolled 316 stainless steel cracking in a MgClg solution boiling at 154°C, Potentials are w.r.t. a saturated calomel electrode. (o/C - open circuit, yUA • 10" amps, mA - 10~J amps). -46-x 3 100% Austenite Partially transformed to Martensite 0 10 20 30 40 50 STRESS INTENSITY ( M P a / m ) 60 FIGURE 18 Effect of stress intensity on crack velocity for Z% cold rolled TRIP steel exposed to MgCl2 solution boiling at 154°C. -47-TABLE IV Variation with Temperature of Region II Velocities for 2% Gold Rolled Austenitic and Partially Martensitic TRIP Steel 154°C 135°C 116°G. Austenitic TRIP V n m/s 3.28xl0"7 2 1.46xl0"7 2 o.550xio" 7 2 Partially Martensitic TRIP V n m/s 3.96xl0" 7 2 1.57xl0" 7 2 0 .656xl0" 7 2 V is the average Region II velocity n is the number of velocity measurements FIGURE 19 Arrhenius plot of the Region II velocities of 2% cold rolled TRIP steel cracking in a h% MgCl_ solution. -49-60 N O R T H C R A C K SOUTH " 2 3 4 TIME (DAYS) FIGURE 20 Variation of Crack Length, load and stress intensity with time for a specimen of 50% cold rolled 310 stainless steel exposed to boiling MgCl2 solution at 154°C. -50-Gracking in the 310 was slower by more than a factor of ten than i t was in either the 316 or TRIP steels at the same stress intensity. After cracking for approximately 1.5mm., the cracks slowed down and continued to crack at a new constant velocity, between one half and one third as fast as the i n i t i a l rate. This change in velocity was quite reproducible and was not accompanied by a change ln corrosion potential. Also, i t was not related to the fatigue pre-cracks since i t occurred even when these were not used. 3.3.3.2 Effect of Temperature on the Crack Velocity The effect of temperature on both the i n i t i a l and sub-sequent velocities is shown in Table V. The mean values are plotted as log^QV against l/T in Fig. 21. The data points for the slower velocity l i e close to a straight line whose gradient gives a value of Q » 18.1±0,8 kcal/g.mole. The two points for the i n i t i a l velocity have a similar temperature dependence. 3.4 Corrosion Potentials 3.4.1 Effect of Stress Intensity on the Corrosion Potential The corrosion potential was found to become more negative (active) on increasing the stress intensity. Fig. 22 shows the change in potential with time of a specimen of 25$ cold rolled 316 at 154°C. as the load was increased in increments of 220N. The i n i t i a l potential was obtained by initiating cracking, unloading the specimen and leaving i t overnight in the solution, during which time cracking continued slowly. On loading tfie specimen the potential decreased (i.e., became more active) quite quickly attaining a new steady state value in less than five minutes. -51-TABLE V Variation with Temperature of Initial and Subsequent Constant Velocities for 50% Cold Rolled 310 Stainless Steel 154°C. 135°C 120°C. Initial velocity-V S.D. n m/s m/s 2.63xl0" 8 0.37xic~ 8 k i . i 5 x i o " 8 2 Subsequent velocity V S.D. n m/s m/s l.OlxlO" 8 O.llxlO" 8 6 0.395xl0" 8 2 0.156x10"8 2 V Is the average Region II velocity S.D. is the standard deviation of V n is the number of velocity measurements FIGURE 21 Arrhenius plot of the crack velocities of 50% cold rolled 310 stainless steel cracking in a k% MgCl, solution. -0-250 ^ -0-3001 ZERO LOAD 220 N 440 N 660 N i i 1 LU d (/> C/) I--I o > _J < h-z LU r-o Q_ Z o t o o o r o u -0-275 -0-325 660 N 65a -0-275 -0-325 200 400 T I M E (S) 880 N 800 1000 T IME (S) 1100 N 660 N 600 I I 00 N 12 00 1300 2000 4000 6000 FIGURE 22 Variation of Corrosion Potential with time as a function of applied load. 25% cold rolled 316 stainless steel ln boiling MgCl 2 solution at 154°C. -54 -On increasing the load, the potential continued to decrease in a similar fashion. On decreasing the load to an Intermediate value, the potential increased immediately, hut required nearly an hour to return to the steady state value. The steady potentials are plotted against stress intensity in Fig. 23. 3.4.2 Effect of Composition and Processing Treatment on the  Corrosion Potential Table VI shows the average corrosion potentials of the different alloys and processing treatments at a temperature of o •!• 154 C. and a stress intensity of 30MPa.m . There was considerable variation in these values and none of them are thought to be more accurate than by±5mV, It can be seen, however, that the TRIP steel is the most active and the 310 the most noble. 3.4.3 Effect of Temperature on the Corrosion Potential Fig. 24 shows the variation of the corrosion potential (E^Qpp) as a function of the temperature for 2$% cold rolled 316 cracking at a stress intensity of 3°MPa..m2. 3.5 Fractography 3.5.1 316 Stainless Steel 3.5.1 .1 Effect of Plastic Deformation The stress corrosion crack path was solely transgranular (T.G.) in the annealed material, whereas both T.G. and Inter-granular (i.G.) cracking were observed in the cold worked material. Al l S.E.M. fractographs are oriented so that the crack propa-gation direction is from left to right. -55--027, -0-28 LU O CO -029 tr a: o o LU -0-30 -0-31 0 10 20 30 40 50 STRESS INTENSITY (MPaVm) FIGURE 23 Effect of stress Intensity on the corrosion potential of Z$% cold rolled 316 stainless steel cracking in boiling MgCl2 solution at 154°C. -0-28 120 130 140 150 160 TEMPERATURE (°C) FIGURE 24 Effect of temperature on the corrosion potential of 2% cold rolled 316 stainless steel cracking in h% MgClg solution at a stress intensity of 25MPa.nJ -56 -TABLE VI Corrosion Potentials of the Various Alloys and Processing Treatments ALLOY AND PROCESSING ECORR TREATMENT V(S.C.E.) , 25% cold rolled 316 - 0 . 2 9 5 25% cold rolled (200°C.) TRIP - 0 . 3 3 1 25% cold rolled (20°C.) TRIP - 0 . 3 3 3 50% cold rolled 310 - 0 . 2 7 0 Values measured at K, = 30MPa.m2 and T = 154°C. -57-Fig. 25 compares cracking under similar conditions in annealed and 2$% and 50$ cold rolled Jl6 stainless steel. The morpho-logical features of the T.G. cracking are similar in a l l three conditions, although the fan shaped features described by Neilson are not so clearly defined in the cold worked samples. Both the transgranular features and the striations on the inter-granular facets were found to match on opposite sides of the cracks as shown in Figs. 26 and 27. This indicates that these features were present when the surfaces were formed. 3.5.1.2 Effects of Electrochemical Potential and Stress Intensity Both the electrochemical potential and the stress intensity influenced the mode of cracking. Experiments were carried out at several different applied potentials to determine the effect of each variable separately, since changing the stress intensity also altered the corrosion potential. Also, i t was observed that the amount of intergranular cracking was greater in the centre of the crack than at the edges. Consequently, the fraction of I.G. cracking was measured as a function of the distance from the edges of the crack when varying E and K^ . The data obtained is displayed on histograms in Fig. 28. It can be seen that the fraction of I.G. cracking increases both with increasing stress intensity independent of applied potential and with decreasing potential independent of stress intensity. Figs. 29 to 33 show some examples of fractography from the specimens used to obtain the data in Fig. 28. Fig. 34 shows an example of the fractography of a crack -58-FIGURE 25 Effect of cold work on the fractography of 316 stainless steel at 15k°C. kOOx. 1 (a) Annealed, K. = lOMPa.m2 (b) 25% cold rolled K-. = 20MPa.m2 (c) 50% cold rolled ^  = 20MPa.m2 -59-FIGURE 26 Matching fractographs from opposite sides of the crack in 25$ cold rolled Jl6 stainless steel. 400x. -6o-FIGURE 2? Matching fractographs from opposite sides of the crack in 2$% cold rolled 316 stainless steel. lOOOx. D I -STRESS INTENSITY (MPa^m) 6 0 -c3 40-•sp 20 • 0 8 0 -6 0 -u CJ CO 4 0 -2 0 -0 -8 0 -c j 6 0 -CJ CO 4 0 -2 0 -2 2 0 | 3 3 - 0 0 L FIGURE 28 Effects of applied potential and stress intensity on the amount and distribution of Intergranular cracking. 2.5$ cold rolled 316 stainless steel exposed to boiling MgClg solution at 154°C. Each histogram shows the variation in cracking mode across the width of the specimen. o CM 6 i in Is-CM O o ro 6 I CJ CO co O > l < z UJ h-O a UJ _J a. Q_ < -62-FIGURE 2 9 Examples of fractographs of 2|# cold rolled 3 1 6 stainless steel at 154°G. ^ - llHPa.mTt E - -0.25V (S.C.E.) 400x. (a) Edge of crack (b) Centre of crack -63 -PIGURE 30 Examples of fractographs of 25$ cold rolled 316 stainless steel at 154°C. K± - 22MPa.m2, E = -0.25V (S.C.E.) 400x. (a) Edge of crack (b) Centre of crack -64-FIGURE 31 Examples of fractographs of 2% cold rolled 316 stainless steel at 154°C. K± = llMPa.m"2", E = -0.3V (S.C.E.) 400x. (a) Edge of crack (b) Centre of crack -65 -FIGURE 32 Examples of fractographs of 25% cold rolled 316 stainless steel at 154°G. KL - 22MPa.m2, E = -0.3V (S.C.E.) 400x. (a) Edge of crack (b) Centre of crack -66-FIGURE 33 Examples of fractographs of 2$% cold rolled 316 stainless steel at 154°C. K± - 33MPa.m2. Centre of crack. 400x. (a) E - -0.325V (S.C.E.) (b) E - -0.35V (S.C.E.) -67 FIGURE 34 Example of fractograph of 2% cold rolled 316 stainless steel cracking without an applied load. 800x. -68-propagating in the absence of an applied load at an open-circuit potential of -0.275V(S.C.E.). There is no evidence of I.G. fracture, which is in agreement with the prediction from the data in Fig. 28. i At very high stress intensities (>55MPa.nrz), patches of ductile fraoture were observed amongst the stress corroded fracture, Fig. 35. The amount of ductile fracture increased i with Increasing stress intensity, above K^ = 55MPa.m2, and was found mostly at the edges of the specimen, Fig. 36. At active potentials, E < -0.325V(S.C.E.) precipitates were observed on the fracture surface, Fig. 37t and X-ray energy analysis of these, conducted in the S.E.M., showed them to contain magnesium. Wilde [493 identified precipitates formed under similar conditions as Mg(OH)g. 3.5.1.3 Effect of Temperature Varying the temperature had l i t t l e effect on the appearance of the fracture surface. Examples of fractographs from specimens of Z$% and 50$ cold rolled 316 stainless steel, cracked at similar stress intensities (K^ «• 25MPa.m2*) at both 154°C. and ll6°G., at open circuit potentials are shown in Figs. 38 and 39. This implies that the variation of fractography with Is similar at different temperatures, but the variation with E must be different at different temperatures, since changing the temperature has been shown to alter the corrosion potential, see Fig. 24. -69 -FIGURE 35 Example of fractograph of 50$ cold rolled 316 stainless steel showing small area of ductile fa i l u r e . K x = 70MPa.m2. 400x . - 7 0 -(A) LJ CC h-O < L_ UJ _J r-O 3 O $5 12 10 8 2 0125 cm FIGURE 36 Distribution of ductile fracture in a specimen of 50$ cold rolled Jl6 stainless steel exposed to boiling MgCl2 solution at 154°C. (a) As a function of position across the specimen. (b) As a function of stress intensity. -71-(a) 0 0 FIGURE 37 Examples of precipitates formed on the fracture surface during cathodic polarization. (a) E = -0.325V (S.G.E.) 800x. (b) E - -0.35V (S.G.E.) 400x. -72-FIGURE 38 Effect of temperature on the fractography of 50$ cold rolled 316 stainless steel. lOOOx . (a) T = 154°C.f K, - 27MPa.m^  (b) T - 116°C., K.. - 27MPa.m2 -73-(a) OO FIGURE 39 Effect of temperature on the fractography of 25$ cold, rolled 316 stainless steel. 400x. (a) T - 154°G.f Kx - 27MPa.m^  (b) T - ll6°G.f K, - 27MPa.m^  -74-3.5.2 TRIP Steel 3.5.2.1 Comparison with the Fractography of the 316 The fractography of the TRIP steel was similar to that of the 316, the major difference being the greater amount of I.G. cracking i n the TRIP cf. Figs. 39 and 43. 3.5.2.2 Effects of Cold Working and the Presence of Martensite The cold worked austenite showed more I.G. cracking than did the annealed material at the same stress intensity. This i s shown i n Fig. 40. The TRIP containing the martensite had a dark coloured fracture surface covered with corrosion product in contrast to the li g h t grey clean looking surface of the TRIP containing only austenite. This made i t rather d i f f i c u l t to obtain sharp fractographs from the p a r t i a l l y martensitic steel. However, the effect of the martensite i s shown i n Fig. 41. It can be seen that there i s more T.G. cracking i n the martensitic TRIP. Also, although the nature of the transgranular surface i s the same for both TRIP materials, i t i s different from that of the 316 i n that i t has a nodular appearance rather than a striated or ridged look. To check that this was not due to excessive corrosion of the TRIP steel in solution after cracking had taken place, an attempt was made to match the features present on opposite sides of the crack. The result i s shown in Fig. 42. It can be seen that the nodules on one side correspond to pits on the other, thus revealing that the nodular features result from the cracking and not from subsequent corrosion. -75-FIGURE kO Effect of cold work on the fractography of the TRIP steel. 400x. (a) Annealed TRIP, Kj- lOMPa.m2, T - 154°G. (b) 25$ cold rolled (at 200°G.), K^ - 9MPa.m2, T - 154°C. FIGURE kl Effect of martensite on the fractography of 25$ cold rolled TRIP steel. lOOOx. (a) Fully austenitic, K - 22MPa.m% T - ll6°C. (D) Containing martensite, K - 22MPa.m% T = ll6°G. FIGURE 42 Matching fractographs from opposite sides of the crack in 25$ cold rolled (at 200°C.) TRIP steel. 800x. -78-3.5.2.3 Effect of Stress Intensity The amount of I.G. cracking in the fully austenitic TRIP increased with increasing stress intensity, as was found for the 316. This is shown in Fig. 43. Also, at lower stress intensities, the T.G. surface shows more ridges and striations and resembles that found in 316, cf. Fig. 34. 3.5.2.4 Effect of Temperature Fig. 44 shows examples of fractographs of the austenitic TRIP cracked at 154°C. and 119°C. at similar stress intensities. The specimen cracked at 119°G. shows more of the nodular type of 1 0 fracture than does the one at 154 G. 3.5.3 310 Stainless Steel 3.5.3.I General Description Cracking in 310 was solely transgranular under a l l the conditions investigated. See Figs. 45 and 46. The surface has a more angular appearance than that of the 316, especially in the case of the annealed material, Fig. 45(a). The stress intensity (approximate because of branching) does not appear to alter the fractography, cf. Fig. 45(b) and 45(c). Also, neither lowering the temperature nor changing the electrochemical potential from the open circuit value of - 0.27V(S.C.E.) to -0.325V(S.C.E.) had any observable effect. See Figs. 46(a) and 46(b) respectively. -79-FIGURE 43 Effect of stress intensity on the fractography of 25$ cold rolled (at 200°G.) TRIP steel at open circuit potentials. lOOOx. (a) No applied load, T = 154°G. (b) ^ - 8.5MPa.m2, T = 154°G. (c) K x - 40MPa.m2, T - 154°C. -80-FIGURE kk Effect of temperature on the fractography of the 25$ cold rolled (at 200°C.) TRIP steel. lOOOx. (a) T = 154°C, Kx - 8.2MPa.m2 (D) T - 119°0., K, - 6.6MPa.m2 -81-FIGURE 45 Effects of cold work and stress intensity on the fractography of 310 stainless steel. 400x. (a) Annealed, Kx - 20MPa.m2, T - 154°C. (h) 50$ cold rolled, ^  - 25MPa.m2, T - 154°C. (c) 50$ cold rolled, K., - 50MPa.m% T - 154°C. -82-FIGURE k6 E f f e c t s of temperature and applied p o t e n t i a l on the fractography of 50$ c o l d r o l l e d 310 s t a i n l e s s s t e e l . 400x. (a) T - 120°G.t K-- 25MPa.m2, E - -0.27V (S.C.E.) (b) T - 154°G.f K 1- 25MPa.m2, E = -0.325V (S.G.E.) -6> 4. DISCUSSION 4.1 Crack Branching The observations made on crack branching can be summarized as followst (i) There was more branching i n the low yield strength alloys, ( i i ) Transgranular cracks branched more than intergranular cracks. ( i i i ) Intergranular cracks branched only when propagating normal to the r o l l i n g direction. The lower yield strength (O ) of the annealed alloys w i l l y result in a larger plastic zone (r ) at the crack t i p since [42], y Consequently, If crack propagation i s dependent on plastic straining at the crack t i p (e.g., to rupture a passive film) then, i n the low yield strength alloys, this w i l l occur over a larger area and in-crease the likelihood of forming more than one crack. The ten-dency to branch, however, must be dependent on more factors than just the size of the plastic zone, since no branching was observed in the 25$ cold rolled 316 at a stress intensity of 25MPa.in2\ (At I 25MPa.m7 the anticipated plastic zone size w i l l be the same as that at which branching occurred in the annealed condition i.e., at " lOMPa.m2). Other factors could be the radius of the crack t i p , the dislocation density and the nature of the active s l i p systems at the crack t i p and how these might affect the s l i p step height. The s l i p step height i s important since i t must be -84-at least as large as the passive film thickness in order to expose the underlying metal. The downward cracking, shown in Fig. 8(a), indicates a strong preference for the I.G. cracks to follow the r o l l i n g direction, in agreement with the observation of Hyatt and Speidel [5] that I.G. cracking followed the grain flow in cold extruded aluminum alloys. The elongation of the grains in the cold rolled alloys w i l l result in an almost straight inter-granular path in the r o l l i n g direction. In the perpendicular direction, an I.G. crack would have to wind back and forth be-tween the grains, making cracking more d i f f i c u l t and providing many opportunities for the crack to branch. Crack branching in high strength aluminum alloys has been examined by Speidel [50]• His observations led him to the conclusion that branching only occurred i n alloys having a relatively isotropic mlcrostructure and only i f the stress intensity was greater than 1.4 Kp, where Kp was the value at which Region II began. The present results are in agreement with Speidel's conditions for branching, except for the annealed alloys, which branched at stress intensities as low as 12MPa.m2 and were certainly cracking in Region I. In explanation of his observations, Speidel proposed that two cracks could form and continue to grow, only i f there was sufficient energy available to maintain each crack with a Region II velocity. If less energy was available, one crack would have a slower velocity than the other and would be quickly s t i f l e d -85-as the region of maximum st r e s s moved away from i t . Thus, the c r i t i c a l energy release rate f o r crack branching would be 2Gp, where Gp i s the G value corresponding to Kp. This r e s u l t s i n a c r i t i c a l s t r e s s i n t e n s i t y of 2.Kp. (See Eqn. l ) . The above argument i s probably v a l i d f o r a wedge loaded D.C.B. specimen, (which i s what Speidel used) since K^ decreases as the crack grows, but i t would not apply to the present arrange-ment because K^ increases. In general, i f a crack branches at a st r e s s i n t e n s i t y of l e s s than 1.4 Kp, then, although the longer crack w i l l grow more r a p i d l y , the shorter one w i l l continue to grow provided i t s stress Intensity remains above K ^ ^ . Regarding the present study, while the value f o r the shorter crack w i l l continue to decrease as the longer crack grows f u r t h e r ahead of i t , the increase i n K^ with crack length and the compliant spring i n s e r i e s with the specimen w i l l both help to delay the decrease i n the s t r e s s i n t e n s i t y of the shorter crack. In essence, the reason that crack branching was observed i n Region I i n the present work, but not i n that of Speidel, may be due to the nature of the specimen geometry and the loading conditions rather than to a d i f f e r e n c e i n the S.C.C. properties of the a l l o y s . 4.2 K i n e t i c s 4.2.1 Region I The cracking rate i n Region I, f o r the 2% c o l d r o l l e d 316 s t a i n l e s s s t e e l , increased r a p i d l y with s t r e s s Intensity, the gradient d log^v/dK^^ being equal to approximately 0.5(MPa.m2)~ . (See F i g . 15). This i s close to the values obtained f o r high -86-strength aluminum alloysj 0.5 to 1.5(MPa.m2)~ [5], and for titanium alloys; < 1.0(MPa.m2)~ [7]. There have been com-paratively few attempts to explain this dependence. Hillig and Charles [51] proposed a model for Region I S.C.C. based on stress activated dissolution. Their treatment was modified by Wiederhorn and Bolz [52] who obtained the relationship V oc exp [(-2V K-jA-n/?)2)/^] where V is an activation volume and /O is the crack tip radius While this has the observed dependence of velocity on stress, the model does not explain the significance of V or the mechanism by which the stress intensity increases the dissolution rate. The cracking which was observed in the unloaded specimens was probably caused by the corrosion product in the crack pre-venting the specimen from completely relaxing when the load was removed [ l ] . Residual stresses from the cold rolling may also have been responsible. 4.2.2 Region II The Region II velocities of the Jl6 and TRIP steels were similar, as were the apparent activation energies (see Figs. 14 and 19). The crack velocity of the 310» on the other hand, was considerably slower although the apparent activation energy was similar, (see Fig. 21). The extent of Region II was dependent on the alloy and the processing treatment. The start of Region II varied con-siderably from specimen to specimen, (see Fig. 15), the average A stress intensity being approximately 20MPa.m2 for both the 316 -87-and TRIP steels. The stress intensity corresponding to the end of Region II was related to the onset of ductile fracture. A comparison of Figs. 13 and 36 shows that for 50% cold rolled 316 ductile fracture was observed on the fracture surface at stress intensities above 65MPa.m% but that the transition 1. from Region II to Region III started at a value of 50MPa.m2. 1. It is possible that some tearing occurred below 65MPa.m2, but that i t was too localized to be observed. While the higher yield strength of the 50% cold rolled 316 extended Region II to higher values of K^ , the lower ductility of the TRIP containing martensite caused i t to f a i l completely at a stress intensity of 1. l i t t l e more than 50MPa.m2, (see Figs. 12, 13 and 18). The reason for the decrease in velocity which occurred in the 310 after the crack had grown by approximately 1-| mm. is not known. The observed crack branching may have been responsible, either by lowering into Region I at the instant of branching, or else, by causing the crack to follow a crooked path which would reduce the average velocity. However, i t was not possible to say with certainty whether either of these effects were taking place. The stress intensity did Increase significantly during crack growth in the 310» as was evident from the appearance of the plastic zone at the crack tip. Even though branching may have had some influence on the stress Intensity and crack velocity, i t is thought that the constant velocities measured were the result of a stress intensity independent process. 4.2.3 Cracking Mechanisms 4.2.3.1 General Considerations Any mechanism of cracking capable of accounting for -88-the Region I and Region II results must satisfy the following conditions s (i) The crack velocity should increase rapidly with stress intensity up to approximately = 20MPa.m2 and then be in-dependent of stress intensity until the onset of ductile fracture. (ii) The crack velocity should be independent of applied potential between -0.25V (S.C.E.) and -0 .30V (S.C.E.) in Region I and between -0.25V (S.C.E.) and -O.325V (S.C.E.) in Region II. At potentials more active than these cracking should stop, ( i i i ) The apparent activation energy for Region II cracking should li e between 15 and 18 kcal/g.mole. Furthermore, i f the mechanism is the same for a l l the materials studied i t must be compatible with the followlngj (iv) 50$ cold rolled 316 cracked faster ln Region II than 2% cold rolled 316. (v) The TRIP steel containing martensite cracked faster in Region II than the fully austenitic TRIP. (vi) The crack velocity was independent of the crack path in 316. (vii) Region II cracking was slower by more than an order of magnitude in the 310 than in either the 316 or TRIP steels. These results differ ln several ways from those found in other systems. Most of the published work on the temperature dependence of Region II, gives apparent activation energies of between 3 and 5 kcal/g.mole, e.g., glasses [4 ] , titanium alloys [7] and high strength aluminum alloys [5 ] . Hyatt and Speidel [5 ] , however, obtained an activation energy of 20 kcal/g.mole -89-for Region II cracking in a 7039-T61 alloy and, in a similar alloy under certain conditions, they observed two stress intensity independent regions, Fig. 47. The values of the apparent a c t i -vation energies, as measured from Fig. 47, are 17.3 kcal/g.mole and 2.6 kcal/g.mole for the slower and faster plateaus respectively. Furthermore, the independence of crack velocity on applied potential was not found for either aluminum or titanium alloys, both of which cracked faster at more noble potentials [5, 7] although in these alloys the applied potential was varied by more than 1 volt. The range of potential over which cracking can occur, ln the present system, i s limited to less than lOOmV [53!• At potentials noble with respect to the corrosion potential the steels p i t rapidly while at active potentials ( i . e . , w.r.t. ECORR^ CTa-ck±ng stops altogether. Time to failure tests for 304 stainless steel in MgClg solution show a change of less than a factor of two in failure times over the available range of potential [54]]. This is consistent with the present results for Regions I and II i f i t is assumed that the i n i t i a t i o n time i s potential dependent. Before discussing stress independent mechanisms which might explain the Region II results, the poss i b i l i t y that the actual stress at the crack t i p does not increase in Region II should be considered. Inglis [55]] showed that the stress at the 4 crack t i p ( o^j_p) was equal to 2o( a / / 3 where o i s the applied stress and /O the radius of the crack t i p . Also, since 4 •= o(fTa)r (see page 3)t the crack t i p stress can be expressed - 9 0 -STRESS INTENSITY (kg - m m ~ 3 / 2 j 10 20 30 40 50 60 70 80 10" 10 tr 1 >-o o _J ai > o < cc o 10" CO O 10 cc o u CO CO Ui cc co -2 10 -3 10 I I , , ! , r _ " A L L O Y 7039-T64 CRACK ORIENTATION:TL (SHORT TRANSVERSE) 5 M AQUEOUS Kl POTENTIAL: -700 m V V S SCE © o 70° C 10 15 20 25 STRESS INTENSITY (MPa.ra2) 90 100 30 10" 10 -2 10" 3 ~ 10 10 -5 10 -6 . — 7 IU H O O _ l LU > o < cc o 2 : o CO O CC cc o o CO CO LU CC FIGURE 47 V-^ plot for 7039-T64 aluminum alloy in 5M aqueous K 1 solution showing two stress intensity independent regions and the effect of temperature (after Hyatt and Speidel [ 5 ] ) . - 9 1 -as ° t i P • V< v>>* Thus, If is increased, the stress at the crack tip may not change if the radius of the crack tip also increases i.e., if the crack is blunted. If Region II is due to crack blunting, then it might be expected that if the yield strength was Increased then the stress intensity at which Region II begins would be greater, resulting in a greater velocity which is in agreement with condition (iv) above. 4.2.3.2 Dissolution Electrochemical dissolution alone does not explain the cracking kinetics very well. If crack growth occurs by dissolution then i t might be expected that anodic polarization would increase the crack velocity. Over the range of potentials in which cracking occurs in MgClg solution, the anodic current density, as measured on a static surface, changes by more than two orders of magnitude [19]» The dissolution rate in Region II could be independent of potential i f the anodic reaction was under diffusion control, but then the activation energy would have to be closer to that for mass transport in the electrolyte, i.e., 3 - 5 kcal/g.mole [56]. Nevertheless, an activation energy of between 15 and 18 kcal/g.mole is in the observed range for dissolution e.g., Staehle reported values between 12 and 18 kcal/g.mole for the dissolution of nickel in lNHgSO^  on static and straining electrodes [19]. The faster velocities in Region II, caused by the cold work in the 316 and the martensite in the TRIP could be due to -92-an increase in the exchange current density resulting from an increase in the number of active sites for dissolution such as at dislocations [573. Nickel, being more noble than iron or chromium, could exert its effect on the cracking rate by raising the reversible electrode potential E Q. This could in turn raise the corrosion potential, (as was observed) and lead to a reduction in the corrosion current, as described by Nikiforuk [293. While this explanation will result in a slower crack velocity for the 310, i t will also require that the velocity be dependent on the applied potential, which was not observed, at least in the case of the 316. If dissolution alternates with a period of film rupture then an electrochemical process can explain the presence of a critica l potential for S.G.G. At a given stress intensity sufficient dissolution must take place at the crack tip to allow build up of enough strain to break the passive film [583. As the potential is made more active w.r.t. eQQRR» the dissolution rate will decrease and the repassivation rate will increase, thereby reducing the total amount of dissolution possible and preventing film rupture. At higher stress intensities more strain would result from a given amount of dissolution, thus requiring a more negative potential to prevent cracking, as was observed. The slowly decreasing velocity observed on changing the potential to -0.35V (S.C.E.), (see Fig. 16) could be related to a change of pH in the crack. The increase In the rate of proton reduction caused by cathodic polarization would raise the pH in -93-the crack. This higher pH, together with the more negative potential, would reduce the dissolution rate, increase the repassivation rate and hence reduce the crack velocity. A reduction in overall dissolution would decrease the supply of hydrogen ions from hydrolysis of the metal ions, causing a further increase in pH and hence a further reduction in crack velocity. 4.2.3.3 Adsorption The cracking kinetics can he explained qualitatively by a process in which an adsorbed species weakens the already stressed metal bonds at the crack tip [59]. In Region II the crack velocity could be controlled by the rate of adsorption of the damaging species. Activation energies for chemisorption can cover a wide range of values, but are frequently of the order of 20 kcal/g.mole [j>6]]. In Region I the cracking rate could be controlled by a deformation process. Deformation at the crack tip could increase the number of adsorption sites, either by breaking a passive film, or else by increasing the number of surface imperfections. An adsorption process is also consistent with the observed independence of crack velocity on applied potential. In Region I a process controlled by deformation would not be affected by potential, provided there were sufficient species available to adsorb. While chemisorption is often restricted to a limited range of potential, It is possible that over the limited range studied there was always more than sufficient sites available having suitable potentials. -94-The critical potential for S.C.C. in 18-8 stainless steel exposed to MgClg solution has been explained by Uhlig and Cook [60] as the potential below which Insufficient chloride ions can adsorb on surface imperfections at the crack tip. Raising could lower the critical potential either by increasing the number of adsorption sites or else, by making these sites more active. The faster cracking rates in Region II caused by the cold work in the Jl6 and the martensite in the TRIP could be due to an Increase in the number of adsorption sites resulting from the increased dislocation density In the Jl6 or the lattice strains produced by the martensite transformation in the TRIP. The slower cracking rate observed in the J10 could be caused by a greater activation energy for chemisorption due to its higher nickel content. The activation energy for chemisorption depends on the heat of adsorption as shown In Fig. 48. A decrease in the magnitude of the heat of adsorption AH will result in ct curve II shifting towards higher potential energies, making Q greater. Heats of adsorption for chloride ions on stainless steels are not available, but a qualitative estimate of the effect of nickel can be obtained from the heats of formation of NiClg, FeClg and CrCl 2 which are - 6 5 . 1 , - 7 2 . 2 and -85.2 kcal respectively [ 6 l ] , The lower affinity of nickel for chloride ions will likely result in a lower heat of adsorption and consequently a greater acti-vation energy. (See Fig. 48). The observed activation energy for the 310 of 18.1 kcal/g.mole is greater than that for the other alloys (15.1 to 16.0 kcal/g.mole). The ratio of the velocities predicted by an Arrhenius rate law, -95-DISTANCE OF ADSORBED UNIT FROM THE SURFACE FIGURE 48 Relationship between the heat of adsorption and the activation energy for chemlsorption (after Glasstone, Eaidler and Eyring [56] ). I i s the potential-energy distance curve for a chloride ion in solution. II are potential energy distances curves for a chemisorbed chloride ion. A smaller heat of adsorption A Ha (b) w i l l l i k e l y result i n a greater activation energy for chemisorption Q (b). -96-V = V q exp (-Q/RT), for values of Q of 15.5 and 18.1 kcal/g.mole is given by: The above calculation assumes that the pre-exponential term V q is the same in both cases. Comparing Tables III, IV and V shows that at 154°C. the 310 cracked approximately 35 times slower than the other alloys, which considering the errors in the apparent activation energies Is in reasonable agreement with the above explanation. 4.2.3.4 Absorption of Hydrogen Diffusion of hydrogen in austenite could be a rate controlling process for Region II. Austin and Elleman [62] obtained a value of 0.6l eV/atom (i.e., 14.2 kcal/g.mole) for the bulk diffusion of tritium in 304 and 316 stainless steels over the temperature range of 25°C. to 222°C. (The result was obtained by injecting tritium into the sample via the reaction ^Li ( n,0t) H^, diffusion annealing and then measuring the tritium concentration profiles, thus avoiding the surface trapping effects found in methods involving permeation rates through membranes [63]). The effect of martensite in the TRIP steel could be to shorten the length of the slow diffusion path through the austenite, diffusion in the b.c.c. a'martensite being faster than in the f.c.c. austenite [63]. The Region II results dp not support a process in which the rate of transformation to martensite at the crack tip is important, because i f this were the case the 25$ cold rolled -97-austenltic TRIP might have been expected to crack faster than the 2% cold rolled 316. In fact, the 316 cracked slightly faster although the difference may not have been significant. (See Tables III and IV). 4.3 Corrosion Potentials The effect of stress on the corrosion potential E C Q R R has been observed by Barnartt and Van Rooyen [64] for 18-8 stainless steel in MgClg solution. They found that EQQpp became less noble after the crack had initiated (on stressed wires) and that the potential continued to decrease until failure occurred. The total potential change was -40mV. They explained their observations by a shift In the anodic polarization curve towards greater currents. Such an explanation is valid for the present observations also. Fig. 49 shows a schematic Evan's diagram Illustrating the effect of stress on the corrosion potential. Curve 1 is the anodic polarization curve representing a film covered surface remote from the crack tip. Curves 2 (a), (b) and (c) represent the anodic dissolution at the film free surface at the crack tip at different stress intensities. Curves 3 (a), (b) and (c) are the composite anodic polarization curves for different values of K^ . If i t is assumed that the cathodic reaction (curve 4) is not significantly affected by the stress intensity then the cor-rosion potential will become less noble on increasing the stress intensity. The increase in the amount of dissolution at the crack tip with K., can be explained in several ways: -98-LOG(CURRENT) FIGURE 49 Schematic Evans diagram illustrating the effect of stress intensity on the corrosion potential. (See text for explanation). -99-(i) The exchange current density may Increase due to the existence of more active sites at higher stresses. (ii) As the crack velocity increases the rate of formation of unfilmed surface will Increase and raise the overall dissolution rate. ( i i i ) If the crack grows by a film rupture/dissolution process then at higher stresses the larger plastic zone will result in a larger area of exposed metal and hence more dissolution, i.e., a wider crack. (iv) At higher stress intensities the effective surface area for dissolution is increased due to the greater roughness of the surface. (See Fig. 7(°))» Of these four possible effects, the f i r s t three would result In a rapid change of Enrtt3_, with K.., whereas the fourth would require time for the topography of the fracture surface to change. Thus, the f i r s t three could explain the almost instan-taneous change of eQQRR o n increasing but only (iv) could explain the slow return to equilibrium on reducing K^ . 4.4 Fractography 4.4.1 General The fractography of the Jl6 indicated that the mode of cracking was subject to the following conditions« (i) I.G. cracking was more probable in the cold worked condition than in the annealed alloy. (ii) The amount of T.G. cracking increased as the applied potential was made more noble. -100-( i l i ) The amount of I.G. cracking increased as the stress intensity was raised. (iv) There was more I.G. cracking in the centre of the crack than at the edges. Nielsen []l"j showed a photograph of a U-bend specimen of 304 stainless steel cracked in 42$ MgClg solution in which the fraction of I.G. cracking was greatest in the last part of the specimen to f a i l . Nielsen did not relate the change in crack path to stress, but instead to the fact that the part which had cracked last was in i t i a l l y in compression and this, together with diffusion of hydrogen into the lattice, caused decohesion of the grain boundaries. Okada and co-workers [65] observed a change from T.G. to I.G. cracking between the f i r s t and last parts to crack of a tensile specimen of 304 stainless steel under constant load in MgCl.2 solution and related this to an increase in stress. These workers also observed a change from I.G. to T.G. cracking on anodic polarization of JlG stainless steel. It is interesting to note that in the case of both alum-inum and titanium alloys the cracking mode changes from I.G. to T.G. as the stress intensity is increased [_5~], 4.4.2 Effect of Applied Potential The effect of applied potential on the cracking mode is compatible with either a dissolution/film rupture process or with an adsorption mechanism. Staehle [vf\ has reported that S.C.C. of Fe-Cr-Ni alloys in concentrated MgClg solutions occurs at potentials close to the transpassive knee of the anodic polar-ization curve, i.e., close to the region where localized film - 1 0 1 -breakdown occurs. Also Payer and Staehle [66] reported that under suitable conditions corrosion pits initiate preferentially at grain boundaries, indicating the less passive nature of the film there. Thus, in terms of a film rupture process, at active applied potentials the grain boundaries may be the only sites where sufficient dissolution can take place between film rupture events to sustain cracking. At more noble applied potentials no sites on the oxide film will be particularly passive and so the effect of the grain boundaries will be less important, the crack path being determined by a film rupture process such as slip. In terms of adsorption the crack path would be dependent on the potential of the adsorption sites. Due to the lattice disorder at the grain boundaries,adsorption sites there would be at more active potentials than sites distant from the grain boundaries. As the potential is made more noble than the criti c a l potential for cracking, the fi r s t sites for adsorption will l i e on the grain boundaries, resulting in I.G. cracking. A further increase in applied potential will allow adsorption to take place on less active sites away from the grain boundaries, resulting in T.G. cracking as well. 4.4.3 Effect of Stress Intensity The effect of stress intensity on the crack path may have been due to the potential at the crack tip becoming more active as the stress intensity was Increased, even though the applied potential was constant. This could occur as a result of an IH drop along the crack, the value of which would increase -102-as the crack velocity Increased. (See explanation below). If the crack path i s determined only by the potential, then from Fig. 28 It can be seen that at -0.25V (S.C .E.) a potential change of at least 50mV. would be required to account for the change in 4 cracking mode caused by an increase i n from lOMPa.nr to 4 30MPa.m . Beck estimated that the IR drop along the stress corrosion cracks ln titanium alloys was greater than 150mV, [[6?]. Assuming a par a l l e l sided crack, the IR drop would be given as a f i r s t approximation byj A E • i . l / c (5) where A E Is the IR drop along the crack 1 i s the average distance between anodic and cathodic sites. I i s the current density at the crack t i p c i s the eleotrical conductivity of the solution. Also i f i t i s assumed that the crack propagates solely by dissolution then the anodic current density (in view of conservation of mass considerations) i s related to the crack velocity by the following expressions 1 - V/O zF/A (6) where p •» density of the alloy z • average valence of the metal ions A m average atomic weight and F « Faraday This does not rule out adsorption since weakening of the metal bonds by adsorption could result in either bond rupture or dissolution. -103-For V - 4 x 10~7m/s,/3 - 8.0 g/cc, z - 2 and A • 56, i «= 1.2 amps/cm2. In order to maintain electroneutrallty at a l l points in the electrolyte there will be a similar current density of anions towards the crack tip. For a total current density of 2.4 amps/cm , AE - 50mV and a conductivity of 2.0 ohnf 1cnf 1 [68], Eqn. (5) gives a value for 1 of 0.4mm. This is not an unreasonable distance in view of the large differences between the anodic and cathodic current densities i.e., the cathode must be considerably larger than the anode. An IR drop effect can also explain the change in crack path across the width of the crack since the potential drop will be less at the edges of the specimen. The above explanation is also valid i f the crack grows by a process in which adsorption results in bond rupture, but in that case the current density might be less since only sufficient current to passivate the sides of the crack would be required. It is also possible that the change in crack path across the width of the specimen could be due to a change in stress state. It is not clear, however, why the plane strain conditions at the centre of the specimen should favour intergranular cracking. Patches of dimple fracture on stress corrosion fracture surfaces, formed at high stresses, have been observed before [69]. That these occur predominantly at the edges of the specimen, in the present case (see Fig. 36), is probably due to the greater shear stress in the plane stress regions at the edges of the specimen. 4.4.4 Effect of Composition The different cracking modes of the three alloys could be due to their different alloy compositions or else to their -104-different corrosion potentials. Specifically, Mo additions have been shown to increase the amount of I.G. cracking in Fe-l6Cr-15Ni alloys [65]. Also, Mo improves the pitting resistance of austenitic stainless steels [70]. A recent paper by Suglmoto and Sawada £?l] Indicates that addition of molybdate ions to the corrosive environment can also inhibit pitting. On the basis of surface analysis, these authors conclude that i t is the adsorption of 2-MoO^  ions from solution which prevents pitting. With regard to adsorption controlled S.C.C., competitive adsorption of MoO^  ions could restrict adsorption of the damaging species to the more active adsorption sites at the grain boundaries. It is also possible that the more active corrosion potentials of the TRIP and yi6 alloys resulted in the greater amounts of I.G. cracking in these alloys since lowering the applied potential favoured I.G. cracking in Jl6. 4.4.5 Effect of Martensite Nodular features similar to those on the fracture surface of the TRIP steel, especially when i t contained martensite, have not been reported before. In view of the fact that they were more prominent at lower temperatures and at higher stress Intensities (see Figs. 43 and 44), I.e., conditions which favour the formation of stress induced martensite, i t Is likely that they are associated with martensite. This implies that the martensite is formed in the fully austenitic TRIP, before or at the same time as the new fracture surface, since the nodular features were found to match on opposite sides of the crack. (See Fig. 42). -105-The fractography associated with a martensite in the case of 304L stainless steel exposed to MgCl^ solution [35] is different from that observed for the TRIP steel. This difference could be related to the higher carbon content of the TRIP steel (0.22 wt. % G as opposed to < 0.03 wt. % G for the 304L). It is possible that precipitation of carbides from the a7phase in the TRIP steel at 154°C. could result in the nodular features because the carbon content exceeds the equilibrium solubility in the b.c.c. a phase. 4.5 Assessment of S.C.C. Models The results have been discussed in terms of three processes thought to be important ln the S.C.C. of austenitic stainless steels, namely, electrochemical dissolution, absorption of hydrogen and adsorption of a damaging species. A film rupture model for S.C.C. can account for most of the present observations with the exception of the effects of applied potential and temperature on the Region II cracking. While a model involving hydrogen absorption can explain some of the kinetic results i t is difficult to see how i t can account for the effect of applied potential on the fractography. A model of cracking involving adsorption is consistent with a l l the present observations although there is not enough known about the details of chemisorption to thoroughly test the model, e.g.» activation energies for adsorption of specific species on the specific alloys. Foley [72] has pointed out that the concentrated solutions encountered in localized corrosion cells will result in the anions -106-and cations being very close together making adsorption quite probable. In 45$ wt. $ magnesium chloride solution of pH2 there are approximately 11 moles/litre of chloride ions, 10~ moles/ —12 lit r e of hydrogen Ions and 10 ~ moles/litre of hydroxyl ions. Thus, the adsorbing species Is very likely to be a chloride ion or perhaps a complex metal chloride ion, although the possibility that i t could be a hydrogen ion cannot be excluded. A model for cracking involving adsorption does not rule out the possibility that other processes such as dissolution, deformation or film rupture may be important. As mentioned previously plastic deformation or film rupture might be necessary to produce sufficient adsorption sites. Also i t is possible that adsorption could result in either bond rupture or dissolution. With reference to Fig. 50 adsorption on atom A could weaken bonds AB, AC and AD by causing a redistribution of the electrons surrounding atom A. This could then result in either the rupture of bonds AB or AC or else in the dissolution of any of atoms A, B or C. Thus, the crack could propagate by adsorption induced bond rupture or adsorption assisted dissolution. -107-O M E T A L A T O M X A D S O R B E D SPECIES FIGURE 50 Schematic representation of the effect of adsorption at the crack tip. (See text for explanation). -108-5. CONCLUSION 5.1 Conclusions The present studies on stress corrosion cracking of 316 and 310 stainless steels, and a TRIP steel in hot aqueous magnesium chloride solutions have led to the following conclusions! (i) V-K^plots of these ductile alloys are best obtained in the cold worked condition with a variation of the double canti-lever beam specimen whose behaviour Is similar to a 3 point bend specimen. (li) The alloys exhibit both stress intensity dependent cracking (Region I) and stress intensity independent cracking (Region II). The cracking rates in Region II at 154°C. of the 316 and TRIP steels (approximately 3.5 x 10 m/s) are considerably faster than that of the 310 (1.01 x 10" 8 m/s). (iii) Crack branching occurs in both Regions I and II, particularly in the annealed condition and when the cracking mode is transgranular. (iv) Cracking in 25$ cold rolled 316 at 154°C. does not occur below a critical potential whose value is dependent. Above this -value cracking is independent of potential over a range of 50mV in Region I and ?5mV in Region II. (v) In Region II, the crack velocities of a l l the alloys obey an Arrhenius rate law, the apparent activation energies varying from 15.1 kcal/g.mole to 18.1 kcal/g.mole. (vi) The corrosion potential of the 25$ cold rolled 316 alloy cracking at 154°C. becomes more active as the stress intensity is increased consistent with a shift of the anodic polarization curve to higher currents. -109-(vil) The amount of Intergranular cracking decreases with increasing alloy content and in 25$ cold rolled 316 i t in-creases with stress intensity and cathodic polarization and is greater in the centre of the crack front. The effects of and position along the crack front may he attributed to a change in potential at the crack tip. (viii) Temperature does not affect the fractography of the 316 and 310 alloys, but at lower temperatures there are nodular features on the fracture surface of the TRIP steel which may be related to the presence of martensite. (ix) While this work does not exclude a film rupture/dissolution mechanism the results are most readily explained by a process in which an adsorbed species modifies the behaviour of surface atoms at the crack tip. 5'2 Suggestions for Future Work (i) Determine quantitatively the effects of composition, cold work and temperature on the Region I kinetics. This may provide some insight Into the rate controlling processes in Region I. (ii) Determine the effect of varying the solution concentration on the cracking rates in Regions I and II and on the activation energy for Region II. This could help to distinguish between dissolution and adsorption controlled cracking. -110-BIBLIOGRAPHY 1. Nielsen, N.A. Corrosion, Vol. 27, p. 173 (1971). 2. Brown, W.F. and Srawley, J.E. Plane Strain Crack Toughness Testing of High Strength Materials, p. 6, A.S.T.M. S.T.P. 410 (1966). 3. Wiederhorn, S.M. J. Amer. Ceramic Soc, Vol. 50, p. 407 (1967). 4. Wiederhorn, S.M. Int. J. of Fract. Mech., Vol. 4, p. 171 (1968). 5. Speidel, M.O. and Hyatt, M.V. Advances in Corrosion Science and Technology, Vol. 2, pp. 115-335» Plenum Press, New York (1972). 6. Speidel, M.O. Met. Trans, Vol. 6A, p. 631 (1975). 7. Blackburn, M.J., Feeney, J.A., and Beck, T.R. Advances in Corrosion Science and Technology, Vol. 3, Plenum Press, New York (1973). 8. Sandoz, G. S.C.C. in High Strength Steels and in Titanium and Aluminum Alloys, pp. 79-145, U.S. Gov. Printing Office (1972). 9. McEvily, A. and Bond, A. J. Electrochem Soc, Vol. 112, p. 131 (1965) 10. Speidel, M.O. Corrosion, Vol. 32, p. 187 (1976). 11. Helfrich, W.J. Stress Corrosion Testing, pp. 21-30, A.S.T.M. S.T.P. 425 (1966). 12. Hodge, J.G., and Miller, J.L. Trans. A.S.M. Vol. 28, p. 25 (1940). 13. Osterholm, R. Paper Trade Journal, p. 26, Oct. 1975. 14. Van Droffelaar, H., and Hudson, G. Canadian Chemical Processing, Vol. 50, p. 78, FEB. and p. 63 MAR. (1966). 15. Grafen, H. Werkstoffe und Korrosion, Vol. 16, p. 344 (1968). 16. Snowden, P.P. J.I.S.I., Vol. 197, P. 136 (1961). 17. Harston, J.D., and Scully, J.G. Corrosion, Vol. 25, p. 493 (1969). 18. Latanision, R.M., and Staehle, R.W. Proc of Conf. on Fund. Aspects of Stress Corrosion Cracking, pp. 214-296, N.A.C.E. (1969). 19. Staehle, R.W. The Theory, of Stress Corrosion Cracking in Alloys, p. 223, J.C. Scully - Ed., N.A.T.0. (1971). -111-20. Copson, H.R. Physical Metallurgy of Stress Corrosion Fracture, p. 247, T. Rhodin - Ed., Interscience, New York (1959). 21. Edeleanu, C. J. of the Iron and Steel Industry, Vol. 175, P. 390 (1953). 22. Blrley, S. Phd. Thesis 1972, University of British Columbia. 23. Baker, H.R., Bloom, M., Bolster, R., and Singleterry, C. Corrosion, Vol. 26, p. 420 (1970). 24. Swann, P.R. Corrosion, Vol. 19, p. 102t (1963). 25. Swann, P.R., and Nutting, J. J. of Inst, of Metals, Vol. 88, p. 477 (I960). 26. Swann, P.R., and Embury, J.D. High Strength Materials, Zackay - Ed., Wiley, New York (1965). 27. Douglas, D.L., Thomas, G., and Rosser, W. Corrosion, Vol. 20, p. 15t (1964). 28. Vermilyea, D.A. Physics Today, p. 23, SEPT 1976. 29. Nikiforuk, T.P. Masters thesis 1976, University of British Columbia. 30. Morek, M., and Hochman, R. Corrosion, Vol. 26, p. 5 (1970). 31. Rhodes, P.R. Corrosion, Vol. 25, p. 462 (1969). 32. Barnartt, S., and Van Rooyen, D. J. of the Electrochem Soc, Vol. 108, p. 222 (1961). 33. Holzworth, M.L., and Louthan, M.R. Corrosion, Vol. 24, p. 110 (1968). 34. Vaughn, D.A., Phalen, D., Peterson, C , and Boyd, W. Corrosion, Vol. 21, p. 315 (1965). 35. Birley, S.S., and Tromans, D. Corrosion, Vol. 27, p. 63 (1971)• 36. Holzworth, M.L. Corrosion, Vol. 25, p. 107 (1969). 37. Mehta, M.L., and Burke, J. Corrosion, Vol. 31, p. 108 (1975). 38. Zackay, V., Parker, E.., Fahr, D., and Bursch, R. A.S.M. Trans. Quart., Vol. 60, p. 252 (1967). 39. McCoy, R.A., and Gerberch, W.W. Met. Trans, Vol. 4, p. 539 (1973). 40. Sedriks, A.J. Corrosion, Vol. 31f p. 339 (1975). -112-41. Irwin, G.R., and Kies, J.A. Welding Journal Research Supplement, Vol. 33, P. 193 (1954). 42. Knott, J.P. Fundamentals of Fracture Mechanics, pp. 108-125, Butterworths, London (1973). 43. Irwin, G.R. J. Appl. Mech., Vol. 24, p. 361 (1957). 44. Gross, B., and Srawley, J.E. Stress Intensity Factors for Three Point Bend Specimens by Boundary Collocation, Technical note D-3092, N.A.S.A. DEC. 1965. 45. Smith, H.R., and Piper, D.E. p. 65 of Ref. # 7. 46. Blau, P.J. Met. Trans. Vol. 7A, p. 463 (1976). 47. Magnani, N.J. Corrosion, Vol. 31# p. 337 (1975). 48. Streicher, M.A., and Ga3ale, I.B. p. 305 of Ref. # 18. 49. Wilde, B.E. J. of Electrochem. Soc., Vol. 118, p. 1717 (1971). 50. Speidel, M.O. p. 35 of Ref. # 19. 51. H i l l l g , W.B., and Charles, R.J. High Strength Materials, p. 682, V. Zackay - Ed., Wilely, New York (1965). 52. Wiederhorn, S.M., and Bolz, L.H. J. of Amer. Ceramic Soc. Vol. 53, P. 543 (1970). 53. Brennert, S. Recent Advances i n Stress Corrosion, Royal Swedish Acadamy of Eng. Sciences, Stockholm (1961). 54. Wilde, B.E., and Kim, G.D. Corrosion, Vol. 28, p. 350 (1972). 55. Inglis, C.E. Trans. Inst, of Naval Archit,, Vol. 55, p. 2 (1913). 56. Glasstone, S., Laidler, K., and Eyring, H. Theory of Rate Processes, McGraw H i l l , New York (1941). 57. West, J.M. Electrodeposition and Corrosion Processes, p. 39, D. Van Nostrand, London (I965). 58. Vermilyea, D.A., and Diegle, R.B. Corrosion, Vol. 32, p. 26 (1976). 59. Uhllg, H.H. pp. 86-91 of Ref. # 18. 60. Uhllg, H.H., and Cook, E.W. J. Electrochem. S o c , Vol. 117, p. 18 (1970). 61. Latimer, W. Oxidation Potentials, Prentice Hall, New York (1952). -113-62. Austin, J.H., and Elleman, T.S. J. of Nuclear Mat. Vol. 43, p. 119 (1972). 63. Jost, W. Diffusion in Solids, Liquids and Gases, pp. 305-307, Academic Press, New York (1952), 64. Barnartt, S., and Van Rooyen, D. J. of Electrochem. Soc, Vol. 108, p. 222 (1961). 65. Okada, H., Hosoi, Y., and Abe, S. Corrosion, Vol. 27, p. 424 (1971). 66. Payer, J.H., and Staehle, R.W. Corrosion Fatigue - Chemistry, Mechanics and Microstructure, p. 211, N.A.C.E. (1971). 67. Beck, T.R. p. 64 of Ref. # 19. 68. Chemical Engineers Handbook, J.H. Perry - Ed., McGraw H i l l , New York (1941). 69. Scully, J.C. p. 127 of Ref. # 19. 70. Streicher, M.A. Corrosion, Vol. 30, p. 77 (1974). 71. Sugimoto, K., and Sawada, Y. Corrosion, Vol. 32, p. 347 (1976). 72. Foley, R.T. J* Electrochem. Soc, Vol. 122, p. 1493 (1975). -114-APPENDIX 1 Calculation of Errors due to the Downward Growth of the Cracks From Fig. Alt sin 3 ~ ?! CB - a sin 3 At) C A cos 3 -£g AC •» a cos 3 OD - OB - (OC2 + CB 2)^ ' - [ ( R + a cos 3 ) + (a sin 3 ) 2 ] ^ - (R2 + 2aR cos 3 + a 2 ) ^ Error i n a « A B - A D « a - (OD - OA) 2 2-4-i.e., a « a + R - (R + 2aR cos 3 + a )* For 3 constant, the fractional error in da/dt, i.e., in V, is the same as the fractional error in a. The % error in V for different values of a and 3 is shown in Table Al. The % error in due to the increase in length of L is equal to the % increase in L, since Oc L, see Eqn. (4). From Fig. Al. AL - CB - a sin 3 AL/L «• a sin 3/L Values of AL/L for different a and 3 values are shown in Table Al. The error in due to the actual ligament length BE, being greater than the measured value DF, has been calculated as the difference ln due to horizontal cracks of lengths AC and AD, i.e., of lengths a cos 3 and [R + 2aR cos 3 + a ) * - R] respectively. The % error in K^ , obtained in this way is shown in Table Al for different values of a and 3. The net error in due to both of the above errors is also shown in Table Al. -115-F-IGURE A l Effect of the downward growth of the cracks on the measurements of V and K^. AB crack length of a 0 a t angle between the crack and the horizontal 0 - centre of scribed scale DB - arc subtended by 0 AD m distance measured as crack length BE m actual ligament length DP m measured ligament length CB s the error in moment arm L OA » distance R TABLE Al Errors due to the Downward Growth of the Cracks Crack Length (cm.) % Error ln Velocity % Error in due to L % Error in YL^ due to Ligament Error Net % Error in Kx 10° 0.5 - 1 . 3 -1 .4 +0.1 - 1 . 3 1.0 - 1 . 1 -2 .7 +0.47 -2 .2 2.0 -0 .9 - 5 . 5 +1.5 -4 .0 20° 0.5 -5 .2 -2 .7 +0.52 -2 .2 1.0 - 4 . 5 -5 .4 +1.81 -3 .6 2.0 -3 .6 -10.8 +5.7 - 5 . 1 N.B. negative errors indicate that the measured value is less than the actual value 

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