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Embrittlement of brass by ammoniacal solutions and mercury Birley, Stuart Samuel 1970

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EMBRITTLEMENT OF BRASS BY AMMONIACAL SOLUTIONS AND MERCURY by STUART S. BIRLEY B.Sc. Dept. of I n d u s t r i a l Metallurgy U n i v e r s i t y of Birmingham, 1965 A THESIS SUBMITTED IN PARTIAL FULFILMENT OF THE REQUIREMENTS FOR THE DEGREE OF MASTER OF APPLIED SCIENCE i n the Department of Metallurgy We accept t h i s thesis as conforming to the required standard THE UNIVERSITY OF BRITISH COLUMBIA January, 1970 In p r e s e n t i n g t h i s t h e s i s in 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 a d v a n c e d d e g r e e a t t h e 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 h e 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 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 b e . g r a n t e d by t h e 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 Metallurgy  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 M a r c h 12, 1970 FOR LOUISE ABSTRACT The influence of a l i q u i d metal and a stress corrosion environment on the mechanical properties of a and 3 brasses was investigated under continuous t e n s i l e loading conditions: s t r a i n rate and grain s i z e (a brass only) were systematic v a r i a b l e s . Increasing s t r a i n rate or decreasing grain s i z e was found to increase the d u c t i l i t y and fracture stress of the. p o l y c r y s t a l l i n e material i n eit h e r environment. Single c r y s t a l studies revealed (1) that grain boundaries are not e s s e n t i a l for embrittlement by either media and (2) that the surface films induced by the environment are mechanically very weak. The fractured surfaces of a brass were examined (1) f o r topographical features using both d i r e c t and i n d i r e c t r e p l i c a electron microscopy and (2) for evidence of t h i n films using a low angle electron d i f f r a c t i o n technique. Crack path i n both environments was i n v a r i a b l y intergranular, and d e t a i l s of fractured surfaces were s i m i l a r . Thin films were detected on the fractured surfaces, and the compositions determined. In general, both environments conferred the same general e m b r i t t l i n g e f f e c t s . I t i s possible to account f o r the current observations by a common cracking mechanism: the development of such a model based on the s l i p step displacement of a passive surface f i l m i s discussed. ACKNOWLEDGEMENTS I am indebted to my research d i r e c t o r , D. Tromans for h i s u n f a i l i n g enthusiastic guidance and encouragement throughout the work. I also extend sincere thanks to my collegues, members of Centre for Materials Research, Research Associates, Staff and Faculty members of the Metallurgy Department for t h e i r advice, to the workshop s t a f f for the i r e f f i c i e n t specimen machining s e r v i c e , to the s t a f f of Mineral Engineering Department, and to a l l others who assisted with the preparation of th i s t h e s i s . I g r a t e f u l l y acknowledge f i n a n c i a l assistance from the National Research Council (grant number A-4534) and the Defence Research Board (grant number 9535-50). ( i i i ) TABLE OF CONTENTS Page No. 1. INTRODUCTION 1 1.1. Analogies between l i q u i d metal embrittlement and stress corrosion cracking 1 1.1.1. Influence of stress 1 1.1.2. E f f e c t of impurities and a l l o y i n g elements on cracking s u s c e p t i b i l i t y 3 1.1.3. S p e c i f i t y of environment 3 1.1.4. Path of cracking 4 1.1.5. Grain s i z e dependence 5 1.1.5. Temperature e f f e c t s 6 1.2. Models applicable to both stress corrosion cracking and l i q u i d metal embrittlement 7 1.2.1. D i s s o l u t i o n models 7 1.2.2. Tarnish and i n t e r m e t a l l i c rupture models 9 1.2.3. Surface energy concepts i n environmental stress assisted fracture 11 1.2.3.1.Evaluation of surface energy 11 1.213.2.Influence of p l a s t i c deformation 13 1.3. Scope of current i n v e s t i g a t i o n 16 2. EXPERIMENTAL 20 2.1. Materials 20 2.2. Tensile specimen preparation 21 2.2.1. P o l y c r y s t a l l i n e a-brass 21 2.2.2. P o l y c r y s t a l l i n e B-brass 22 2.2.3. Single c r y s t a l s a-brass 22 2.2.4. Single c r y s t a l s B-brass 23 2.3. Tensile t e s t i n g procedure and conditions 23 (iv) Table of Contents (Contd), Page No. 2.3.1. Testing procedure 23 2.3.2. Testing conditions 25 2.3.2.1. P o l y c r y s t a l l i n e Materials 25 2.3.2.2. Single c r y s t a l s s 26 2.4. Fractography 26 2.5. R e f l e c t i o n electron microscopy 27 2.6. Metallography 27 2.6.1. Crack path determination 27 2.6.2. Intergranular penetration determination 27 3. RESULTS AND OBSERVATIONS 30 3.1. Single c r y s t a l t e n s i l e tests 30 3.2. P o l y c r y s t a l l i n e t e n s i l e tests 31 3.2.1. a-brass 31 3.2.1.1. S t r a i n rate e f f e c t s 31 3.2.1.2. Grain s i z e e f f e c t s 32 3.2.1.3. Incremental loading tests 33 3.2.1.4. Penetration tests 34 3.2.2. B-brass 35 3.2.2.1. S t r a i n rate e f f e c t s 35 3.2.2.2. Grain s i z e e f f e c t s 35 3.3. O p t i c a l metallography 35 3.3.1. Crack path 35 3.3.2. Penetration 36 (v) Table of Contents (Contd). Page No. 3.4. R e f l e c t i o n electron microscopy 36 3.5. Fractography 37 3.5.1. P o l y c r y s t a l l i n e a-brass at Room temperature 37 3.5.2. P o l y c r y s t a l l i n e a-brass - high temperature Hg 38 3.5.3. P o l y c r y s t a l l i n e 3-brass - room temperature 38 3.5.4. Single c r y s t a l s 38 3.6. Analogies of t e n s i l e behaviour i n mercury and Mattsson's s o l u t i o n - a summary 39 4. DISCUSSION 90 4.1. Discussion of aforementioned cracking mechanisms i n terms of present observations 90 4.1.1. Surface energy mechanism 90 4.1.2. Tarnish rupture mechanism 92 4.1.3. Di s s o l u t i o n Processes 94 4.2. Development of a cracking mechanism 94 4.2.1. Predictions of the cracking mechanism 97 4.2.2. Quantitative elements of a cracking mechanism 101 5. CONCLUSIONS 110 BIBLIOGRAPHY ' 111 APPENDIX I 116 APPENDIX II 117 APPENDIX III 122 (vi) LIST OF FIGURES No. Description Page 1. P o t e n t i a l Energy vs. Interatomic separation 18 2. Film Rupture mechanism 19 3. Tarnish rupture mechanism 19 4. Copper-Zinc phase diagram " 28 5. Specimen dimensions and d e s c r i p t i o n 29 6. Resolved shear stress vs. shear s t r a i n f o r a-brass s i n g l e c r y s t a l s deformed i n neutral Mattsson's s o l u t i o n 40 7. Resolved shear stress vs. shear s t r a i n f o r a-brass s i n g l e c r y s t a l s deformed i n mercury 41 8. Resolved shear stress vs. shear s t r a i n for B-brass sin g l e c r y s t a l deformed i n Mattsson's s o l u t i o n (pH 6.8) 42 9. Influence of mercury on the load extension curve of B-brass sin g l e c r y s t a l s 43 10. Load-extension curves of a-brass deformed i n neutral Mattsson's s o l u t i o n at a seri e s of crosshead speeds 44 11. Load-extension curves of a-brass deformed i n Mattsson's s o l u t i o n at a seri e s of crosshead speeds (20u grain dia.) 45 12. Load-extension curve for a-brass deformed i n Mattsson's s o l u t i o n (pH 6.8) at a seri e s of loading rates (large grain size) 46 13. Load-extension curves of a-brass deformed i n mercury as function of grain s i z e and s t r a i n rate 47 14. D u c t i l i t y of a-brass vs. pH of Mattsson's s o l u t i o n for a seri e s of crosshead speeds 48 15. Influence of environment on the r e l a t i o n s h i p between d u c t i l i t y of a-brass and s t r a i n (20u grain dia.) 49 ( v i i ) L i s t of Figures (Contd).. No. Des c r i p t i o n Page 16. Elongation vs. s t r a i n rate as function of grain diameter for a-brass i n mercury 50 17. Elongation at crack i n i t i a t i o n vs. s t r a i n rate f o r a-brass deformed i n Mattsson's s o l u t i o n and mercury 51 18. Load extension curves of a-brass deformed i n several environments for two crosshead speeds (20p grain dia.) 52 19. E f f e c t of grain s i z e on load extension curve of a-brass i n ambient a i r 53 20. E f f e c t of grain s i z e and environment on the i n i t i a l stages of the load extension curves of d-brass deformed at a crosshead speed of 0.001"/min. 54 21. E f f e c t of grain s i z e and environment on the load-extension curves of a-brass deformed at crosshead speed^-of 0.001"/min. 55 22. V a r i a t i o n of frac t u r e stress and flow stress with grain s i z e f o r a-brass i n a i r 56 23. Fracture stress vs. Inverse square root of grain d i a . of a-brass embrittled by Mattsson's s o l u t i o n . 57 24. Fracture stress vs. inverse square root of grain d i a . for a-brass embrittled by mercury at d i f f e r e n t s t r a i n rates 59 25. Incremental loading tests performed on three p o l y c r y s t a l l i n e a-brass specimens i n mercury 60 26. Load r e l a x a t i o n curves of a-brass i n mercury and Mattsson's s o l u t i o n 62 27. Load extension curves of B-brass deformed i n water at a series of crosshead speeds 63 28. Load extension curves of B-brass deformed i n several environments at crosshead speeds of 0.1'Vmin and 0.01"/min. 64 L i s t of Figures (Contd). No. Description Page 29. Influence of mercury on the load elongation curve of p o l y c r y s t a l l i n e 3-brass f or several crosshead speeds 65 30. D u c t i l i t y of 3-brass vs. pH for several crosshead speeds 66 31. Influence of environment on the r e l a t i o n s h i p between d u c t i l i t y of 3-brass and crosshead speed 67 32. Intergranular cracking of a-brass i n mercury Mattsson's solution (X120) 68 33. Ring Pattern - Gold standard 69 34. Ring Pattern - Mercury fractured a-brass at 100°C 69 35. Spot Pattern - Mercury fractured a-brass at 100°C 69 36. Ring Pattern - Mattsson's s o l u t i o n fractured a-brass 70 37. As 36, showing broad ri n g 70 38. Ring Pattern - Tarnish 71 39. Ring Pattern - Tarnish 71 40. Du c t i l e fracture dimples i n a i r fractured a-brass (X600) 75 41. Elongated fr a c t u r e dimples i n a i r fractured a-brass (X600) 75 42. Intergranular cracking of a-brass i n Mattsson's s o l u t i o n (X300) 76 43. Intergranular cracking of a-brass i n mercury (X300) 76 44. Ductile/intergranular t r a n s i t i o n region i n Mattsson's s o l u t i o n fractured a-brass 77 45. a-brass fractured i n mercury at ambient temperature (X400) 78 L i s t of Figures (Contd). (ix) No. 46. 47, Description a-brass fractured i n mercury at ambient temperature (X16,000) a-brass f r a c t u r e d ' i n Mattssori's s o l u t i o n (pH6.8)-RT (X2000) Page 78 79 48. a-brass fractured i n Mattsson's s o l u t i o n (pH6.8)-RT (X5800) 79 49. 50. 51. 52. a-brass fractured i n Mattsson's s o l u t i o n (PH6.8)-RT (X4,000) 79 a-brass fractured i n Mattsson's s o l u t i o n (pH6.8)-RT (X10,200) 79 Intergranular cracking i n Mattsson's s o l u t i o n (XI,500) 80 Transgranular region i n Mattsson's s o l u t i o n (X2,000) 80 53. Fine background d e t a i l i n Mattsson's.solution (X5,800) 80 54. Grain boundary product i n Mattsson's s o l u t i o n (X6,500) 80 55. Intergranular fracture of a-brass by Hg at 100°C (X1500) 81 56. 57. Intergranular f r a c t u r e and s t r i a t i o n s i n a-brass deformed i n Hg at 100°C (X4,000) Surface Features of a-brass fractured i n Hg at 100°C (X2000) 81 82 58. Surface Features of a-brass fractured i n Hg at 100°C (X8000) 82 59. Surface features of a-brass fractured i n Hg at 100°C (X2000) 82 60. Surface features of a-brass fractured i n Hg at 100°C (X2000) 82 (x) L i s t of Figures (Contd). No. Description Page 61. Surface features of a-brass fractured i n Hg at 100°C (X2000) 83 62. Surface features of a-brass fractured i n Hg at 100°C (X10,000) 83 63. Probe topography of 100°C mercury fractured a-brass showing surface features at seri e s of magnifications 84 64. Zinc X-ray scan by probe of area i n F i g . 65 85 65. Area containing mercury globule and pyramids 85 66. Mercury X-ray scan by probe of area i n f i g . 65. 85 67. Intergranular cracking of B-brass by mercury (X255) 86 68. Cracking of B-brass i n Mattsson's s o l u t i o n (X300) 87 69. Cracking of B-brass i n Mattsson's s o l u t i o n (X300) 87 70. Low Mag photo, of a-brass s i n g l e c r y s t a l s 88 71. Low mag. photo, of B-brass sin g l e c r y s t a l s 89 72. Topography of fractured surface of specimen B i n f i g . 71 89 73. Topography of fractured surface of specimen B i n f i g . 71 (X225) 89 74. Topography of fractured surface of specimen B i n f i g . 71 (X225) 89 75. Cracking time vs. pH (load relaxation)- : 104 76. Film shearing and parting process 105 77. Model for cracking 105 78. E f f e c t of s l i p step on exposure of fresh metal and oxide thickness 106 79. Notch geometry for two grain sizes 107 80. Cracking time vs. temperature 108 (xi) L i s t of Figures (Contd). No. Description Page 81. Geometrical/Mathematical model of c r i t i c a l notch 107 82. Log G G vs. log time p l o t s 109 A l . Fatigue l i f e v a r i a t i o n with pH. 121 ( x i i ) LIST OF TABLES No. Description Page I E f f e c t i v e surface energies (after Rostoker) 58 II Penetration/Immersion test r e s u l t s 61 III Analysis of d i f f r a c t i o n rings -Hg at 100°C 72 IV Analysis of d i f f r a c t i o n rings - Mattsson's s o l u t i o n 73 V Analysis of d i f f r a c t i o n rings - v i s i b l e t a r nish 74 VI A c t i v a t i o n energy data 99 VII Calculations of x c from experimental data 103 A l Fatigue conditions 119 A2 Test Results 120 1. INTRODUCTION S p e c i f i c environments promote the premature cracking of appro-p r i a t e metals and non-metals under su i t a b l e conditions of t e n s i l e s t r e s s . The influence of environment p a r t i c u l a r l y on the behavior of stressed metals i s a long standing concern of industry since normally d u c t i l e material may f a i l i n a catastrophic manner i n the presence of the unique embrittling medium. H i s t o r i c a l l y , e arly studies were exc l u s i v e l y confined to copper base a l l o y s because of the acute problem of "season cracking" and " f i r e - c r a c k i n g " to the brass trade, but increasing attention i s being focused on high strength s t e e l s , s t a i n l e s s s t e e l s , and aluminum and titanium a l l o y s , important engineering materials of today. Many s i m i l a r i t i e s between the cracking of stressed metals i n aqueous environments and i n l i q u i d metals have been recognised (1,2,3) o r i g i n a t i n g the suggestion (2) that both stress corrosion cracking and l i q u i d metal embrittlement enjoy a common cracking mechanism. 1.1. Analogies between l i q u i d metal enbrittlement and stress corrosion cracking Influence of stress 1.1.1. - 2 -It i s generally acknowledged that both stress and a unique environment are required c o n j o i n t l y f or the occurrence of l i q u i d metal embrittlement or stress corrosion cracking. Stresses either externally applied or r e s i d u a l from cold working (4) or welding (5) operations may induce cracking provided they possess t e n s i l e components. Uniaxial compressive stresses do not encourage the phenomena (6). Although plastic p r e s t r a i n has been shown to influence the cracking of an age-hardenable aluminium a l l o y when immersed under stress i n a s a l i n e s o l u t i o n or mercury (2), macroscopic p l a s t i c deformation does not appear to be a p r e r e q u i s i t e f or the occurrence of cracking since fracture has been observed i n annealed brass (2) and other a l l o y s (7) stressed below t h e i r y i e l d points. The time to f a i l u r e does, however, decrease with increasing magnitude of applied t e n s i l e stress and such " s t a t i c f a t i g u e " behavior i s a common a t t r i b u t e of l i q u i d metal embrittlement and stress corrosion cracking (2). Under continuous t e n s i l e loading conditions, the l i q u i d metal or aqueous environment has no e f f e c t on the y i e l d point and the work •= hardening rate of p o l y c r y s t a l l i n e material (8), but serves to l i m i t the elongation at which fracture occurs. Decreasing the loading rate i n l i q u i d metals (9) further reduces the d u c t i l i t y which r e f l e c t s the observed decrease i n fracture s t r e s s ; however, e f f e c t s of s t r a i n rate have not been reported for tests performed i n stress corrosion media. Very l i t t l e work has been concerned with s i m i l a r e f f e c t s on si n g l e c r y s t a l s . In fatigue, 70/30 brass exhibits reduced l i f e i n mercury (10). and ammoniacal environments (11). - 3 -1.1.2. E f f e c t of impurities and a l l o y i n g elements on s u s c e p t i b i l i t y  to cracking S u s c e p t i b i l i t y of a metal to embrittlement i s profoundly i n f l u -enced by the presence of impurities and a l l o y i n g elements. Uhlig (12) has asserted that pure metals are completely immune to stress cracking and although only one instance of cracking of pure magnesium (13) and pure copper (14) has been reported, high p u r i t y , aluminium has never been known to f a i l . Bailey (15) suggests that the cracking tendency i s very minute but not n i l , i n f e r r i n g that f a i l u r e may occur on prolonged immersion i n the environment. In l i q u i d metals, however, pure metals f a i l r e a d i l y (16) at high stresses. An addition of impurities or a l l o y i n g elements while i n v a r i a b l y increasing the cracking tendency of pure metals, may e i t h e r retard, stimulate, or unchange the attack of a l l o y s depending on the a l l o y / environment system. "Pure" 70/30 Cu/zinc i s just as vulnerable as commercial a-brass to ammonia or mercury (17). Increasing the zinc content of brass increases the cracking tendency (18) and small additions of P, As, Sb, S i , Al and Ni to pure copper form a l l o y s which are prone to enbrittlement i n ammonia solutions, but further additions of the same elements (19) improve resistance to f a i l u r e . Addition of 36 separate elements to an a brass was performed by Wilson (20) and not one was found to have an accelerating e f f e c t on the cracking, supporting observations of other workers (18,21,22). 1.1.3. S p e c i f i t y of environment Comparatively few l i q u i d metals and aqueous solutions promote the cracking of stressed metals and a l l o y s , and the action of an environment - 4 -i s u sually s p e c i f i c to one p a r t i c u l a r metal or group of a l l o y s having a common base solvent, or confined to a small number of such a l l o y groups. There i s no si n g l e chemical or l i q u i d metal that causes cracking of a l l metals and a l l o y s , and no one metal or a l l o y i s susceptible to a l l environments known to induce cracking. Hence ammonia vapours and solutions (15) are s p e c i f i c a l l y embrittling to copper base a l l o y s , l i q u i d gallium (9) to aluminum, and l i q u i d bismuth (2) to copper, whilst sodium chloride (23) w i l l embrittle aluminum base and magnesium base a l l o y s and some s t e e l s , and l i q u i d metal mercury attacks copper and copper base-, aluminum base- and titanium base-alloys and c e r t a i n s t e e l s (9). 1.1.4. Path of cracking Stress corrosion cracking may occur i n a t r a n s c r y s t a l l i n e or intergranular fashion depending on the system involved. Embrittle-ment produced by l i q u i d metals however i s almost i n v a r i a b l y intergranular, but t r a n s c r y s t a l l i n e cracking has been noted once i n a p o l y c r y s t a l l i n e aluminum a l l o y stressed u n i a x i a l l y i n a mercury-zinc l i q u i d amalgam (24), and i s possible i n si n g l e c r y s t a l s . It i s well: established that a-brass i s embrittled i n an i n t e r -c r y s t a l l i n e fashion by mercury i n constant load conditions but the, s i t u a t i o n f o r ammonia i s less s e t t l e d . Cracking i n ammonia i s pre-dominantly intergranular but a change to the transgranular mode has been reported when the zinc content exceeds 30% (19), or when a t h i r d element such as aluminum (5,25) i s present. Mattsson (26) however demonstrated that the path of cracking of a binary copper -37.2% zinc a l l o y i n an ammonical copper sulphate so l u t i o n i s dependent on the pH - 5 -of the environment, intergranular f a i l u r e occurring i n the pH range 6.3-7.3 with transgranular cracking at both lower and higher pH values. Although Hoar and Booker (27) supported Mattsson's thesis that i n t e r -granular attack i s associated with the formation of the black oxide tarnish which i s exclusive (in the above system) to the aforementioned pH range, Johnson and Leja (28), Tromans et a l , (29) and L a h i r i (30) have obtained intergranular cracking at pH values exceeding 7.3, con-cluding that t a r n i s h formation i s not a p r e r e q u i s i t e of such a fracture mode. Tensile tests on p o l y c r y s t a l l i n e material i n stress corrosion :'X media were f i r s t reported by Coleman et a l (3) who concluded that crack i n i t i a t i o n was incident at the grain boundary/surface i n t e r f a c e as analogous to m e t a l / l i q u i d embrittlement couples (9). However, crack i n i t i a t i o n i s not n e c e s s a r i l y dependent on grain boundaries as indicated by the stress corrosion cracking of a-brass si n g l e c r s y t a l s i n ammoniacal environments (12,18) and the embrittlement of zinc, cadmium and t i n s i n g l e c r y s t a l s by a v a r i e t y of l i q u i d metals (31). 1.1.5. Grain s i z e dependence S u s c e p t i b i l i t y to cracking of u n i a x i a l l y stressed a-brass i n mercury or ammonia increases with grain s i z e (9,12,18) and, i n f a c t , Edmunds (18) has observed a d e f i n i t e grain s i z e dependency of f r a c t u r e s t r e s s . The e f f e c t of cold work on fracture through a change i n grain diameter i s influenced by the type and the amount of deformation, and i n addition contributions from r e s i d u a l stresses, d i s t o r t i o n of grain shape and preferred o r i e n t a t i o n must be accounted f o r (15). - 6 -Although sin g l e c r y s t a l s of a-brass crack i n ammoniacal environ-ments (12,18), Edmunds (18) observed no e f f e c t i n mercury. However, t h i s single report of an a-brass monocrystal i n mercury remains to be substantiated. Several workers have observed a l i n e a r c o r r e l a t i o n of fracture stress to inverse square root of grain diameter i n the t e n s i l e t e s t i n g of copper and copper base all o y s i n l i q u i d metals, and of s t a i n l e s s s t e e l and magnesium al l o y s i n stress corrosion media. This e f f e c t has been analysed (2,3,8,9) i n terms of the Stroh-Petch equation for d i s l o c a t i o n nucleated cracks y i e l d i n g e f f e c t i v e surface energies of the correct order of magnitude for b r i t t l e fracture [N.B. Fracture s t r e s s , i n t h i s connection, r e f e r s to that stress given by the point of deviation of the s t r e s s - s t r a i n curve i n the environment from that curve i n a i r , and not to the stress at f i n a l fracture.] 1.1.6. Temperature e f f e c t s Generally, increasing the temperature of the system reduces the cracking time of stressed metals immersed i n aqueous, gaseous or l i q u i d media, with no e f f e c t occurring below the solidus i n the case of embrittlement by l i q u i d metals. [The observation that zinc s i n g l e , c r y s t a l s wetted with mercury undergo a d u c t i l e - b r i t t l e t r a n s i t i o n (9) at 150°C may be a t t r i b u t e d to the incidence of general corrosion.] Several attempts have been made to determine the a c t i v a t i o n energies f o r crack propagation. Hence i t has been proposed (32) that the migration of s i n g l e and double vacancies i s involved i n the cracking of unworked and worked aluminum allo y s r e s p e c t i v e l y , stressed i n a 3% NaCl s o l u t i o n i n the temperature range 25-70°C. Johnson (33) obtained a curve i n an Arrhenius pl o t f o r the Mattsson s o l u t i o n cracking of; ;brass - 7 -suggesting a gradual change i n one r a t e c o n t r o l l i n g step to another, but since the upper temperature was l i m i t e d to 63°C by the v o l a t i l i s a t i o n of ammonia no d e f i n i t e conclusions were made. An extremely low a c t i v a -t i o n energy (1000 cal/mole j has been noted f o r an aluminum a l l o y wetted with mercury i n the range -28°C to + 90°C (34) . Nichols and Rostpker (2) have observed the above mentioned general e f f e c t of temperature on the delayed f a i l u r e of 70/30 brass comparing a decrease of f a i l u r e time i n ammonia with a decrease of f r a c t u r e s t r e s s i n mercury. Other temperature e f f e c t s are considered i n s e c t i o n 4.2.1. : 1 . 2 ^ Y ; ., Models a p p l i c a b l e to both s t r e s s corrosion cracking and l i q u i d metal embrittlement A u n i f i e d mechanism of i n i t i a t i o n arid propagation f o r both st r e s s c o r r o s i o n cracking and l i q u i d metal embrittlement must account f o r the aforementioned observations [and, indeed, may be relevent to environ-mental a s s i s t e d fracture;:of glasses (35) and polymers (36)]. Consequently, i n s p i t e of the multitude of models proposed and reviewed (15,37-40) f o r s p e c i f i c systems by previous workers, only those capable of being applied to both l i q u i d metal embrittlement and s t r e s s c o r r o s i o n cracking w i l l be discussed. 1.2.1. D i s s o l u t i o n models D i s s o l u t i o n models were o r i g i n a t e d to account f o r the combined a c t i o n of st r e s s and c o r r o s i o n . - 8 -A strained metal generally corrodes more r a p i d l y than an un-strained one (41), and where no f i l m i s present on the metal surface accelerated corrosion might be caused by the f a c t that an atom i n a strained structure w i l l require less energy f o r d i s s o l u t i o n than one occupying a stable p o s i t i o n : t h i s i s i n f e r r e d from the f a m i l i a r potential-energy distance curve shown i n f i g u r e 1. P r e f e r e n t i a l attack might be expected at d i s l o c a t i o n s , grain boundaries and crack tips, but the energy d i f f e r e n t i a l due to s t r a i n i n g i s probably i n s u f f i c i e n t to induce s i g n i f i c a n t d i s s o l u t i o n . Anodic d i s s o l u t i o n models are more popular than the simple " s t r e s s - a s s i s t e d " d i s s o l u t i o n mechanism. Electrochemical attack may be associated with compositional heterogeneities induced at the metal surface by the action of s t r e s s , for example, by the segregation of solute to stacking f a u l t s (42), p r e c i p i t a t i o n on s l i p planes (43) or destruction of short range order (44) depending on the a l l o y under consideration. It has also been suggested (45) that an a l l o y contains highly anodic continuous paths e.g. grain boundaries and that stress acts as a tearing agent. By t h i s token such a material must have s u s c e p t i b i l i t y to intergranular corrosion i n the absence of s t r e s s , but t h i s conditions i s not always met (12) . Hoar, Hines and West (46,47) therefore propose that these i n t r i n s i c anodic areas are required only f o r i n i t i a t i n g a sharp crack, and that consequent pro-pagation i s incident by d i s s o l u t i o n of stress-induced anodic material at the crack t i p . Pardue et a l (48) considered fracture to consist of a s e r i e s of r e p e t i t i v e stages involving d i s s o l u t i o n - i n i t i a t i o n and short - 9. -range b r i t t l e crack propagation. Accelerated cracking of a l l o y s possessing a surface coating led to the development of f i l m rupture models (49,50), according to which the basic cause of cracking i s the l o c a l rupture of passive surface films by emergent s l i p steps (figure 2) and prevention of repassivation i n t h i s area by further p l a s t i c deformation. These mechanisms are not confined to transgranular cracking and i n f a c t , the f i l m i s more l i k e l y to be ruptured at grain boundaries. In summary, i t i s a general b e l i e f that the k i n e t i c s associated with present d i s s o l u t i o n models are not compatible with the observed rapid cracking r a t e s . However, anodic d i s s o l u t i o n mechanisms are i popular f o r the slower i n i t i a t i o n stage of f r a c t u r e . 1.2.2. Tarnish and i n t e r m e t a l l i c rupture models The mechanism of stress corrosion cracking i n a t a r n i s h producing s o l u t i o n i s believed (51) to be d i f f e r e n t from the f i l m rupture-d i s s o l u t i o n mechanism described above, occurring by the repeated formation and rupture of a t a r n i s h f i l m which forms within the metal (52,53). The b r i t t l e f i l m (figure 3A) forms at the specimen surface and i s fractured (3B) at a s u f f i c i e n t l e v e l of applied s t r e s s . Further s t r e s s i n g opens the crack, blunting the t i p (3C) and the environment enters to react with the f r e s h l y exposed metal to form l o c a l l y (3D) more ta r n i s h , which i n turn ruptures (3E). This process repeats u n t i l f i n a l mechanical fracture occurs. [Note that the f i l m involved i n t h i s mechanism i s of much greater thickness than that associated with the anodic f i l m rupture mechanism.] - 10 -This model may account f o r the observation of s t r i a t e d fractured surfaces (39) and the transgranular f a i l u r e of a-brass monocrystals i n ammonia (18) and i s consistent with p r e f e r e n t i a l grain boundary tarnish penetration (53) to propagate the intergranular cracking of a-brass. Tromans et a l (29) have pointed out, though, that a black tarnish does not form on the fractured surfaces of a c i r c u m f e r e n t i a l l y notched t e n s i l e specimen of p o l y c r y s t a l l i n e a-brass deformed i n Mattsson's sol u t i o n and therefore a t a r n i s h rupture model would not appear to apply to the crack propagation stage. Extending the model to l i q u i d metal embrittlement, cracking might occur as depicted i n fi g u r e 3 by the repeated formation and rupture of an i n t e r m e t a l l i c compound r e s u l t i n g from an i n t e r a c t i o n of the l i q u i d metal environment at the material surface. However, Rostoker (9) has the opinion that t h i s process would not occur r a p i d l y enough to account for the f a s t observed cracking rate. Also, cracking i s promoted above the temperature of formation of i n t e r m e t a l l i c compound (9) i n the brass mercury system. See Appendix I. Tarnish or i n t e r m e t a l l i c rupture may be only applicable to the crack i n i t i a t i o n stage: the p r e f e r e n t i a l grain boundary penetration (53) and subsequent b r i t t l e rupture of a t a r n i s h may account f o r grain boundary i n i t i a t i o n i n p o l y c r y s t a l l i n e a-brass. Occurrence of i n t e r -granular penetration i s quite common i n m e t a l / l i q u i d metal systems (9,54) and has been observed by Robertson (55) for the case of annealed a-brass contacted with mercury at 350°C. However, penetration was not induced i n copper or brass at room temperature or 170°C, and the e f f e c t was a t t r i b u t e d to the formation of an i n t e r m e t a l l i c compound which - 11 -prevented penetration. The observation by Greenwood (56) of mercury penetration into annealed copper at room temperature seems to b e l i e t h i s suggestion. 1.2.3. Surface energy concepts i n environmental stress a s s i s t e d  fr a c t u r e Surface energy mechanisms have been advanced to explain the phenomena of hydrogen embrittlement (57) stress corrosion cracking (3, 58) and l i q u i d metal embrittlement (59) . The proposed mechanism i s based on the concept that s p e c i f i c adsorption of surface active species from the environment w i l l decrease the i n t e r f a c i a l energy associated with the nucleation and propagation of a b r i t t l e crack, and, accordingly, might account f o r the s p e c i f i c i t y of media, the observed b r i t t l e crack paths, f a i l u r e i n tarnishing and non-tarnishing s o l u t i o n s , and fracture of non-metallic materials. Attempts have been made to examine the : e f f e c t of environment on the magnitude of surface energy, but i t i s d i f f i c u l t to obtain d i r e c t proof that the basic mechanism does operate. 1.2.3.1. Evaluation of surface energy The G r i f f i t h (60)equation °APP 7TC indicates the stress (CT^pp) required to propagate an e x i s t i n g e l a s t i c crack of length 2c i n a material of Youngs Modulus E. C l e a r l y , a reduction i n magnitude of the surface energy term y w i l l permit the crack - 12 -to extend under a smaller applied stress [y i s the energy of the s o l i d with respect to i t s own vapour i f fracture i n i t i a t e s i n the material whereas for surface nucleated cracks, y i s the i n t e r f a c i a l energy of the s o l i d with respect to i t s own oxide or other adsorbed f i l m ] . Surface energies associated with the crack propagation of glasses can be as low as 10% of the t o t a l fracture energy (61), the discrepancy a r i s i n g from the p l a s t i c energy expended i n the process. The s i t u a t i o n of surface energy being equal to fracture energy i s approached only i n the case of a few non-metallic materials (62) and, i n f a c t , has been used as a d e f i n i t i o n of b r i t t l e fracture (61). Thus for the cracking of metals, any determination of a reduction of magnitude of y by environmental e f f e c t s using t h i s approach must have contingency for p l a s t i c deformation. Considerations of the crack i n i t i a t i o n stage have also ignored the contribution from p l a s t i c energy. Rostoker and his colleagues (9) demonstrated d i r e c t l y that environmenta-lr induced b r i t t l e fracture of metals i s associated with reduced surface energy, by means of the Stroh-Petch (63,64) equation for d i s l o c a t i o n i n i t i a t e d cracks v i z . a a + KD -1/2 F o where a F fracture stress a functional constant o D K grain diameter 1-v G r i g i d i t y modulus poissons r a t i o - 13 -Rostoker employed dynamic loading tests and obtained p l o t s of fracture stress versus inverse square root of grain diameter. From the slopes of these p l o t s , y was estimated to be 280 ergs/sq cm for the embrittlement of brass by mercury, 157 ergs/sq cm for the fracture of a s t a i n l e s s s t e e l i n MgC^ s o l u t i o n and 93 ergs/sq sm for the cracking of a Mg/6A1 a l l o y i n NaCl-K^Cr^O^ s o l u t i o n . Since the estimated surface energies of these metals with respect to t h e i r own vapours ; are 1500, 1000 and 500 ergs/sq cm r e s p e c t i v e l y (3,64) i t would appear that a reduction i n surface energy had been ef f e c t e d by the presence of the corresponding environments. Similar calcuations have not been, performed f o r the fr a c t u r e of brass i n ammonia. 1.2.3.2. Influence of p l a s t i c deformation on surface energy The assumption that the surface energy"term, y, represents the energy to provide two f r e s h l y exposed undeformed surfaces i s not v a l i d , since as already indicated, fracture i s associated with l o c a l p l a s t i c deformation. Hence, y may be regarded as an apparent surface energy term to which p l a s t i c d i s t o r t i o n a l energy i s a major contributor, and the Orowan modification (65) to the G r i f f i t h equation of the form a = / 2 E(Y+P) APP TTC where P i s the p l a s t i c energy, has been applied. Since the estimates 3 6 of y and P are 10 and 10 ergs/sq cm r e s p e c t i v e l y , there appears no reason why the reduction of surface energy should d i r e c t l y reduce the stress required f o r b r i t t l e crack propagation. / 2|P TTC - 14 -S t o l o f f and Johnston (59) considered the events occurring i n the region of the crack t i p , as an a l t e r n a t i v e to the energy-balance approach adopted by Orowan and G r i f f i t h , proposing that adsorped species would reduce the cohesive strength of atoms def i n i n g the crack t i p to f a c i l i t a t e cracking by bond rupture, while increasing p l a s t i c deforma-t i o n would round o f f the crack p r o f i l e to retard cracking by decreas-ing the stress concentration. Their r e l a t i o n s h i p APP 4 T r a c o where p = crack t i p radius a Q = equilibrium l a t t i c e spacing does account for the decrease i n fracture stress by reduction of surface energy: although t h i s equation resembles the G r i f f i t h formula f o r the instance of the sharpest possible crack ( a Q = p), i t has the important d i s t i n c t i o n that the p l a s t i c deformation renders a dimension-less r a t i o ( — ) which m u l t i p l i e s , and not an energy term which i s a o added to, y« The model, however, contributes l i t t l e to the understanding of the basic embrittlement mechanism since no notice i s taken of other possible processes occurring at the crack t i p . For example, adsorbed species might d i f f u s e into the bulk metal and retard d i s l o c a t i o n motion l o c a l l y and thereby prevent a l t e r a t i o n of the p r o f i l e . On the other hand, d i s s o l u t i o n of the metal might increase or decrease the value of - 15 -p to decrease or increase, r e s p e c t i v e l y , the stress concentration a r i s i n g from the defect. A useful extension of frac t u r e theory i s being applied by design engineers f o r predicting the r e l i a b i l i t y of material already containing defects f o r use as castings, and fabricated vessels, girders, sheet and plate i n a l l types of environments. The Orowan modification to the G r i f f i t h equation f o r b r i t t l e crack propagation i s aAPP 7TC Rearranging:-CAPP ^ 1 / 2 = ^ = K e t l - A where K = stress i n t e n s i t y factor Substituting 2P by G, the "fracture toughness" and squaring:-K 2 = EG K i s a measure of the "fracture toughness", G, which i s the resistance of a material of a s p e c i f i c composition and structure to the propagation of an e x i s t i n g crack. The important feature i s that G i s independent of material dimensions f o r plane s t r a i n conditions and, thus, r e s u l t s of the e a s i l y experimentally determined K for specimens may be applied to larger s t r u c t u r a l components of i d e n t i c a l material to be used at the same temperature and equivalent s t r a i n r a t e s . From a knowledge of K, and equation A, the maximum allowable defect of s i z e 2c at the expected operative stress o- may be determined. : - 16 -In summary, modifications of fracture theory to allow f o r p l a s t i c deformation have been based s o l e l y on crack propagation c r i t e r i a and are unable to t e s t the hypothesis that reduced surface energy f a c i l i t a t e s environmental stress cracking, and moreover the present techniques of fracture mechanics despite a p r o f i c i e n c y f o r predicting the performance of defective material unfortunately shed no l i g h t on the fundamental mechanism of cracking although a recent a r t i c l e (66) suggests otherwise. 1.3. Scope of current i n v e s t i g a t i o n a-Brass was selected f o r the study material because (1) of i t s h i s t o r i c a l i n t e r e s t (2) i t s deformation c h a r a c t e r i s t i c s are well known (3) the embrittlement of the p o l y c r y s t a l l i n e form by ammoniacal environ-ments and by l i q u i d metals have been the subject of many experimental investigations and (4) almost every mechanism of stress corrosion cracking has, at some time, been applied to the brass/ammonia system. At the onset of the current work, i t was hoped to systematically, investigate the t e n s i l e behaviour of a-brass and g-brass monocrystals and p o l y c r y s t a l s deformed i n ammonia and mercury i n order to extend the analogies between stress corrosion cracking and l i q u i d metal embrittlement i n the d i r e c t i o n of continuous t e n s i l e loading conditions with a view to illuminating the mechanisms of i n i t i a t i o n and propagation of environmental stress cracks. '• Although the influence of s t r a i n rate and grain s i z e of a-brass deformed i n these environments was examined, the occurrence of a large invariant grain s i z e i n g-brass confined the study of t h i s phase to the e f f e c t of s t r a i n rate and thus extinguished the a t t r a c t i o n of - 17 -including, i n t h i s work, the a d d i t i o n a l influences of a d i f f e r e n t c r y s t a l structure and zinc composition. Since few te s t s of si n g l e c r y s t a l s i n stress corrosion or l i q u i d metal environments have ever been performed, a programme was designed to investigate the influence of Mattsson's s o l u t i o n and mercury on the y i e l d point and work hardening rate of a- and 3-brass monocrystals and to e s t a b l i s h (1) i f grain boundaries are e s s e n t i a l to the environmental a s s i s t e d stress cracking of brasses, (2) the influence of surface films on dislocation egression. Fatigue t e s t s (appendix II) were performed to throw a d d i t i o n a l l i g h t on the conjoint r o l e of crack p r o f i l e and environment. 13 -- 19 -T - Film Fresh Surface Revealed J F i g . 2 Film Rupture by s l i p step Emergence If 0 A f Tensile Stress F i g . 3 Tarnish Rupture Process 2. EXPERIMENTAL 2.1. Materials a-Brass was received as hard drawn bar of 3/8" diameter. Results of analyses performed by Coast-Eldridge are summarised below:-Element Ave Composition(%) Copper 70.47 Zinc 29.38 Iron 0.1 Lead 0.01 S i l i c o n 0.01 S i l v e r 0.005 Tin 0.01 Titanium 0.01 Gold, Vanadium Trace Others <0.001 • The binary Cu-Zn phase diagram i s i l l u s t r a t e d i n fi g u r e 4. The ammoniacal s o l u t i o n provided for the majority of tests were of the type used by Mattsson (26) i . e . containing 0.04 M CuSO^ SE^O and 1.5 M (NH ) SO. adjusted to the desired pH with approximately - 21 -2 N NH^OH. Separate stock solutions of 6 M (NH4)2SC>4 and 0.5 M CuS0 4.5H 20 were made up using b o i l e d double d i s t i l l e d water, and the actual d i l u t i o n and mixing of these solutions were c a r r i e d out before each series of t e s t s . pH measurements were made to an accuracy of i 0.005 pH units with a Corning model 10 pH meter, and three p r i n c i p l e pH values were investigated i . e . 6.8, 8.0 and 5.5. Care was taken to min-imise the influence of CO^ on cracking by (1) making up the solutions with b o i l e d double d i s t i l l e d water (2) by purging the tes t solutions with oxygen immediately p r i o r to use and (3) by storing solutions i n fla s k s with e i t h e r ground glass stoppers or potassium hydroxide U-tubes. A saturated s o l u t i o n of mercurous n i t r a t e was preferred to m e t a l l i c mercury f o r the work since i t s use f a c i l i t a t e d wetting of the brass by (1) d i s s o l v i n g the surface oxide and (2) depositing m e t a l l i c mercury on the brass by chemical displacement providing a "true*' i n t e r f a c e between s o l i d brass and l i q u i d mercury. ; 2.2. Tensile specimen preparation 2.2.1. P o l y c r y s t a l l i n e a-brass C y l i n d r i c a l unnotched t e n s i l e specimens were machined from the bar stock to the dimensions shown i n figu r e 5A,annealed (sheathed with brass f o i l to minimise d e z i n c i f i c a t i o n e f f e c t s ) , and cooled, descaled i n cold 40% n i t r i c acid, and the gauge necks polished to a 0000 emery f i n i s h . A f t e r extensive studies on short off-cuts of bar stock to determine the correspondence of heat treatment to grain s i z e , batches of speci-ments were annealed at temperatures ranging from 450°C to 700°C f o r - 22 -appropriate periods i n the range 1/2-6 hrs. The grain s i z e of each specimen was determined by an intercept method using a photomicrograph of an end o f f - c u t . Since annealing twins act as b a r r i e r s to most di s l o c a t i o n s (67) these were included i n the counts. Note.that specimens i n the i n i t i a l stages of the work were annealed f o r 1 hr. at 550°C :or fo r 3 hrs. at 650°C and grain s i z e control was effected only by repro-d u c t i b i l i t y of furnace conditions and are probably subject to minor grain s i z e v a r i a t i o n . , 2.2.2. P o l y c r y s t a l l i n e 3-brass 1/2" Diameter rods of 3-brass containing 49-50% zinc were produced i n the laboratory: appropriate quantities of a-brass and 99.999% zinc were melted;in quartz capsules f o r 3 hrs. at 980°C andthe product quenched v e r t i c a l l y i n t o water producing bubble free metal rod surfaces. Small off-cuts were annealed f o r 1 1/2 hrs at 550°C and hot r o l l e d to , a thickness of 0.025" with several short intermediate anneals at the i n i t i a l working temperature of 650°C. Specimens of dimensions i l l u s t r a t e d i n fig u r e 5B were blanked from cold s t r i p , deburred, e l e c t r o -polished i n chromic-acetic acid and stored i n a desiccator u n t i l required. The average grain diameter of the specimens was i n v a r i a b l y i n , the range 600-1000 microns: the thermomechanic treatments a v a i l a b l e could not confer grain s i z e control because of (1) rapid grain growth above 450°C and (2) i n a b i l i t y to work the ordered b r i t t l e material below 450°C. 2.2.3. Single c r y s t a l s a-brass Bar stock was cold swaged to 0.160" diameter and 6" lengths were encapsulated i n 7 mm i.d. vycor tubing forming shallow points - 23 -at the seeding end of the capsules: sin g l e c r y s t a l s were grown u t i l i s i n g a Bridgman technique with a capsule throughput speed of 2" per hour and furnace hot zone temperature of 980°C. Back r e f l e c t i o n X-ray photographs were obtained from each end of a p a r t i c u l a r c r y s t a l to determine f i r s t l y whether or not the c r y s t a l was singular and secondly the o r i e n t a t i o n of the t e n s i l e a x i s . Short lengths were removed from each end of the monocrystal and the cen t r a l portion was bisected to provide a p a i r of rods of c i r c u l a r cross section and approximately 2" long: gauge necks were c a r e f u l l y machined using the jewellers lathe and electropolished to the dimensions in, figure 5D. 2.2.4. Single c r y s t a l s g-brass Monocrystals.were prepared from 1/2" diameter cast rods (end-sectioned and annealed for 1 1/2 hours at 550°C) i n a manner s i m i l a r to that f o r a-brass save that the cast rods were hot swaged to 0.160" diameter with several intermediate anneals, at the working temperature of 650°C. 2.3. Tensile t e s t i n g procedure and conditions 2.3.1. Procedure A l l t e n s i l e tests were ca r r i e d out using the f l o o r model instro n machine at ambient temperatures except where stated. P o l y c r y s t a l l i n e specimens were f i t t e d as depicted i n fig u r e 5E into a teflon-pyrex c e l l enclosing the gauge neck or notch, and the environment introduced p r i o r to load a p p l i c a t i o n . In the case of sing l e c r y s t a l s , f l e x i b l e , t e f l o n tubing (figure 5F) replaced the teflon-pyrex c e l l and the ; - 24 -environment was introduced into the top of the c e l l by i n j e c t i o n through a hypodermic needle. Two specimens were tested under each set of conditions except i n the case of the sing l e c r y s t a l s . In the i n i t i a l stages of the work pH measurements were made before and a f t e r each t e s t but t h i s procedure was discontinued a f t e r the values were found to change by less than 0.1 pH unit during the course of the longest t e s t . Gauge and notch dimensions were determined before and a f t e r t e s t i n g by means of a t r a v e l l i n g microscope. Total elongations and elongation at frac t u r e i n i t i a t i o n were calculated from the associated autographic load time p l o t s . For 3-brass, elongation was preferred to reduction of area f o r assessing the e f f e c t of environment since deter-mination of cross s e c t i o n a l area was d i f f i c u l t due to the small (3%) amount of p l a s t i c deformation involved p r i o r to fracture i n some instances. A f t e r fracture each specimen was rinsed i n water and stored f o r fractographic or metallographic examination. Load time p l o t s corresponding to each specimen were converted., to load elongation p l o t s and compared, and i n the case of the grain s i z e e f f e c t s , p l o t s of fracture stress versus inverse square root of grain diameter were constructed. In the only high temperature tests which were performed, a heating tape and r e f l u x arrangement maintained the environment at i t s b o i l i n g point and 15 minutes were allowed to elapse f o r conditions to s t a b i l i s e p r i o r to a p p l i c a t i o n of the y i e l d load. A small volume of b o i l i n g concentrated mercurous n i t r a t e was added and cracking followed auto- . graphically i n load r e l a x a t i o n conditions. . - 25 2.3.2. Conditions 2.3.2.1. P o l y c r y s t a l l i n e material Experimental parameters were varied according to the following table. Material Grain Size Environment Size (microns) Crosshead Speed ("/min) a-brass a-brass(notched) 18-420 Approx.20 18-420 Approx.20 18-420 18-450 18-450 18-450 18-420 Mattsson's pH 6.8 Mattsson's pH 5.5 Mattsson's pH 8.0 Boiled double d i s t i l l e d water Ambient a i r 0.0002 to 0.1 0.0002 to 0.1 0.001 0.0002 to 0.1 0.001 0.0002 to 0.1 0.0002 to 0.1 Mercurous n i t r a t e 0.0002 to 1.0 Ambient a i r 0.001 -brass(unnotched) "Invariant" 600-1000 Mattsson's pH 6.8 0.001 to 0.1 5.5 8.0 Boiled double d i s t i l l e d water Ambient a i r Dry a i r S i l i c o n e o i l Mercurous n i t r a t e 0.02 to 10.0 The following a d d i t i o n a l experiments were performed:-- 26 -(1) Penetration E f f e c t s 96 micron grain diameter unnotched specimens were tested at s t r a i n rates of 0.1, 0.01, 0.001 per min i n Mattsson's s o l u t i o n (pH 6.8) and at 0.1, 0.001, and 0.0002 per min i n mercurous n i t r a t e s o l u t i o n a f t e r immersion i n the same for 7 days at room temperature. In addition, one specimen was coated with mercury, immersed i n a constant temperature bath for 6 days at 100°C, tested at room temperature at a s t r a i n rate of 0.001 per min i n mercurous n i t r a t e s o l u t i o n and compared with the control specimen without the mercury coating. Load r e l a x a t i o n notched tests were performed at 100°C i n mercurous n i t r a t e s o l u t i o n and compared to r e s u l t s of ambient temperature t e s t s . (2) Incremental loading tests One unnotched 96 micron specimen was loaded at a rate of 0.001"/min in mercurous n i t r a t e s o l u t i o n u n t i l cracking occurred. A s i m i l a r specimen was loaded i n a i r to t h i s p a r t i c u l a r stress when mercurous n i t r a t e was introduced. Any increase i n elongation on further loading was recorded. 2.3.2.2. Single Crystals a- and B-brass sin g l e c r y s t a l s were tested at a crosshead speed of 0.001"/min i n Mattsson's s o l u t i o n at pH 6.8 and i n mercurous n i t r a t e . 2.4. Fractography Topographical features of the fractured surfaces of t e n s i l e specimens were examined using the scanning attachment to the electron probe microanalyser and f i n e r d e t a i l s of such surfaces were observed using both d i r e c t and i n d i r e c t r e p l i c a electron microscopy. In addition, - 27 -fractured surfaces of sing l e c r y s t a l s were investigated employing low magnification photography. 2.5. R e f l e c t i o n electron microscopy (100 KV) Fractured surfaces of notched, large grained, a-brass specimens deformed i n Mattsson's s o l u t i o n and i n b o i l i n g water containing traces of s o l i d mercurous n i t r a t e . (The fractured surfaces of unnotched specimens contained some regions of d u c t i l e f r a c t u r e , and were of too small cross s e c t i o n a l area f o r the a p p l i c a t i o n of t h i s technique.) 2.6. Metallography 2.6.1. Crack path determination Fractured p o l y c r y s t a l l i n e a-brass t e n s i l e specimens were sectioned l o n g i t u d i n a l l y with a jewellers saw and c a r e f u l l y prepared f o r metallographic examination to determine the crack path. 2.6.2. Unstressed intergranular penetration determination Annealed 96 micron grain diameter a-brass rod off-cuts were . contacted with mercurous n i t r a t e s o l u t i o n and Mattsson's s o l u t i o n at room temperature f o r 1 week, and with mercurous n i t r a t e at 100°C f o r 6 days, and were sectioned f o r evidence of intergranular penetration.. - 28 -70/30 51/49 0 10 20 30 40 50 60 70 80 wt % Zinc F i g . 4 Copper-rich End of Copper-Zinc Equilibrium Diagram M B A -22. \ C y l i n d r i c a l Tensile 3 / 8 D I A 0.15DIA 16 thread/inch B Min,DIA 0.25" 3/ 8" DIA 0 0.200'»DIA 0.1Q0"DIA or 0.120"DIA LIQUID ENVIRONMENTS IPyrex C e l l Teflon Tape Teflon Bung S f r i p T e a s i l e Fatigue Specimens Single C r y s t a l Tensile "Teflon Tape ;jl=R Teflon Tubing n Wire Teflon Tubing Clamp F i g . 5 Dimensions of Specimens and C e l l Arrangements forfmechanical Testing. 3. RESULTS AND OBSERVATIONS 3.1. Single c r y s t a l t e n s i l e tests Total elongations for a-brass of 80 or 90% i n a i r were reduced to 30 or 40% i n eit h e r mercury or neutral Mattsson's s o l u t i o n , both environments having approximately equal e f f e c t s . Figures 6 and 7 reveal that both ammoniacal copper sulphate and mercury have no s i g n i f i c a n t e f f e c t on the y i e l d point and i n i t i a l work hardening rate of a-brass; the departure of the 'environmental' curve from that i n a i r i s att r i b u t e d to cracking. Values of C.R.S.S. are comparable with those disclosed by -4 Jamieson and S h e r r i l l (68) f o r Cu/28.5 Zn at a s t r a i n rate of 10 per sec, Ardley and C o t t r e l l (69) f o r Cu/28.8 Zn at 10 ^ per sec, ;. -4 P h i l l i p s and Robertson (70) f o r Cu/30 Zn at 10 per sec, and by Murakami and Ikai (71) for Cu/31.2 Zn at an undisclosed s t r a i n r a te. Serrated y i e l d i n g was evident i n one or two cases at very high s t r a i n s . Crystals of B-brass exhibited 60-70% t o t a l elongation i n a i r , 30 or 40% i n ammoniacal copper sulphate, and no p l a s t i c deformation at a l l i n mercury. Figure 8 shows that ammonia did not a f f e c t the y i e l d point or i n i t i a l work hardening behaviour, and figu r e 9 reveals the • d r a s t i c embrittling action of mercury. A sharp y i e l d point was not exhibited by B-brass and the C.R.S.S. was i n approximate accord with - 31 -that of Ardley and C o t t r e l l (69) for the same s t r a i n rate of 10 ^ per sec. In contrast to the a-brass sin g l e c r y s t a l s , the 3-brass dried very quickly a f t e r removal from the mercurous n i t r a t e s o l u t i o n , possibly due to inward d i f f u s i o n of mercury or i n t e r m e t a l l i c formation at the surface. The r e l a t i v e l y invariant y i e l d stresses and i n i t i a l work hardening rates of a-brass i n Mattsson's s o l u t i o n i n d i c a t e that the oxide has no e f f e c t on d i s l o c a t i o n egression from the c r y s t a l suggesting that the th i n t a r n i s h possesses very poor mechanical properties. This i s supported by Donaldson's (72) tests i n a i r of a-brass sin g l e c r y s t a l s , bearing a thick coat of dry t a r n i s h . 3.2. P o l y c r y s t a l l i n e t e n s i l e tests 3.2.1. a-Brass 3.2.1.1. S t r a i n rate e f f e c t s Load-extension curves from autographic load-time data f o r a-brass of defined grain s i z e deformed at a series of s t r a i n rates i n Mattsson's so l u t i o n of pH 6.8, 5.5 and 8.0, and i n mercury are shown i n figures 10-13. Contingent upon s l i g h t grain s i z e v a r i a t i o n due to problems with furnace temperature r e p r o d u c i b i l i t y , i t i s noted that neither of these environments nor s t r a i n rate influences the y i e l d point or work hardening r a t e , but only l i m i t the elongation at which fracture occurs, i n accord with previous observations (8,9) relevent to other systems. In the range examined, s t r a i n rate did not influence properties i n d i s t i l l e d water or ambient a i r . - 32 -E f f e c t of environment and s t r a i n rate on d u c t i l i t y i s reported i n figures 14-16. The f a m i l i a r U shaped (28,29,33) dependence on pH i s noted from fig u r e 14, and f i g u r e 15 represents the same r e s u l t s i n an a l t e r n a t i v e manner to include the action of a i r and water. The influence of s t r a i n rate i n mercury (figure 16) on elongation at crack i n i t i a t i o n i s compared with s i m i l a r properties i n ammonia i n f i g u r e 17, and i s supported by figure 18 which compares the e f f e c t of several environments on the load-extension curve. While mercury and ammonia have s i m i l a r s t r a i n rate e f f e c t s on the y i e l d point, work hardening rate and d u c t i l i t y , the former environ-ment appears the more e f f e c t i v e embrittling agent. 3.2.1.2. Grain s i z e e f f e c t s Load-extension curves f o r a-brass as a function of grain s i z e deformed at constant s t r a i n rate i n a i r and ammoniacal copper sulphate s o l u t i o n compose.figures 19-21. Similar curves f o r mercury have been presented i n f i g u r e 13. Decreasing grain s i z e at constant s t r a i n rate diminishes uniform elongation and elongation at fracture i n a i r , and whilst increasing the elongation at crack i n i t i a t i o n i n mercury . and a l k a l i n e Mattsson's s o l u t i o n , has l i t t l e e f f e c t i n acid (pH 5.9) : or neutral Mattsson's s o l u t i o n . This behaviour i s also noted i n Mattsson's pH 6.8 s o l u t i o n and i n mercury for other s t r a i n rates (figures 10, 12, and 13). No grain s i z e e f f e c t s on elongation under continuous t e n s i l e loading i n ammoniacal copper sulphate have hi t h e r t o been reported but i n mercury Rostoker et a l (9) have observed s i m i l a r e f f e c t s whilst Rosenberg and Cadoff (8) found that the elongation of - 3 3 -s t r i p specimens increased with increasing grain s i z e . It i s noted from fi g u r e 2 2 - 2 4 that y i e l d stress and fracture s t r e s s , i n a l l environments examined, varies l i n e a r l y with inverse square root of grain diameter, as previously observed ( 2 , 3 , 8 , 9 , 7 3 ) i n other metal/environment couples; the fracture stress r e l a t i o n s h i p has more scatte r than the y i e l d stress r e l a t i o n s h i p . The surface energy (y) associated with fr a c t u r e , obtained from the slopes of these p l o t s (table 1 ) are comparable with values quoted i n section 1 . 2 . 3 . 1 . Since the values corresponding to ammmoniacal and mercurous environments are lower than those experimentally determined surface energy of a-brass with respect to i t s own vapour i . e . 1 5 0 0 ergs/sq cm ( 2 ) , i t appears that these environments reduce the surface energy associated with j formation of d i s l o c a t i o n nucleated cracks. However several cursory observations are indicated: ( 1 ) The slope corresponding to environmental fracture i s almost equal to that for d u c t i l e f a i l u r e , i n which the low value of y i s meaningless. ( 2 ) Y i n ammoniacal copper sulphate i s less than that i n mercury when the reverse would be consistent with a surface energy mechanism since mercury i s more embri t t l i n g . ( 3 ) y i - n mercury varies systematically with loading rate. An invariant condition would be more suited to a surface energy model. 3 . 2 . 1 . 3 . Incremental loading tests Figure 2 5 reveals that 2 3 / 4 % and 1 1 / 2 % further extension occurred a f t e r introduction of mercurous n i t r a t e at stresses X and Y r e s p e c t i v e l y . An i n i t i a t i o n period should be incompatible with a surface energy mechanism unless time i s required for the formation of a - 34 -c r i t i c a l concentration of species as may be the case of ammonia (28, 33) . Since the mercurous n i t r a t e s o l u t i o n i t s e l f was saturated, immediately wets the surface and p r e c i p i t a t e s mercury on the brass surface, i t i s u n l i k e l y that a surface energy mechanism i s operative:" an i n i t i a t i o n time i s more consistent with d i s s o l u t i o n or i n t e r m e t a l l i c / t a rnish rupture type mechanisms. 3.2.1.4. Penetration tests Immersion test r e s u l t s are presented i n Table I I . P r i o r immersion i n Mattsson's s o l u t i o n or mercury at room temperature did not influence the mechanical properties but the s i n g l e elevated temperature mercury coated immersed specimen exhibited reduced d u c t i l i t y and increased y i e l d point possibly pointing to the rapid inward d i f f u s i o n of mercury at high temperatures with d i s s o l u t i o n occurring followed b y the p r e c i p i t a -t i o n of a b r i t t l e product at a lower temperature. A marked reduction i n i n i t i a t i o n time was observed i n the high temperature load r e l a x a t i o n tests (figure 26) i n mercury. The i n i t i a l stages of the autographic load time plots obtained when preparing both mercury and ammonia fractured specimens for electron d i f f r a c t i o n are s i m i l a r i n shape and comparable with those of Tromans, Dowds and Leja (29) for brass i n ammoniacal copper sulphate. Such a decrease of i n i t i a t i o n time might be consistent with increasing metal d i s s o l u t i o n rates or penetration rates i n the order ammonia at room temperature, mercury at room temperature, and mercury at 100°C. In a l l cases including continuous t e n s i l e loading and load-relaxation - 35 -conditions whenever mercurous n i t r a t e contacted a-brass f o r periods greater than a few minutes, the environment adopted a blue c o l o r a t i o n . 3.2.2. g-Brass 3.2.2.1. S t r a i n rate e f f e c t s The influence of environment and s t r a i n rate on the load extension curves of p o l y c r y s t a l l i n e g-brass are shown i n figures 27-29, and the e f f e c t on t o t a l elongation i n figures 30 and 31. The behaviour i n Mattsson's s o l u t i o n resembles that of a-brass save that g-brass appears more b r i t t l e , and even cracked prematurely i n ambient a i r , i n agreement with Bailey (74) . Note that d i r e c t comparison with a-brass i s not too meaningful because of (1) problem of grain s i z e control , (2) d i f f e r e n t specimen geometry and dimensions (3) large grain s i z e with respect to specimen gauge area . (4) s t r i p specimens tend to tear providing s l i g h t l y over-exaggerated elongation values. Mercury severely embrittled g-brass (figure 29) and even at very high loading rates fracture occurred below the y i e l d s t r e s s . 3.2.2.2. Grain s i z e e f f e c t s The invariant microstructure of g-brass prevented any i n v e s t i g a t i o n of grain s i z e e f f e c t s . 3.3. Optical metallography 3.3.1. Crack path Intergranular cracking was noted i n a-brass specimens fractured i n - 36 -ammonia ( a l l pH values) and mercury, and the shape of the crack was independent of environment. The representative photograph (figure 32) indicate features commonly observed. 3.3.2. Penetration No intergranular penetration of black tarnish or of mercury into annealed a-brass was observed i n the work. The presence of mercury which flowed over the prepared polished surface, i n any case, masked any penetration along the grain boundaries. McEvily and Bond (53) have noted such penetration of black t a r n i s h but no studies of room temperature mercury grain boundary penetration of brass appear to have been performed. 3.4. R e f l e c t i o n electron microscopy Electron d i f f r a c t i o n patterns were obtained from the gold standard (figure 33) and the fractured surfaces of p o l y c r y s t a l l i n e a-brass deformed i n b o i l i n g d i l u t e mercurous n i t r a t e (figure 34 and 35) and i n ammoniacal copper sulphate (figures 36 and 37). Patterns of the black tarnish compose figures 38 and 39 and analyses of a l l r i n g patterns in, tables I I I , IV and V. Rings from the high temperature mercury fractured specimen (table III) correspond to a t h i n zinc layer spread over a large area, whilst the d i f f r a c t e d spots indicated a l o c a l concentration of product o of l a t t i c e parameter 9.6 A, and those given by two patterns not shown 0 0 o yielded l a t t i c e parameters 10.9 A and 11.4 A. That of CuHg i s 9.415 A (75). - 37 -Rings (table IV) from specimens fractured i n ammoniacal copper sulphate s o l u t i o n i n v a r i a b l y corresponded exactly with those of Cu^O (76) the major contituent of the t a r n i s h (table V) even though the fractured surface was not black, but a darkened yellow. One pattern (figure 37) c l e a r l y shows a broadened t h i r d band i n d i c a t i n g the coincidence of two r i n g : p a t t e r n s e.g. C^O with a higher oxide or brass Since the other rings (77,78) comprising the patterns of these materials were masked by Cu^O l i n e s , there i s a p o s s i b i l i t y of existance of. these compounds i n addition to C^O on the fractured surface, even though they could not be p o s i t i v e l y i d e n t i f i e d . Moreover i t was demonstrated that a p r i o r t h i n layer of anodic or cathodic product could prevent the formation of the black t a r n i s h i n Mattsson's s o l u t i o n and reproduce the observed darkened yellow fractured surface (N.b. i n no way i s i t suggested that the electrochemical conditions p r e v a i l e n t during the stress corrosion cracking of a-brass were simulated.) 3.5. Fractography 3.5.1. P o l y c r y s t a l l i n e a-brass - room temperature a-Brass fractured i n a i r exhibited cup and cone f a i l u r e and electron fractbgraphs i n figures 40 and 41 reveal the f a m i l i a r character-i s t i c s of such d u c t i l e f a i l u r e . Intergranular f a i l u r e i n ammonia (figure 42) resembles that i n mercury (figure 43) and f i g u r e 44 represents the t r a n s i t i o n region from d u c t i l e mechanical f a i l u r e to ammonia induced intergranular f r a c t u r e . Attempts to resolve f i n e r d e t a i l s of mercury fractured brass by d i r e c t r e p l i c a studies were f o i l e d by the unbiquitous presence of the l i q u i d metal on the surface. Room temperature fracture i n mercury, (figures 45 and 46) compares with that i n ammonia (figures 47-54); an - 38 -intergranular path and s l i p l i n e s feature i n both environments and i n ammonia, there i s a f i n e l y d e t a i l e d surface and corrosion product, of which several aspects are shown. One sole area of possible trans-granular fracture was observed (figure 52). 3.5.2. P o l y c r y s t a l l i n e a-brass - high temperature mercury Intergranular f a i l u r e (figure 55), s l i p s t r i a t i o n s (figure 56), and i n addition p e c u l i a r pyramidal type features proud of the surface (figures 57-62) were noted. The observation of such features on fr e s h l y fractured surfaces by means of the probe (figure 63) indicates that the e f f e c t was not associated with the r e p l i c a techniques. Attempts to q u a l i t a t i v e l y determine the composition of the feature yielded ambiguous r e s u l t s . A zinc scan (figure 64) of a selected area (figure 65) containing a mercury globule and several of these " a r t i f a c t s " • revealed an apparent loss i n zinc, and a mercury scan (Fig. 66) indicated no d i f f e r e n c e from the background but markedly less concentration than the mercury globule. However the apparent reduced zinc content could be a t t r i b u t e d to d e f l e c t i o n of the beam by the features. 3.5.3. P o l y c r y s t a l l i n e 3-brass - room temperature 3-brass f a i l e d by an intergranular path i n mercury (Fig. 67) and ammonia (Figs. 68 and 69). Transgranular features are also evident i n these l a t t e r fractographs. Presence of mercury on the surface (Fig. 67) prevented d e t a i l e d examination. 3.5.4. Single c r y s t a l s Transgranular cracking of a-brass i n a i r , mercury and neutral' Mattsson's s o l u t i o n , and of 3-brass i n a i r and mercury are shown i i i - 39 -Figures 70 and 71. Low magnification photography of the mercury embrittled 3-brass specimen (Figure 72) reveals cleavage steps on the fra c t u r e d surface: f i n e r d e t a i l s are i l l u s t r a t e d by the probe f r a c t o -graphs - Figs. 73 and 74. and the macroscopic .crack plane was close to {110} 3.6. Analogies of t e n s i l e behaviour i n mercury and Mattsson's s o l u t i o n - a summary Li q u i d metal or str e s s corrosion environments do not influence the y i e l d point or i n i t i a l work hardening rate of p o l y c r y s t a l l i n e , or s i n g l e c r y s t a l s of, a- or 0-brass but l i m i t the elongation at f r a c t u r e . In the p o l y c r y s t a l l i n e m a t e r i a l , i n c r e a s i n g the crosshead speed serves only to extend t h i s l i m i t , not acting on the y i e l d point or work hardening r a t e . Decreasing the grain s i z e ( of a-brass) increases elongation at f r a c t u r e i n mercury and Mattsson's s o l u t i o n (pH 8.0). P r i o r immersion of a-brass i n mercury or Mattsson's s o l u t i o n has no e f f e c t on the mechanical pro p e r t i e s determined i n these media: -1/2 a vs D l i n e a r c o r r e l a t i o n i s exhibited i n both environments, and f r a c t u r e was i n t e r g r a n u l a r . Load r e l a x a t i o n behaviour i n mercury at room temperature and 100°C was s i m i l a r to that i n ammonia, and t h i n f i l m s corresponding to cor r o s i o n products were found on fr a c t u r e d surfaces of specimens deformed i n Mattsson's s o l u t i o n and i n b o i l i n g water containing traces of s o l i d mercurous n i t r a t e . F i g . 10. Load-Extension curves of a-brass deformed i n Mattsson's Solution (pH6.8) at a Series of Crosshead Speeds. 1900 i800 200 600 Load lbs. LO 500 500 300 200 200 0.0002 10 pH 5.5 Crosshead Speed /mm 20 30, % E l . 40 50 . pH 8.0 Crosshead Speed"/mir 1 0.05 0.01 > 0.001 °' 1 0.0002 • • 10 20 30 40 50 % E l . F i g . 11 Load-Extension curves of a-brass deformed i n Mattsson's Solution at a Series of Crosshead Speeds (Grain diamter<s20vi) 800 700 600 500 Load l b s . 400 300 200 100 V 0.0002 _ J 10 0.05 pH 6.8 Crosshead Speeds "/min \ » » \ 0.1 _j 20 30 60 70 F i g . 12. 40 50 % P l a s t i c Elongation Load-Extension curves of large grained a-brass deformed i n Mattsson's Solution (pH 6.8) at a Series of Crosshead Speeds. Load lbs 800 700 600 500 400 300 200 100 Strain Rates A 1.0 per min B 0.1 per min C 0.001 per min D 0.0002 per min 0 6 7 8 9 10 % P l a s t i c Elongation 11 12 13 14 F i g . 13 Load-Extension curves of a-brass deformed i n Mercury as Function of Grain Size and S t r a i n Rate. - 48 -5.5 pH 8 F i g . 15 D u c t i l i t y of a-braas vs. pH of Mattsson's Solution f o r a Series Crosshead Speeds (Grain Dia approx. 20y.) A Boiled double d i s t i l l e d water B pH 8.0 Mattsson's Solution -C pH 5.5 Mattsson's Solution D pH 6.8 Mattsson's Solution Log S t r a i n Rate(per min) F i g . 15 Influence of environment on the relationships between d u c t i l i t y of a-brass and S t r a i n Rate (grain diameter = 20u) - 50 -F i g . 16 Elongation vs S t r a i n Rate as Function of Grain Diameter for a-brass i n Mercury. log S t r a i n Rate (per min.) - 51 --4 -3 -2 -1 Log S t r a i n Rate (per min.) Fi g . 17 Elongation at Crack I n i t i a t i o n vs S t r a i n Rate for a-brass deformed i n Mattsson's Solution and Mercury. Crosshead Speed 0.01"/ min Crosshead Speed O.OOtJVmin Fi g . 18 Load-Extension Curves of a-brass deformed i n several Environments for two Crosshead Speeds (Grain Diameter 20y) Load lbs 1000 900 800 700 600 500 400 300 200 100 CrosshMad Speed 0.001"/min. F2230y G 300y 0 10 15 20 25 30 35 40 45 50 % P l a s t i c Extension F i g . 19 E f f e c t of Grain Size on Load Extension Curve of a-brass i n ambient A i r . - 5 4 -F i g . 20 E f f e c t of Grain Size and Environment on the early stages of the Load-Extension Curve of a-brass (deformed at a cross-head speed of 0.001"/min.) F i g . 21 E f f e c t of Grain Size and Environment on the Load Extension Curve of a-brass deformed at Crosshead Speed of 0.001"/min. - 57 -Stress k . s . i . 90 80 70 60 50 40 30 20 10 Crosshead Speed 0.001"/min. A Mattsson's pH8800 V pH 5.9 © pH 6.8 X . Y i e l d Stress Indicates two Readings 8 10 12 14 16 18 20 22 24 26 28 - 1 - -l* (Grain Diameter) 2 cm 2 F i g . 23 Fracture Stress vs. D 2 of a-brass embrittled by Mattsson's Solution. - 58 -Table I E f f e c t i v e surface energies (Rostoker's analysis) Environment y (ergs/sqcm) Mattsson's s o l u t i o n (pH 6.8) 235 (pH 5.9) 301 (pH 8.0) 625 Mercurous n i t r a t e - 1.0"/min 310 crosshead speed rt . / n r i r 0.1 /min 400 0.001"/rain 464 0.0002"/min 412 A i r - unnotched A i r - notched 262 315 Hi- 5 9 -@ S i x a i n RAte of 1.0 per min. O " " o.i;" :: A ' . *' Y i e l d Stress 0.001 0.0003 4 6 8 10 -h -h (Srlain Diameter) cm 12 14 16 18 20 22 24 F i g . 24 Fracture Stress vs D ^ for a-bcass" embrittled Hg. 0 1 2 3 . 0 1 2 3 4 5 0 1 2 3 4 5 6 7 . % P l a s t i c Elongation F i g . 25 "Incremental Loading" Test'performed on 3 p o l y c r y s t a l l i n e a-brass Specimens .. A Load i n Hg to Fracture at X - ' . B-Loaded i n A i r to X, add Hg ,.. Doad-to Fracture at Y C Loaded i n A i r to Y, add Hg , Load to Fracture at Z - 61 -Table II Penetration/ Immersion Tests Specimen Treatment Test Conditions Y i e l d Point %Uniform Temp.°C Environ-ment Crosshead speedC"/min) lbs Elongation Immerse 7 days 25°C Mattsson pH 6.8 25 Mattsson .1 240 45 M I I J M I I .01 245 21 I I II I I M .001 240 . 5.5 Immerse 7 days 25°C HgNO^solution 25 HgN03 .1 240 - 5' II I I it .001 240 3.5 , . .... ti • ii ii .0002 245 2 Immerse 6 days b o i l i n g water Hg coating 25 HgN03 .001 350 0_ I I I I ii no coating ; — •• 25 HgN03 .001 245 3.2 Load lbs 1300 4 00 300 100 Mattsson's pH 6.8 Mercury at Room Temperature 10 15 Time (minutes) F i g . 26 Load-Relaxation Curves of a-brass i n ftattsson's Solution (pH 6.8) and Mercury 20 Fig.27 Load-Extension Curves of B-brass deformed i n bo i l e d double d i s t i l l e S d w a t e r at a seri e s of Crosshead Speeds ("/min.) • « _ , — , — , , . , ' • r ON o :• \ 20 so •/;> so " so o 10 20 30 40 Percentage P l a s t i c Extension Percentage P l a s t i c Extension F i g . 28 Load-Extension Curves of g-brass deformed i n several Environments at two Rates. Elongation % F i g . 29 Influence of Mercury on Load-Elongation Curve of p o l y c r y s t a l l i n e B-brass f o r indicated Crosshead Speeds. - 66 -6-Brass I Crosshead Speed 6.1 "/min . 2 " " 0.01 "/min 3 " " 0.005 "/min 4 " " 0.001 "/min 30 r 25 h Percentage j Elongation I F i g . 30 Relationship between d u c t i l i t y and pH of Mattsson Percentage Elongation 40 30 20 10 A pH 8.0 Mattsson's Solution. B pH 6.8 C pH 5.5 D Boiled double d i s t i l l e d water. Log Crosshead Speed("/min) : Dry A i r S i l i c o n e O i l Ambient A i r D A B C - 3 ON F i g . 3 1 Influence of environment on the r e l a t i o n s h i p between d u c t i l i t y of Beta Brass and crosshead speed. - 68 -F i g . 32 Intergranular cracking of a-brass (xl20) Etch FeCl„ (Representative of cracking produced by mercury and ammonia). - 69 -F i g . 34 F i g . 35 F i g . 33 Au standard F i g . 34 Mercury fractured a-brass (ring pattern) F i g . 35 Mercury fractured a-brass (spot pattern) - 70 -r i g . J D F i g . 37 F i g . 36 Ring pattern a-brass-Mattsson's s o l u t i o n F i g . 37 Ring pattern a-brass-Mattsson's s o l u t i o n (showing broad 3rd ring) - 71 -F i g . 38 Black tarnish F i g . 39 Black tarnish - r i n g pattern ri n g pattern - 72 -Table III Analysis of Rings from fractured Surface of a-brass deformed i n Mercury at 100°C Camera Constant = 1.33 inch A* = 2R.d 2R = diameter of r i n g d = interplanar spacing Ring No 2R(inch) d ( X ) ^ I n t e n s i t y Zinc l i n e s 1 0.54 .'. 2.41 med./weak 002 d-= 2.47 A5 2. : 0.58' 2.29 med./weak 100 2.30 3 0.70- 1.9 strong 1 01 2.09 4 ; 0 .74 1 .79 med. 102 1.68 5' 1.06 1.26 med. 110 1.33 6 . 1.25 - 1.06 med. 202 1.05 7 1.62 0.82 med. 2 12 0.82 - 73 -Table IV Analysis of Patterns from fractured Surface of a-Brass deformed i n  Mattsson's s o l u t i o n Camera Constant = 1^33 inch ft 2R.d. 2R = diameter of r i n g r d = interplanar spacing Ring No 2R(inch) Intensity d(ft) Cu 20 l i n e 1 0.«4 weak 3.02 110 d = 3.02ft 2 0.S5 strong 2.41 Mo 2.47 3 0.63 medium . 2.11 200 2.13 4 0.74 weak ' 1.79 2 11 1.74 5 0.88 weak 1.51 220 1.51 6 1.01 weak 1.31 311 1.29 7 1.08' . weak 1.23 2 22 1.23 8 1.L4 weak 1.16 400 •1.07 9 1 .36 weak 0.97 331 ~ 0 .98 10 1.40 med./weak 0.95 420 0 .96 1 1 1.52 weak 0.87 422 0.87 12 1.60 "• weak 0.83 5 11 0.82 1 0.44 weak 3.02 110 3.02 0.54 ~ / If 2 strong 2.46 2.47 2.16-3 band 0 .615 medium 2.06 2 00 2.13 4 0.645 weak 1.75 211 1.74-5 0.76 weak 1.51 220 1.51 6 1.00 . weak 1.31 311 1.29 7 i.o6; weak 1.09 400 1.07 8 1.14 weak 0.99 331 0.98 9 1.34 weak 0.87 422 0.87 - 74 -Table V Analysis of Tarnish Rings Camera constant = 1.'33 inch A* = 2 R d 2R = diameter of r i n g d = interplanar spacing : Ring No 2R(inch) Intensity d(A) C U 2 O l i n e s 1 0.43 medium 3.09 110 d = 3.02 A 2 0.52 v. strong 2.54 111 2.47 3 0.602 m-strong 2.20 200 2.13 4 0.74 weak 1.78 211 1.77 5 0.86 medium-strong 1.54 220 1.51 6 1.0 medium-strong 1.33 311 1.29 7 1.04 v.weak 1.27 222 1.23 1 0.43 3.09 2 0.53 v. strong 2.5 3 0.61 2.18 4 0.745 1.78 as above 5 0.855 strong 1.56 6 1.005 1.32 7 1.05 v.weak 1.26 - 75 -F i g . 40 Duc t i l e dimples i n a i r - f r a c t u r e d a-brass. (x600 probe). F i g . 41 Elongated fracture dimples i n a i r -fractured a-brass.(x600 probe). Fig. 43 Intergranular cracking of a-brass i n mercury. (x300). F i g . 44 T r a n s i t i o n region from d u c t i l e -intergranular fracture i n Mattsson's s o l u t i o n fractured a-brass.(x300) . - 78 -F i g . 45 Direct r e p l i c a of a-brass fractured i n Hg at room temperature (x400) F i g . 46 Direct r e p l i c a of a-brass fractured i n Hg at room temperature (xl6000). - 79 -Fi g . 49 F i g . 50 Fig . 47 a-brass-fractured Mattsson's s o l u t i o n (x 2000) Fig- 48 " " " " ( x 5 > 800) Fig- 49 " " " ( x 4,000) Fig- 50 " " " " ( x 10,200) - 80 -Dig. F i g . 54 Fig . 51 a-brass fractured i n Mattsson's s o l u t i o n (x 1,500) Fig- 52 " » ( x 2,000) F i g . 53 " " " ( x 5 ) 8 0 0 ) Fig- 54 " " " ( x 6 ) 5 0 0 ) F i g . 56 Intergranular f a i l u r e of a-brass i n Hg at 100°C (x4000). _ 82 -- 84 -x 600 x 1200 x 2500 F i g . 63 Probe topography of 100°C mercury fractured a-brass showing surface features at a seri e s of magnifications. - 85 -F i g . 66 Zinc scan. Selected area of 100°C HgNO^ fractured a-brass specimen for probe X-ray examination (x 2500). Mercury scan. F i g . 64 F i g . 65 F i g . 66 - 86 -67 Intergranular cracking of g-brass by mercury (x 255) . - 8 7 -- 88 -F i g . 70 Low magnification (x5) photograph of a-brass s i n g l e c r y s t a l s fractured i n : A a i r B Ammonia C mercury - 89 -F i g . 73 and 74 Topography of fractured surface of specimen 6(x225) 4. DISCUSSION 4.1. Discussion of aforementioned mechanisms of environmental stress a s s i s t e d cracking i n terms of present observations 4.1.1. Surface energy mechanism A surface energy mechanism of crack i n i t i a t i o n based on d i s l o c a t i o n nucleated cracks has been applied to the cracking of s i n g l e and poly-c r y s t a l s . There i s l i t t l e doubt that b a r r i e r s to d i s l o c a t i o n s , a p r e r e q u i s i t e of t h i s mechanism, are a v a i l a b l e i n the form of i n t e r -secting s l i p bands, grain boundaries and oxide films (79). In order f o r Rostoker's mechanism based on d i s l o c a t i o n p i l e ups at grain boundary-surface i n t e r s e c t i o n s to be operative, p r o v i s i o n must be made for intergranular cracking: t h i s is. achieved by assuming (1) a grain boundary w i l l "open up" under the action of a s u i t a b l y large t e n s i l e component of the d i s l o c a t i o n p i l e up stress or (2) an i n i t i a t e d crack w i l l enjoy a short transgranular passage p r i o r to s e l e c t i n g an intergranular path: the observations of such transgranular regions at the edge of 3-brass specimens (Fig. 68) and a-brass (80) support the l a t t e r t h e s i s . (No such d e t a i l was observed i n a-brass i n the .present work.) -1/2 It was shown i n s e c t i o n 1231 that Rostoker has c i t e d o f vs D - 91 -l i n e a r i t y as evidence for a surface energy mechanism since the analysis predicts y values of the correct order of magnitude f o r b r i t t l e f a i l u r e : the current work questions the a s s e r t i o n on two grounds. F i r s t l y , the slopes corresponding to fracture i n ammonia or mercury approximate to that i n a i r i . e . for d u c t i l e f a i l u r e , i n which the low value of y i s meaningless. It i s therefore f o r t u i t o u s that the analysis y i e l d s y of the correct order of magnitude f o r b r i t t l e f a i l u r e . Secondly, since CQ increases r a p i d l y with p l a s t i c deformation, the -1/2 equation Op = O q + KD i s not v a l i d f o r fracture at grain s i z e : variant s t r a i n s . The operation of a surface energy initiation mechanism i s consistent with instant fracture on introduction of the environment at an appropriate stress provided there e x i s t s a s u f f i c i e n t concentration of surface a c t i v e species. The i n i t i a t i o n period i n load r e l a x a t i o n t e s t s of a-brass i n Mattsson's s o l u t i o n may be associated with the. time f o r b u i l d i n g up of a c r i t i c a l concentration of complex ions (33) but such 'an explanation i s u n satisfactory i n terms of a concentrated mercurous n i t r a t e environment i n both load-relaxation or "incremental" loading conditions since wetting i s almost instantaneous. On a f i n e scale, ammonia fractured surfaces were uniform i n back-ground but whether smooth enough to be associated with surface energy b r i t t l e propagation i s debatable; moreover, the random appearance of > some compound on the ammonia fractured surfaces i s inconsistent with the continuous operation of t h i s model. The fractured surfaces of mercury specimens deformed at high temperatures were also of uniform appearance with a random dispersed compound, but since these features could have - 92 -formed at room temperature a f t e r a d i s s o l u t i o n r e a c t i o n at the elevated temperature, no c r i t i s m could be l e v e l l e d at a surface energy mechanism from t h i s viewpoint. Although many observations of stress corrosion cracking and l i q u i d metal embrittlement have been explained by an adsorption-bond rupture mechanism of crack propagation (section 1.2.3.), a model which i s of a q u a l i t a t i v e nature i s of l i m i t e d value only. 4.1.2. Tarnish rupture mechanism The t a r n i s h rupture mechanism has been applied to both intergranular and transgranular crack i n i t i a t i o n . As o r i g i n a l l y proposed, the process has explained the fracture of a sing l e c r y s t a l of a-brass r e q u i r i n g stress concentrations associated with p i l e ups at the oxide-metal i n t e r f a c e to rupture a thick hard b r i t t l e oxide f i l m . Since d i s -l o c a t i o n egression does not appear to be impeded by such a f i l m , the o r i g i n a l concept i s c l e a r l y i n v a l i d : the ta r n i s h rupture mechanism i s not, however, necessa r i l y d i s q u a l i f i e d but modifications based on a s l i p step displacement of a weak powdery f i l m are required. A p r i n c i p l e shortcoming of the t a r n i s h rupture mechanism i s that s p e c i a l p rovision must be made for intergranular f a i l u r e : a p p l i c a t i o n to intergranular cracking assumes p r e f e r e n t i a l cracking of the t a r n i s h i s occurrent at grain boundaries, and indeed fr a c t u r e at these s i t e s i s expected to be e a s i l y incurred by grain boundary rumpling or stress .concentrations associated with d i s l o c a t i o n p i l e ups at intersections of the boundaries with the oxide/metal i n t e r f a c e . A l t e r n a t i v e l y , the fract u r e of p r e f e r e n t i a l l y penetrated grain boundary t a r n i s h i s assumed - 93 -to t r i g g e r cracking, i n which case, the observed i n i t i a t i o n time i s considered equal to the time elapsed during the d i f f u s i o n c o n t r o l l e d (53) t a r n i s h entrenchment to the appropriate depth. It i s debatable whether such a small penetration would be e f f e c t i v e i n i n i t i a t i n g cracking. Moreover penetration of the black t a r n i s h into the grain boundaries (of annealed brass) i s not a common occurrence (section 3.3.2). The absence of a v i s i b l e black t a r n i s h on the fractured surfaces of a-brass i n continuous t e n s i l e loading conditions i s only compatible with the o r i g i n a l t a r n i s h rupture process i f the repeated stages of ta r n i s h formation and rupture occur so r a p i d l y that the t a r n i s h b u i l d up i s small, allowing only f o r the formation of the observed thinner Cu^O f i l m . However, the fac t that the f i l m does not increase i t s thickness i . e . darken V':'rapidly on prolonged (29) immersion i n neutral Mattsson's s o l u t i o n suggests that i t s physical c h a r a c t e r i s t i c s are. d i f f e r e n t from those of the v i s i b l e black t a r n i s h , so as to r e s t r i c t d i f f u s i o n of relevent species: therefore the t a r n i s h rupture mechanism would require at least a modification of the C^O producing mode at the fractured surfaces. As o r i g i n a l l y proposed, the ta r n i s h rupture mechanism i s not applicable to non-tarnishing solutions. However, since s i m i l a r behaviour ( i . e . deformation curves, d i f f r a c t i o n patterns and fractography) was exhibited i n both types of s o l u t i o n , i t i s probable that the mechanism can be applicable to non-tarnishing solutions with only a s l i g h t modifi-cation. The l i q u i d metal analogy has been discussed (section 1.2.2). In the - 94 -present work, d e t a i l s of r e a c t i o n products on room temperature mercury fractured surfaces, and of intergranular penetration were indeterminate and provided no c r i t i s m of the mechanism. The i n t e r m e t a l l i c rupture process does not apply at elevated temperatures and since there i s no sharp t r a n s i t i o n (51) i n cracking behaviour at the i n t e r m e t a l l i c compound formation temperature, i t i s doubtful whether the compound plays any s i g n i f -icant r o l e (other than that of a passive f i l m ) . See also Appendix I. 4.1.3. D i s s o l u t i o n D i f f r a c t i o n patterns of reaction product on the fractured surfaces o f f e r evidence of d i s s o l u t i o n . The occurrence of i n i t i a t i o n periods i n load r e l a x a t i o n and 'incremental loading' conditions, and the observed fractographic features are consistent with the operation of a d i s s o l u t i o n mechanism. Stress must be operative i n add i t i o n to d i s s o l u t i o n to account for the observed cracking rates (section 1.2.1). The chemistry of d i s s o l u t i o n of a - b r a s s i n these solutions i s ; both complex and co n t r o v e r s i a l and has been discussed elsewhere (26,28,30, 33,37,81). 4.2. Development of a cracking mechanism In accounting f o r the observations, the mechanism must include three e s s e n t i a l features:-(1) p l a s t i c deformation (2) d i s s o l u t i o n (3) s i m i l a r mechanism operating i n both i n i t i a t i o n and propagation i stages as suggested by the observation of Tromans et a l . (29) that both stages exhibit the same U shaped rate dependence on pH( F i8- 7 - 95 -It has been shown that p r i o r immersion of a-brass i n Mattsson's s o l u t i o n has no e f f e c t on the mechanical properties determined by-continuous loading t e s t s . The f i l m thickens with time (53,82) reaching a depth of 5u a f t e r 4 hrs immersion compared with less than lu on immediate t e s t i n g , therefore the cracking mechanism i s independent of surface f i l m thickness. Since over 92% of s l i p steps i n a-brass monocrystals at few % o s t r a i n are smaller than 500 A i n depth (83) i . e . much less than the thickness of the visible t a r n i s h f i l m , i t i s suggested that one s i n g l e s l i p step emergence process i n p o l y c r y s t a l s can only allow the access of environment to fresh metal to ensure further d i s s o l u t i o n i f the f i l m 'parts' rather than 'shears' (Fig. 76). At t h i s juncture, a simple cracking model i s considered: d i s l o c a t i o n s p i l e up at a prerequiste b a r r i e r (Figure 77A) and d i s s o l u t i o n occurs at the metal surface, penetrating into the metal (Fig. 77B) , As b a r r i e r #1 i s removed, s l i p occurs (Fig. 77C), and the f i l m parts. D i s s o l u t i o n of the f r e s h l y exposed metal occurs (Fig. 77D), followed by s l i p step emergence (Fig. 77E). This process can repeat u n t i l ; f i n a l mechanical fr a c t u r e . Since the oxide f i l m i s r e a d i l y parted by p l a s t i c deformation, and penetration of t a r n i s h into a-brass i s very expeditious (53,82) i n i t i a l l y , the f i r s t s l i p process cannot account forthe observed i n i t i a t i o n times (15-30 mins) i n load r e l a x a t i o n conditions. Moreover i t i s u n l i k e l y that one step w i l l produce any detectable change in, properties. Thus a new concept for " i n i t i a t i o n time" i s introduced: i t i s suggested that the ' i n i t i a t i o n period' corresponds to the s e t t i n g - 96 -up of a notch (Fig. 77F) co n s i s t i n g of several successive steps of metal d i s s o l u t i o n and s l i p step parting of a f i l m , and thus corresponds to a very slow rate of crack growth. It i s assumed that the marked increase i n cracking associated with notch formation occurs because stress i s concentrated i n such a manner to produce severe p l a s t i c deformation (perhaps c l o s e r and/or larger s l i p steps) ahead of the , advancing crack t i p . The c r i t e r i a f o r establishment of the notch are not known at present. Attention i s drawn to the following:-(1) the dissolved surface must show decreasing r e a c t i v i t y with time, otherwise cracking w i l l not be l o c a l i s e d . (2) Ou^O forms as an anodic product (S3) and the exposed bare metal at the crack t i p i s always anodic to the thick v i s i b l e t a r n i s h f i l m during cracking. (3) cracking depends on the s l i p plane spacing: .planes c l o s e r together w i l l allow a f a s t e r o v e r a l l cracking rate since d i s s o l u t i o n i s permitted to proceed i n several short stages rather than i n fewer stages over large distances. Although a c r i t i c a l s l i p step height i s not a p r e r e q u i s i t e of cracking, the magnitude of s l i p step may be important i n i n f l u e n c i n g such factors as the amount of fresh metal exposed to the environment at each step, oxide thickness (see F i g . 78) and therefore electrochemical variables such as anodic current density. S t r a i n and grain s i z e are expected to influence s l i p step height and s l i p l i n e spacing. Swann and Nutting (84) and Fourie and Wilsdorf (83) agree that s l i p l i n e spacing decreases and s l i p step height increases with increasing s t r a i n . - 97 -L i t t l e information (83,84) pertaining to the e f f e c t of grain s i z e on the d i s t r i b u t i o n of s l i p l i n e spacing and s l i p step height i n t h i s material i s a v a i l a b l e . However, for a constant s t r a i n , d i s l o c a t i o n density increases systematically (but s l i g h t l y ) with decreasing grain s i z e (85) and i t i s therefore assumed that, for a given strain,.there i s , s t a t i s t i c a l l y a smaller s l i p l i n e spacing (and s l i p step height) i n the smaller grains. 4.2.1. Predictions of the model In material of a constant grain s i z e , the e f f e c t of increasing stress i s to induce closer s l i p steps and therefore increase cracking tendency. Grain s i z e e f f e c t s are d i f f i c u l t to r a t i o n a l i s e i n the present uncertainty of d e t a i l . The s e t t i n g up of a p a r t i c u l a r notch depth i n material of two d i f f e r e n t grain sizes i s considered. See Figure 79. The model predicts that i n a p a r t i c u l a r environment a macroscopically i d e n t i c a l notch i s established i n a shorter time period i n the smaller grain sized material: thus the f i n e grain material i s more susceptible to f a i l u r e than the coarse grained structure i f s l i p step s i z e does not influence d i s s o l u t i o n k i n e t i c s . This s i t u a t i o n has, i n f a c t , been observed i n the continuous t e n s i l e loading experiments of Rosenburg and Cadoff (8). I f increasing s l i p step s i z e increases d i s s o l u t i o n k i n e t i c s and/or i f a smaller c r i t i c a l notch depth " i n i t i a t e s " f racture i n the coarse grained material, the coarser structure w i l l be more susceptible to cracking. This r e f l e c t s (with increasing grain size) i n shorter i n i t i a t i o n periods and shorter times to complete f a i l u r e i n load r e l a x a t i o n - 98 -conditions, and decreased elongation at fracture i n continuous loading conditions. In any event, i t i s anticipated that there w i l l be only a marginal e f f e c t of grain s i z e i . e . , i n continuous t e n s i l e loading t e s t s , f r a c t u r e w i l l be i n i t i a t e d at c l o s e l y s i m i l a r s t r a i n s . It i s -1/2 because of t h i s small s c a t t e r that the vs. D p l o t s exhibit a high degree of c o r r e l a t i o n . . pH i s assumed to influence the corrosion rate and s o l u b i l i t y of Cu^O. Foraconstant stress and grain s i z e ( i . e . s l i p step d i s t r i b u t i o n ) 'penetration' i s most rapid at pH 6.8, r e s u l t i n g i n a U shaped dependence of i n i t i a t i o n time with pH. T r a n s f e r r i n g the e f f e c t to con-tinuous t e n s i l e loading conditions also r e s u l t s i n the U shaped r e l a t i o n -ship (Fig. 14) since the e f f e c t i s simply to displace the f a i l u r e to higher stresses. There may be an a d d i t i o n a l e f f e c t of pH, however, i n that at pH of 8.0 and 5.0 the oxide f i l m may not be as passive, permitting a more general corrosion. This v a r i a t i o n of d u c t i l i t y with pH w i l l be observed at most s t r a i n rates (Fig. 14). The e f f e c t of increasing the crosshead speed i s to apply a p a r t i c u l a r stress i n a shorter time period. Although t h i s w i l l decrease the i n t e r s l i p l i n e spacing favouring rapid cracking through many d i s s o l u t i o n steps, i t w i l l also decrease the time a v a i l a b l e for s e t t i n g up of the notch. At the-higher rates of t e s t i n g , the metal exhibits d u c t i l e f a i l u r e (Fig. 14,15). The model predicts intergranular cracking because grain boundaries are l i k e l y to be the strongest b a r r i e r s to d i s l o c a t i o n s . There i s no need to postulate p r e f e r e n t i a l grain boundary penetration of the v i s i b l e -'99 -black tarnish or d i s s o l u t i o n of one component p r e f e r e n t i a l l y concentrated at the grain boundaries. Cracking of s i n g l e c r y s t a l s i s possible when ba r r i e r s of s u f f i c i e n t strength e.g. i n t e r s e c t i n g operative s l i p systems are created. The model predicts that cracking rate may be governed by deforma-t i o n processes or factors a f f e c t i n g the growth of the oxide f i l m . Based on the observation that the v i s i b l e black oxide i s porous (52), i t i s assumed that transport of species through the f i l m to the brass i n t e r f a c e rather than i o n i c d i f f u s i o n , i s a major contributor to f i l m growth. A c t i v a t i o n energies estimated from Figure 80 suggest that transport processes i n the l i q u i d are rate c o n t r o l l i n g at the high temperatures and deformation at low temperatures with a gradual change from one to the other over the temperature range, e.g. See Table VI In the case of l i q u i d metal embrittlement, no sudden change i n behaviour i s noted at the l i m i t i n g temperature of formation of the i n t e r m e t a l l i c compound, suggesting the i n t e r m e t a l l i c assumes no other, part i n the cracking process than that of a weak passive f i l m . However, i n contrast to stress corrosion cracking, i t i s anticipated that the surface f i l m severely l i m i t s further penetration into the metal; therefore, there may exist a high degree of s e n s i t i v i t y of cracking time to stress such that cracking time w i l l be slow when the thickness of the surface f i l m i s less than the i n t e r s l i p l i n e spacing, and very, rapid when i t exceeds the same (No information i s a v a i l a b l e pertaining to thickness of a mercury i n t e r m e t a l l i c f i l m on a-brass). In either case, the cracking w i l l be d i s s o l u t i o n c o n t r o l l e d . The d i s s o l v i n g zinc leaves a remnant Cu r i c h plus i n t e r m e t a l l i c layer i n l i q u i d metal embrittle-ment and a C U 2 O layer i n ammonia. See Appendix I I I . 100 TABLE VI Estimated Activation Energies System Activation Energies Lower Limit Higher Limit of temperature  a-brass - Mattsson's pH 6.8 (37) a-brass - Mattsson's (non-tarnish) (37) a-brass - Mattsson's pH 6.8 (3) a-brass - Mattsson's pH 6.8 (33) Be/Cu in Hg pH 6.8 (3) Be/Cu in Mattsson's pH 6.8 (3) 22 kcals/mole 7 kcals/mole 36 " 5 . " 6 " 2 " 14 " 4 " 20 " 3 " 12 " 3 " N.B. (i) Activation energy for creep of copper (Cross s l i p at 300°K) is experimentally determined at 23 kcals/mole (86). ( i i ) Activation energy for diffusion of cupric tetrammine ions in aqueous solution estimated at 2.9 kcals/mole (87). ( i i i ) Activation energy for diffusion of Hg in Hg is 1.0-1.8 kcal/mole (88). - 101 -Increasing zinc content of a-brass increases both the i n t e r s l i p l i n e spacing (84) and tarnishing rate (27) . I f the r i s e i n tarnishing rate i s greater than the increase i n i n t e r s l i p l i n e spacing, and provided that the surface f i l m i s s u f f i c i e n t l y passivating, the model can p r e d i c t that cracking time w i l l decrease with increasing zinc contents. 4.2.3. Quantitative elements of crack i n i t i a t i o n The model predicts that the c r i t i c a l notch f or fra c t u r e " i n i t i a t i o n " comprises of several i n d i v i d u a l d i s s o l u t i o n stages determined by the s l i p plane spacing x^. See F i g . 81. The time, t , taken f o r one step of x^ i s given by + !/2 x. = at l where a i s a constant. Assume that d i f f u s i o n i n the environment permits the same fresh m e t a l - l i q u i d i n t e r f a c i a l conditions to p r e v a i l at the s t a r t of each d i f f u s i o n step. Thus the i n i t i a t i o n time T £ = Nt, where N i s the number of d i s s o l u t i o n stages, i . e . x. T = (—) N c a J and since the c r i t i c a l crack length x = Nx. 2 x - 102 -Let a c r i t i c a l stress i n t e n s i t y be reached; the frac t u r e s t r e s s , Op i s re l a t e d to the c r i t i c a l crack length by a*(x ) 1 / 2 = K • f v cJ where K i s a constant r e l a t e d to the "fracture toughness" (section 1.2.3.2.). Thus ° f - N l / 4 a l / 2 T c Data from the curves for constant grain s i z e a-brass deformed i n continuous t e n s i l e loading conditions i n Mattsson's s o l u t i o n (Figs. 10-12) and mercury (Fig. 13) at a series of crosshead speeds, has been replotted i n the form of log vs. log T £ ( F i g . 82). The slopes corresponding to neutral Mattsson's s o l u t i o n are i n approximate accord with the above equation, but the mercury environment induced penetra-1/4 t i o n given by x^ = at on the basis of the model. It i s known (53, 82) that the ta r n i s h exhibits d i f f u s i o n c o n t r o l l e d growth, but no data i s a v a i l a b l e f o r mercury penetration. In view of the following examination of the assumptions, the exact r e l a t i o n s h i p observed i s fo r t u i t o u s . The above equation has several obvious shortcomings:-(1) the growth rate of each step w i l l not be the same i f there i s i n s u f f i c i e n t d i f f u s i o n i n the l i q u i d to prevent stagnancy of environment i n the crack. - 103 -(2) since decreases with increasing s t r a i n (83,84) i n continuous loading conditions, x £ w i l l consist of unequal x^. In load r e l a x a t i o n conditions, the assumption x = Nx. i s more reasonable. r c 1 (3) the foregoing equation assumes that K i s constant; since fr a c t u r e toughness changes with s t r a i n , K cannot be constant. (4) N i s not constant. • , • , The simple mathetmatical model doe's, however, pr e d i c t not unreason-able c r i t i c a l notch depths based on the s t r a i n variant i n t e r s l i p l i n e spacing data of Swann and Nutting (84) and u t i l i s i n g the ta r n i s h growth equation of McEvily and Bond (53). See Table VII. Estimates of x i n continuous t e n s i l e conditions do not take into c. account variance of x^ during deformation, therefore values f o r load r e l a x a t i o n conditions are tabulated using data of Tromans et a l . (29) and two values (84) of i n t e r s l i p l i n e spacing. TABLE VII Calculations of Xc from experimental data Crosshead Speed Fracture•Strain x. I t N X c "/min % o A sees y 0.0002 2.5 4,500 81.0 138 62 0.001 6.0 4,000 64.0 84 34 0.005 15.0 2,000 16.0 168 33 0.01 22.0 1,300 6.8 288 37 0.02 27.0 1,000 4.0 302 30 0.05 35.0 800 2.5 252 20 Load Relaxation Conditions: I n i t i a t i o n Period X. t N X l c o A sees y 1200 sees 4,000 64 18-19 8 II 1,000 4 300 30 - 104 -Cracking Time (Mins) 1 • * 1 • • ' 6.2 6.4 6.6 6.8 7.0 7.2 7.4 pH Mattsson's Solution Figure 75. Cracking time vs. pH (after Tromans, Dowds, Leja [29]). y T - 105 -Shearing i Parting Figure 76. Film Shearing and Parting Processes Figure 77. Model for Cracking. (1,2,3,4,5,6 Represent S l i p planes containing d i s l o c a t i o n s p i l e d up at b a r r i e r ) . Large S l i p Step Small S l i p Step A. P r i o r to i n i t i a l displacement B. D i s s o l u t i o n between s l i p planes p r i o r to second displacement C. Second displacement reveals larger amount of f r e s h l y exposed metal surface and oxide thickness on crack sides. Figure 78. E f f e c t of s l i p step si z e on exposure of fresh metal and on oxide thickness. - 107 -Figure 81. Mathematical Model f o r C r i t i c a l Notch. 3.0 - 108 -Temperature 1/T °K (x 10 - 3) 3.5 4.0 • l u a-brass - Mattsson's pH 6.8 a-brass non-tarnishing Mattsson's n „ . Be-Cu i n Hg. a-brass Mattsson's pH 6.8 Be-Cu i n NH3 a-brass - NH3 Pugh and Montague (37) ^ Pugh.and Montague (37) A Nichols and Rostoker ( 2 ) O Johnson (33) Q Nichols and Rostoker ( 2 ) , ft Nichols and Rostoker (2 ) -2 1 1/t n (sees) Figure 80. V a r i a t i o n of Cracking Time with Temperature. >rcury 5. CONCLUSIONS 1. P o l y c r y s t a l l i n e a-brass exhibits analogous behaviour i n a l i q u i d metal and a stress corrosion environment i n continuous loading conditions. P o l y c r y s t a l l a i n e B-brass exhibits s i m i l a r s t r a i n rate e f f e c t s and fractography i n these environments. Single c r y s t a l s of both a- and B-brasses f a i l prematurely i n mercurous n i t r a t e and neutral Mattsson's s o l u t i o n . 2. The stress a s s i s t e d cracking of a-brass i n ammoniacal copper sulphate and mercury can be accounted f o r by a common cracking mechanism: a modification of the t a r n i s h rupture model, based on s l i p step displacement of a passive surface f i l m has been developed. -1/2 3. Plots of a_ vs. D cannot be used as evidence for a surface f energy mechanism. ,4. A low angle electron d i f f r a c t i o n technique has been s u c c e s s f u l l y applied to the analysis of t h i n films on fractured surfaces. T 111 -BIBLIOGRAPHY 1. SMITH, W.R. and FORSYTH, P.E.J. Me t a l l u r g i a V. 34 p. 245-6 (1946). 2. NICHOLS,H. and ROSTOKER, W. Trans ASM V o l . 56 p. 494-506 (1963). 3. COLEMAN, E.G., WEINSTEIN, D. and ROSTOKER, W. Acta Met V. 9, p. 491-6 (1961) . 4. HEYN, E. J . Int. Metals, V ol. 12, p. 3 (1914). 5. 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Corrosion V. 19, p. 102 (1963) . 45. Ref. 6, p. 323. 46. HOAR, T.P. and HINES, J.C. "Stress Corrosion Cracking and Embrittle-ment", p. 107-125, Wiley, (1956). 47. HOAR, T.P., and WEST, J.M. Proc. Royal Soc. Series A, V. 268, p. 304-315 (1962). 48. PARDUE, W.M.,' BECK, F.H., and FONTANA, M.G. Trans A.S.M. V. 54, p. 539-548 (1961). 49. LOGAN, H.L. J . Res. Nat. Bureau of Standards, V. 48, p. 99 (1952). 50. CHAMPION F.A. in "Internal Stresses of Metals and A l l o y s " , I n s t i t u t e of Metals, London, p. 468 (1948). 51. PUGH, E.N. and WESTWOOD, A.R.C. P h i l . Mag. p. 167 (1966). 52. FORTY, A.J. and HUMBLE, P. P h i l . Mag. V. 8, p. 247 (1963). - 114 -53. McEVILY, A.J., J r . , and BOND, A.P. J . Electrochemical Soc,.Vo. 112, p. 131 (1965). 54. SMITH, C.S. Trans A.I.M.E. V. 175, p. 15-51 (1948). 55. ROBERTSON, W.D. Trans A.I.M.E., p. 1190-1191 (1951). 56. GREENWOOD, J.N. J . Inst. Metals, p. 177-178 (1953). 57. PETCH, N.J. P h i l . Mag. p. 331 (1956). 58. McCLEAN, D. "Mechanical properties of metals", p. 383, Wiley, (1962). 59. STOLOFF, N.S. and JOHNSTON, T.L. Acta Met, Vol. 11 p. 251-256 (1963). 60. GRIFFITH, A.A. Trans Royal Soc. A, V. 163 (1921). 61. LINGER, K.R. Private Communication with: Univ. B r i t i s h Columbia 62. WESTWOOD, A.R.C. and HITCH, T.T. J . App. Physics. V. 34 p. 3085 (1963). 63. STROH, A.N. Proc. Royal Soc. A. V. 232, p. 548 (1955). 64. PETCH, N.J. J . Iron and Steel I n s t i t u t e , V. 86, p. 456 (1957-58). 65. . OROWAN, E. Welding Journal, V. 34, p. 1575 (1955). 66. SPRETNAK, J.W. and GRIFFIS, C.A. Corrosion, V. 25, No. 5, p. 193 (1969). 67. ARMSTRONG, R, CODD, I, DOUTHWAITE, R.M. and PETCH, N.J. P h i l . Mag. Vo. 7, p. 45 (1962). 68. JAMIESON, R.E. and SHERRILL, F.A. Acta Met. V. 4, p. 197 (1956). 69. ARDLEY, G.W. and COTTRELL, Proc- Royal Soc. Series A, V. 219 p. 328 (1953) . 70. PHILIPS, W.L. J r . and ROBERTSON, W.D. Trans A.I.M.E. p. 406 June, 1958. 71. MURAKAMI, Y. and IKAI, Y. Trans Jap Inst. Metals Vol. 8, No. 4 p. 246, (1967). - 115 -72. DONALDSON, K.C. Met 455 Project Dept. Metallurgy, Univ. Br. Columbia (1966). 73. PREECE, CM. and WESTWOOD, A.R.C. Trans A.S.M. Quty V. 62, p. 419 (1969) . 74. BAILEY, A.R. Ref. 15, p. 124. 75. A.S.T.M. Powder D i f f r a c t i o n F i l e - Card No. 4-0811 f o r CuHg. 76. " " - " tt tt 5-0667 f o r Cu 20. 77. 11 11 " " 11 4-0836 f o r Cu 78. " " " - i i - i i 3-0879 f o r 6CuO Cu 20 (approx.) 79. KRAMER, I. and DEMER, L.J. "E f f e c t s of environment on mechanical properties of metals" Progress i n Material Science Vol. 9, No. 3. Pergamon Press, Oxford. 80. TROMANS, D. and NUTTING, J . Corrosion V. 21, p. 143 (1965). 81. BLACKWOOD, A.W. and STOLOFF, N.S. Trans A.S.M. Quty V. 62, p. 677-689 (1969). 82. PUGH, E.N., MONTAGUE, W.G. and WESTWOOD, A.R.C. As r e f . 37, p. 20. 83. FOURIE, J.T. and WlLSDORF, H.G.F. Acta Met. Vol. 7, p. 334 (1959). 84. SWANN, P.R. and NUTTING, J . J . Inst. Metals, V. 90, p. 133 (1961). 85. CHARNOCK, W. P h i l . Mag. V. 18, p. 89 (1968). 86. GARAFALD, F. "Fundamentals of Creep and Creep Rupture i n Metals" p. 78, MacMillan, 1965. 87. BURKIN, A.R. "Chemistry of Hydrometallurgical Processes" p. 53, 1966. 88. Metals Reference Handbook, Smithells ,Butterworth, 1967. 89. HANSEN, "Constitution of Binary A l l o y s " , 1958. - 116 -APPENDIX I Rostoker a c t u a l l y states (9) that embrittlement does not occur i n systems below the i n t e r m e t a l l i c compound formation temperature, and hence no embrittlement of copper by mercury occurs below 100°C. However, one may expect CuHg formation (89) i n brass. Rostoker apparently resolved the c o n t r o v e r s i a l point by claiming X-ray d i f f r a c t i o n analysis f a i l e d to detect CuHg on the surface of brass exposed to Hg, even a f t e r heating to. above 100°C and cooling. The electron d i f f r a c t i o n patterns of the present work support the formation of the i n t e r m e t a l l i c compound. It i s shown by the approximate c a l c u l a t i o n below that there i s a s i g n i f i c a n t increase i n volume when the i n t e r m e t a l l i c compound i s formed. . Thus even i f an i n t e r m e t a l l i c compound i s formed, i t should not completely hinder further d i s s o l u t i o n of zinc atoms because volume str a i n s produced w i l l probably lead to cracking and/or f l a k i n g of the i n t e r m e t a l l i c layer. Approximate c a l c u l a t i o n : -Data from CuHg Card i n ASTM Powder D i f f r a c t i o n Index (75) o L a t t i c e Parameter>c= 9.4 A No. of CuHg u n i t s / u n i t c e l l o s 25 (9 4x10 ) Volume occupied by 1 Cu atom i n the CuHg cpd = V„ - ' -— cc L UCuHg o Z b L a t t i c e Parameter of 70/30 brass = 3.676 A Volume occupied by 1 Cu atom i n 70/30 brass = ^ ^ — cc. Ratio C u a = 0.53. CuHg - 117 -APPENDIX II Corrosion Fatigue of Copper and ot-Brass Only one instance of stress corrosion of pure copper has been reported (14) and OFHC copper d i d not s u f f e r embrittlement under continuous t e n s i l e loading conditions i n te s t s performed i n Mattsson's s o l u t i o n concurrent to the main work. The e f f e c t of fatigue deformation was therefore examined. The e f f e c t of pH of the environment was recorded since t h i s parameter severely influences the fatigue l i f e of a-brass (11, f i g . A l ) . Procedure 3/8" O.F.H..C. copper and a-brass rods were machined to fatigue specimens ( f i g 5C) which were annealed f o r 1 hour at 550°C and f i n i s h e d to 0000 grade emergy. Three tests were performed under each set of conditions shown i n Table AI using a r o t a t i n g beam bending fatigue machine (at 1700 r.p.m.). The environment was introduced into the c e l l through a hypodermic needle between the specimen and t e f l o n tubing ( f i g . 5G) and not through the t e f l o n wall so as to eliminate a i r bubbles. Results are given i n Table A2. Comments The fatigue l i f e of a-brass was reduced by a f a c t o r of between 6 and 10 by neutral Mattsson's s o l u t i o n , i n accord with Northcott (11). It i s apparent that both the i n i t i a t i o n and propagation stages are accelerated by the presence of t h i s environment. - 118 -Each environment has the e f f e c t of reducing the fatigue l i f e of copper by a f a c t o r of 2-2 1/2 from that i n pure water. Neither ammonia nor pH v a r i a t i o n (over the range examined) had any s i g n i f i c a n t e f f e c t on fatigue l i f e . 119 TABLE A l Testing Conditions Material Environment Stress (psi) a-brass Mattsson's s o l u t i o n (pH 6.8) D.D.B:. Water 30,000 30,000 Copper Mattsson's s o l u t i o n (pH 6.8) " 5.5 " " 8 . 0 " " 4.9 " 4.4 NaCl " 5.0 CuS0 4 " 4.7 D.D.B. Water 21,000 A l l tests were performed at ambient temperatures. - 120 -TABLE A2 Results Material Environment Fatigue L i f e (cycles) a-brass D.D.B. Water 162,210 181,800 180,170 Mattsson's pH 6.8 30,510 16,920 22,100 30% l i f e i n water (55,000 c.) 93,600 then f a i l i n Mattsson's pH 6.8 77,850 65,810 Copper D.D.B. Water 236,870 225,470 193,950 Mattsson's pH 6.8 75,490 100,470 81,380 " " 8.0 81,740 103,170 90,370 " " 5.5 110,420 81,240 64,360 " " 4.9 110,470 86,320 91,980 " " 4.4 100,400 120,560 98,810 0.04 M CuSO 86,310 110,120 93,700 NaCl 76,870 108,310 97,770 r 10 pH Mattsson's Solution Figure A l . E f f e c t of pH of Mattsson's Solution on Fatigue L i f e of a-Brass r- 122 -APPENDIX III (1) Assuming the re a c t i o n proceeds p r i n c i p a l l y by d i s s o l u t i o n of zinc, i t i s required to determine i f the zinc contained between adjacent s l i p planes can f e a s i b l y d i s s o l v e i n a surface layer of Hg. Consider an annulus of mercury (width T cm) surrounding a 70/30 brass specimen of radius R cm. Assume un i t length of specimen l i q u i d 23 M * * ^ 2 T r R T d u - X 6 x 10 , , u No. of atoms ^  _Hg d„ = density Hg 200 g I f s o l i d s o l u b i l i t y of zinc i n Hg at R Temp = X% „„ X2-rrRt d H 6 x 10 Max No. of zinc atoms held i n volume = n n n  100 x 200 In 70/30 brass - No. of zinc atoms per unit c e l l = 0.3 x 4 " ° 3 - these occupy (3.676 A) 1 2 - No. of zinc atoms/cc brass = 3 -24 (3.676) xlO Consider d i s s o l u t i o n of an annulus of a-brass to depth P V, 2i:RP cc brass 1 2 No. of zinc atoms = 2irRP x 3 -24 (3.676)° x 10 Z 4 Assuming a l l the zinc goes into s o l u t i o n , and the s o l u t i o n i s only h a l f saturated 2400 P x 200 T = d H g x 6 x X x (3.676) 3 Assigning values of X = 5, d^ g = 13.6 ( r e f s . 88,89) - 123 -T = 3.03 x 10 2P A reasonable value f or P = 2000CA (distance between s l i p planes) hence, T^=.6 x 10" 3 ": which i s a reasonable estimate of a "surface l a y e r " thickness. However, i n very d i l u t e solutions of mercurous n i t r a t e ( i . e . very l i t t l e mercury on specimen surface), saturation w i l l soon be reached and penetration w i l l be slow. o (2) It i s shown that zinc atom can d i f f u s e 1000 A i n an extremely short time at ambient temperature Assume D = 0 .85 .x 10~ 4 e c m 2 / s e c r e f . (88) = 0.16 x 10" 4 Assume x^a?/Dt cm 2 x t = — sec I D " 1 0 0.16 x 10-t ^ i r 10 ^ sees, 

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