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Stress corrosion cracking of aluminum alloys Pathania, Rajeshwar Singh 1970

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STRESS CORROSION CRACKING OF ALUMINUM ALLOYS  BY  RAJESHWAR SINGH PATHANIA B.Sc. (Hons.) D e l h i U n i v e r s i t y , 1963 B.E. (Met.) I n d i a n I n s t i t u t e of S c i e n c e , Bangalore, 1965  A THESIS SUBMITTED IN PARTIAL FULFILMENT OF THE REQUIREMENTS FOR THE  DEGREE OF DOCTOR OF PHILOSOPHY  i n the Department of METALLURGY  We accept t h i s t h e s i s as conforming required  to the  standard  The U n i v e r s i t y o f B r i t i s h F e b r u a r y , 1970.  Columbia  In  presenting  this  an a d v a n c e d  degree  the  shall  I  Library  further  for  scholarly  by h i s of  agree  this  written  thesis  in  at  University  the  make  tha  it  for  financial  is  A p r i l 6,  1970  of  Columbia,  British for  by  the  gain  Columbia  shall  not  the  requirements  reference copying of  I  agree  and  copying or  be a l l o w e d  for  that  study.  this  thesis  Head o f my D e p a r t m e n t  understood that  Metallurgy  The U n i v e r s i t y o f B r i t i s h V a n c o u v e r 8, C a n a d a  of  for extensive  p u r p o s e s may be g r a n t e d It  fulfilment  available  permission.  Department o f  Date  freely  permission  representatives. thesis  partial  or  publication  without  my  ABSTRACT The s t r e s s c o r r o s i o n behaviour of p r e c i p i t a t i o n hardened Al-9Mg, Al-22Zn i n aqueous environments  and Al-3Mg-6Zn a l l o y s has been s t u d i e d and e t h a n o l .  The  stress  corrosion  s u s c e p t i b i l i t y d e f i n e d as the r e c i p r o c a l of f a i l u r e  time  has been i n v e s t i g a t e d as a f u n c t i o n of a l l o y - e n v i r o n m e n t i s o t h e r m a l aging treatment, m i c r o s t r u c t u r e , a p p l i e d and temperature  u s i n g smooth and notched  l o a d t e s t s , l o a d - r e l a x a t i o n t e s t s and environments  have been used  a l s o been c a r r i e d out i n aqueous  tensile  specimens.  stress,  Constant  tensile tests i n different  to e v a l u a t e the s t r e s s  c h a r a c t e r i s t i c s of aluminum a l l o y s .  system,  corrosion  A l i m i t e d study of Mg-9A1 has  environments.  The p r o c e s s of s t r e s s c o r r o s i o n g e n e r a l l y c o n s i s t e d of t h r e e p a r t s : 1) A slow i n i t i a t i o n stage 2) a r a p i d p r o p a g a t i o n stage 3) mechanical  f r a c t u r e due to t e n s i l e o v e r l o a d .  the i n i t i a t i o n time was The  With a few e x c e p t i o n s ,  g r e a t e r than the p r o p a g a t i o n time.  c r a c k i n i t i a t i o n and p r o p a g a t i o n r a t e s were s t r e s s  and  t h e r m a l l y a c t i v a t e d and c o u l d be expressed by a g e n e r a l e q u a t i o n of the form Rate = A  q  a  11  . exp  (^|)  where a i s the a p p l i e d t e n s i l e s t r e s s , Q i s the apparent energy  of the r a t e c o n t r o l l i n g p r o c e s s and A  a g i v e n a l l o y - e n v i r o n m e n t system. the r a t e c o n t r o l l i n g p r o c e s s was  The  q  and n are c o n s t a n t s f o r  apparent  different  activation  a c t i v a t i o n energy  i n the two  of  environments.  I t a l s o changed between i n i t i a t i o n and p r o p a g a t i o n s t a g e s .  The  aluminum a l l o y s when ranked  s u s c e p t i b i l i t y were:  1) Al-3Mg-6Zn, 2) Al-9Mg, 3) Al-22Zn.  a l l o y s which were g i v e n heat o f coherent  i n order of i n c r e a s i n g  treatments  or p a r t i a l l y coherent  c o r r e l a t i n g to the  phases, were found  The presence  to be most prone  to s t r e s s c o r r o s i o n c r a c k i n g . The  environments p l a c e d i n an o r d e r o f i n c r e a s i n g a g g r e s s i v e -  ness were d e s s i c a n t - d r i e d a i r , double  d i s t i l l e d water, e t h a n o l , ambient  a i r , d e i o n i z e d water and NaCl/K^CrO^ s o l u t i o n . s u s c e p t i b l e aluminum a l l o y s was  found  to be  The  d u c t i l i t y of ,  s i g n i f i c a n t l y decreased  by  NaCl/K^CrO^ and d e i o n i z e d water at low s t r a i n r a t e s and enhanced by dessicant-dried a i r . Fractography  showed the c r a c k i n g to be i n t e r g r a n u l a r i n  aluminum a l l o y s and t r a n s g r a n u l a r i n the Mg-Al a l l o y . c o r r o s i o n s u r f a c e was w h i l e the m e c h a n i c a l l y dimples  caused  fractured surface exhibited s l i p  unsatisfactory.  steps  and  examined i n l i g h t and was  Models i n v o l v i n g e i t h e r d i s s o l u t i o n o r  a l o n e were a l s o inadequate  found  of  t o be  deformation  i n e x p l a i n i n g the p r e s e n t r e s u l t s .  Therefore  p o s t u l a t e d which i n v o l v e s the g e n e r a t i o n o f a  continuous  path o f chemical h e t e r o g e n e i t y by s h e a r i n g and p r e c i p i t a t e s f o l l o w e d by t h e i r d i s s o l u t i o n . i n the deformation from i n i t i a t i o n  appearance  by v o i d f o r m a t i o n .  hydrogen c h a r g i n g experiments and o t h e r evidence  model was  stress  c h a r a c t e r i s e d by a rough o r corroded  The hydrogen mechanism o f c r a c k i n g was  a new  The  to  process  l i n k up of  The  coherent  rate controlling  i s b e l i e v e d to change d u r i n g the  propagation.  step  transition  - iv The p o s t u l a t e d model i s c o n s i s t e n t w i t h t h e p r e s e n t r e s u l t s but i t s f u r t h e r development must a w a i t b e t t e r knowledge o f the k i n e t i c s o f d i s s o l u t i o n o f p r e c i p i t a t e s and t h a t o f d e f o r m a t i o n p r o c e s s e s a t the c r a c k t i p .  -  V -  TABLE OF CONTENTS Page Chapter 1.  Chapter 2.  Chapter 3.  INTRODUCTION  1  1.1.  C r i t i c a l Review of Stress Corrosion Theories  3  1.2.  Film Rupture Mechanisms  4  1.3.  Mechanisms Involving Precipitate Distribut i o n and Precipitate Free Zone  5  1.4.  Stress Concentration Hypotheses  7  1.5.  Environmental  Factors  9  1.6.  The Unresolved Points  9  1.7.  Objectives of the Present Work  10  EXPERIMENTAL PROCEDURE  12  2.1.  Starting Materials and Alloy Preparation...  12  2.2.  Fabrication by Rolling  13  2.3.  Extrusion  13  2.4.  Heat Treatment  15  2.5.  Polishing, Degreasing  2.6.  Constant  2.7.  Load Relaxation Tests  19  2.8.  Strain Rate Tests  19  2.9.  Hydrogen Charging Experiments  20  2.10.  Optical Microscopy  20  2.11.  Electron Microscopy  21  RESULTS AND OBSERVATIONS  22  3.1.  22  and Tensile Tests....  Load Stress Corrosion Tests  Isothermal Aging of Al-9Mg  17 17  - vi Page 3.2.  Al-3Mg-6Zn  22  3.3.  Mg-8.5A1  26  3.4.  Stress Corrosion Testing Techniques .......  26  3.4.1.  Stress Corrosion S u s c e p t i b i l i t y . . . .  30  3.4.2.  Constant Load Tests  30  3.4.2.1.  Al-9Mg  30  3.4.2.2.  Al-3Mg-6Zn  34  3.4.2.3.  Mg-8.6A1  34  3.4.3.  3.4.4.  Strain Rate Tests  39  3.4.3.1.  Al-2.6Mg-6.3Zn  39  3.4.3.2'.  Al-21.5Zn  42  Load Relaxation Tests f o r Different Alloy-Environment Systems  3.5.  The pH of the NaCl/K Cr0  3.6.  Activation Energy Analysis  50  3.6.1.  Al-21.5Zn  52  3.6.2.  Applied Stress vs Failure Time ....  55  3.6.3.  Al-2.6Mg-6.3Zn  57  3.6.4.  Al-8.6Mg  61  3.6.5.  Apparent Activation Energy During  2  4  Solution  42  Propagation  50  61  3.7.  Hydrogen Changing Experiments  62  3.8.  Electron Microscopy  68  3.8.1.  P r e c i p i t a t i o n D i s t r i b u t i o n After Aging  68  - viiPage  3.8.2. Chapter 4.  Fractography  77  DISCUSSION 4.1.  93  Precipitation Hardening and Mechanical Properties  93  4.1.1.  Al-9Mg Aged at 200°C  93  4.1.2.  Al-3Mg-6Zn Aged at 160°C  95  4.1.3.  Al-21.5Zn Aged at Room Temperature.  96  4.1.4.  Mg-8.5A1 Aged at 200°C  97  4.1.5.  Aged Hardening Mechanism  98  4.1.6.  Effect of Aging on Mechanical Properties of Aluminum Alloys  4.2.  Parameters Affecting  100  Stress Corrosion  Cracking i n Aluminum Alloys 4.2.1.  Stress Corrosion Cracking and Microstructure i n Al-9Mg Alloys  4.2.2.  105  ...  Applied Stress and Stress Corrosion Susceptibility  4.2.3.  108  Applied Stress and Stress Corrosion Susceptibility  4.3.  107  Stress Corrosion S u s c e p t i b i l i t y and Microstructure i n Al-3Mg-6Zn Alloy  4.2.5.  106  Nature of S l i p and Stress Corrosion Susceptibility  4.2.4.  105  Stress Corrosion S u s c e p t i b i l i t y of Mg-8.5A1  Ill 118  - viii Page 4.4.  Variation i n Stress Corrosion S u s c e p t i b i l i t y with the Alloy System 4.5.1.  120  I n i t i a t i o n and Propagation of Stress Corrosion Cracks i n NaCl/K^CrO^ . Solution  4.5.2.  123  The Relative Lengths of I n i t i a t i o n and Propagation Times i n NaCl/K^CrO^ Solution  4.6.  The Effect of Different Environments on the Stress Corrosion Behaviour of Aluminum Alloys  130  4.6.1.  Stress Corrosion i n Alcohols  136  4.6.2.  Effect of Environment on Mechanical Properties of Al-2.6Mg-6.3Zn  4.6.3.  4.6.4.  137  E f f e c t of Environment and Strain Rate on Mechanical Properties of  ;  Al-21.5Zn  140  Effect of Strain Rate on Serrated Yielding  4.7.  129  140  Temperature Dependence of the Stress Corrosion Process  143  4.8.  The Hydrogen Mechanism  145  4.9.  The Requirements of a Satisfactory Model...  154  4.9.1.  Models Involving Either Dissolution or Deformation and Their Drawbacks  4.10. The Proposed Model  155 158  - ix Page 4.10.1. The Observations Explained by the Proposed Model  165  4.11. Conclusions  167  4.12. Suggestion f o r Further Work  168  Appendix I  1^9  Appendix II  177  Bibliography  180  - X -  LIST OF FIGURES Figure 1. 2.  Page Types of t e n s i l e and stress corrosion specimens used i n the present work  14  Stress corrosion c e l l s used i n the the present work  16  3.  Modified creep stand f o r constant load tests  18  4.  P r e c i p i t a t i o n hardening of Al-9Mg, aged at 200°C  23  5.  P r e c i p i t a t i o n hardening of Al-3Mg-6Zn, aged at 200°C.  24  6.  P r e c i p i t a t i o n hardening of Mg-8.6A1, aged at 200°C...  25  7.  Failure time vs aging time f o r Al-9Mg aged at 200°C.  32  8.  Applied stress vs f a i l u r e time for Al-9Mg aged f o r 7 hours at 200°C . Effect of aging time at 160°C on the f a i l u r e time of Al-3Mg-6Zn i n NaCl/K Cr0  35  Applied stress vs f a i l u r e time f o r the aged Al-3Mg6Zn a l l o y  36  11.  Failure time vs aging time f o r Mg-8.6A1 aged at 200°C  37  12.  Effect of applied stress on the f a i l u r e time of solut i o n treated and quenched Mg-8.6A1  38  Strain rate vs uniform and t o t a l elongation f o r solut i o n treated and quenched Al-2.6Mg-6.3Zn a l l o y  40  Strain rate vs uniform and t o t a l elongations f o r an overaged Al-2.6Mg-6.3Zn a l l o y  41  Effect of d i f f e r e n t media on crack i n i t i a t i o n and propagation times i n Al-2.6Mg-6.3Zn a l l o y i n solution treated and overaged conditions  45  Effect of d i f f e r e n t media on i n i t i a t i o n and propagation times i n Al-21.5Zn alloy  49  Logarithm of crack i n i t i a t i o n rate vs 1/T for Al-21.5Zn in ethyl alcohol  53  9.  2  10.  13. 14. 15.  16.  17.  4  33  - xi Figure 18.  19. 20. 21.  Page Logarithm of crack i n i t i a t i o n rate vs 1/T f o r Al-21.5Zn i n ethyl alcohol '  54  The effect of applied stress on log f a i l u r e time of Al-21.5Zn a l l o y i n 99.9% ethyl alcohol  56  Logarithm of crack i n i t i a t i o n rate vs 1/T for an overaged Al-2.6Mg-6.3Zn a l l o y  58  Logarithm of crack i n i t i a t i o n rate vs 1/T f o r Al-2.6Mg-6.3Zn i n NaCl/K Cr0 solution and water  59  2  4  22.  Logarithm of crack i n i t i a t i o n rate vs 1/T for aged Al-8.6Mg a l l o y i n ethyl alcohol  23.  Logarithm of propagation rate vs 1/T for Al-2.6Mg6.3Zn i n NaCl/K Cr0 and D. water  63  Logarithm o f propagation rate vs 1/T for Al-Mg-Zn and Al-Mg alloys i n ethyl alcohol  64  Logarithm of propagation rate vs 1/T for Al-21.5Zn i n 95% ethyl alcohol  65  Logarithm of propagation rate vs 1/T for Al-21.5Zn i n 99% ethyl alcohol  66  Al-9Mg aged 7 hours at 200°C. d i s t r i b u t i o n , 10,000 x  Matrix p r e c i p i t a t e 70  28.  Al-9Mg aged 7 hours at 200°C.  G.b. ppt 10,000 x  29.  Al-9Mg aged 5 hours at 200°C. tions, 20,000 x  G.b. ppt and d i s l o c a -  2  24. 25. 26. 27.  4  30a,b Al-9Mg aged at 200°C.  Showing the p r e c i p i t a t e free  zone 31a-d Al-9Mg aged at 200°C.  70  71 Precipitate d i s t r i b u t i o n  72  32a,b-Difference i n p r e c i p i t a t e size between the aged A109Mg and Al-3Mg-6Zn  73  33a,b Al-3Mg-6Zn aged at 160°C for a) 1 hour b) 3 hours ...  74  34a-c Al-9Mg aged 5 hours at 200°C. Strained a) 1.3% b) 4%, c) 4% 35. Mg-8.6A1 aged for 1 hour at 200°C. Transgranular stress corrosion crack 40 x  75 76  - xii Figure 36.  Page Mg-8.6Zn, solution treated and water quenched. Transgranular stress corrosion crack 80 x  76  Aged Al-22Zn a l l o y , stress corroded i n NaCl/K CrO at 10,000 p s i . Complete stress corrosion f a i l u r e . . . .  81  Al-22Zn, hydrogen charged for 30 minutes and stress corroded i n NaCl/K C r 0 at 10,000 p s i  84  Al-2.6Mg-6.3Zn, solution treated and water quenched, stress corroded i n NaCl/F^CrC^ at 55,000 p s i . Predominantly stress corrosion surface  87  Aged Al-9Mg, stress corroded i n NaCl/I^CrG^ at 7,000 p s i . Predominantly stress corrosion surface...  89  Al-22Zn stress corroded i n 99.9% ethanol at 30,000 p s i . Stress corrosion and t e n s i l e overload modes of f a i l u r e  91  42.  Phase diagrams of Al-Mg and Al-Zn systems  92  43.  Plates of equilibrium Mg^yAl-^2 phase with basal habit plane . . .  37a-e  38a-h  4  39a-e  40a-d 41a-f  44. 45.  Schematic v a r i a t i o n of y i e l d stress with aging time for an age-hardening system  102  Log-log plot of applied stress vs f a i l u r e time for Al-9Mg i n NaCl/K Cr0  113  Log-log plot of applied stress vs f a i l u r e time for Al-9Mg i n NaCl/K Cr0  114  Log-log plot of applied stress vs i n i t i a t i o n time for Al-22Zn i n ethanol  116  Log-log plot of applied stress vs propagation rate for Al-22'Zn i n ethanol  117  Applied stress vs f a i l u r e time for three systems studied under constant load i n NaCl/K Cr0  122  Three types of load relaxation curves observed i n the present work  125  Variation i n stress i n t e n s i t y factor Kj during c r i t i c a l crack propagation  128  2  46.  2  47. 48. 49.  4  4  2  50a-c 51.  99  4  sub-  - xiii Figure  Page  52.  I l l u s t r a t i o n of the proposed model  53a-h  Tensile fracture surface of aged, notched Al-22Zn a l l o y fractured i n dessicant dried a i r at crosshead speeds of 20 in/min (a-e) and 0.01 in/min (f-h).  170  Notched, aged Al-2.6Mg-6.3Zn a l l o y pulled to fracture in a i r at 0.01 in/min  173  Aged Al-9Mg pulled to fracture i n a i r at 0.01 in/min.  , -  54a-f 55a,b  159  7I  - xiv LIST OF TABLES Table  Page  1.  Chemical Analysis of Fabricated Alloys  12  2.  Extrusion Condition of Alloys  15  3.  Summary of Stress Corrosion Tests and Testing Conditions  27  4.  Effect of Strain Rate on D u c t i l i t y and Yield Strength, of Al-21.5Zn i n Dry A i r and Deionized Water  42  Crack I n i t i a t i o n and Propagation Times for Al-3Mg-6Zn in NaCl/K Cr0 Solution  43  Crack I n i t i a t i o n and Propagation Times i n Different Environments  44  Crack I n i t i a t i o n and Propagation Times f o r Al-8.6Mg in NaCl/K Cr0 Solution  46  Crack I n i t i a t i o n and Propagation Times f o r Mg-8.6A1 i n NaCl/K Cr0 Solution  47  Effect of Different Media on the I n i t i a t i o n and Propagation Times f o r Al-21.5Zn Alloy  48  Activation Energies of Crack I n i t i a t i o n f o r Al-21.5Zn i n Ethanol-Water Mixtures  52  Activation Energies f o r Al-21.5Zn From Log Rate/T vs 1/T Plot  55  Hydrogen Charging Experiments on Al-21.5Zn i n NaCl/ K C r 0 at 10,000 p s i  67  Hydrogen Charging Experiments on Aged Al-2.6Mg-6.3Zn Alloy i n NaCl/K Cr0 at 76,500 p s i  68  5.  2  6. 7.  2  8.  2  9. 10. 11. 12.  2  13.  4  4  4  4  2  14.  4  Relative Fractions of I n i t i a t i o n and Propagation Times of Different Alloys i n NaCl/K Cr0 Solution  131  Relative Lengths of I n i t i a t i o n and Propagation Times i n Various Alloy Environment Systems  132  The Apparent Activation Energies for Crack I n i t i a t i o n and Propagation and the Possible Rate Controlling Processes  146  2  15. 16.  4  - XV  -  Table 17.  Page Apparent Activation Energies of Crack and Propagation f o r Aluminum Alloys  Initiation 155  18.  Proposed Activation Energies  162  19.  The Correlation Coefficient and Error i n Apparent Activation Energy from the Arrhenius Plots  178  - xvi -  ACKNOWLEDGEMENTS The author wishes to thank his research adivsor Dr. D. Tromans for his keen interest and valuable advice throughout this project. Thanks are also extended to the members of his Ph.D. committee, Dr. Peters and Dr. Lund and to other members of the department for t h e i r useful suggestions.  The assistance of the technical s t a f f p a r t i c u l a r l y  Mr. Arvid Lacis i s appreciated. The f i n a n c i a l support  i n the form of a National Research  Council Scholarship i s g r a t e f u l l y acknowledged.  Errata  p.  6  last  p.  8  l i n e 18  reiterate  p. 10  line 2  but i s immune  p. 13  line 2  homogenized  p. 20  line 2  part  p. 33  Abcissa  F a i l u r e time sec. 10  p. 50  line 9  alloys  p. 54  Ordinate  log I n i t i a t i o n Rate/T  line  p. 57  enhanced  4  t  f  , -n = k a x  p. 61  l i n e 20  fatigue i n aluminum  p. 62  l i n e 22  c a r r i e d out  p. 96  l i n e 11  comparable ;  p. 120  line 9  matrix  l i n e 14  overaged  l i n e 24  O.lu radius  p. 150  .0  INTRODUCTION  1.  Stress corrosion cracking can be defined as the rupture of metal, taking the form of cracks that may occur under the conjoint influence of a s p e c i f i c corrosive environment and applied or residual t e n s i l e stresses (1) well below the ultimate t e n s i l e strength.  A  "crack" can further be defined as a penetration, i n which the depth i s several order of magnitude greater than the width.  Failures promoted  by cross sectional area reduction by corrosion, are excluded  from the  term stress corrosion cracking (SCC). Metals of high strength and d u c t i l i t y can f a i l c a t a s t r o p h i c a l l y by stress corrosion cracking.  The phenomenon i s insidious, because  there may be no macroscopic indications of impending f a i l u r e and yet i t may result i n the explosion of a b o i l e r , the f a i l u r e of a s t e e l hook on a building s i t e or the f a i l u r e of landing gear struts on an airplane (2).  Many stress corrosion f a i l u r e s are undramatic and include leaks  in tanks, unspectacular f a i l u r e s of bolts and f a i l u r e of stainless steel vessels used i n various industries. Some other features which make stress corrosion cracking a major problem are (3) : (i)  Alloys subject to cracking are normally considered passive  - 2 and may  take years to be corroded to unusable dimensions and yet i n the  presence of a stress may (ii) environments may  f a i l i n several hours.  Extremely small quantities of apparently cause cracking; e.g. water may  innocous  cause cracking i n  aluminum alloys and high strength s t e e l s ; s a l t water may  stress corrode  austenitic s t a i n l e s s s t e e l and titanium a l l o y s . (iii)  The stress corrosion cracking of aluminum and magnesium  alloys assumes special importance i n view of the fact that these alloys are widely used i n a i r c r a f t f i t t i n g s and i n applications where a high strength to weight r a t i o i s required.  The most susceptible aluminum  alloys are those that can be strengthened by heat treatment and these are,of course,most useful as load carrying members i n aerospace industry (2), automotive applications, r a i l r o a d and marine applications and chemical process application. The objective of this project was  to study the stress corrosion  characteristics of aluminum base a l l o y s , s p e c i f i c a l l y Al-9Mg, Al-3Mg-6Zn, and Al-21.5Zn with a view to developing cracking.  a mechanism of stress corrosion  A limited study of Mg-8.5A1 was  also carried out.  The  p r e c i p i t a t e i n Al base a l l o y i s believed to be anodic to the matrix whereas i n Mg base alloy i t i s cathodic  (4)  (5) to the s o l i d solution.  No general mechanism of stress corrosion i s believed to exist (6).  There are a number of processes operating under s p e c i f i c metallurgical  and environmental conditions and only a few of these operating may  be necessary for cracking to occur.  Stresses may  together  give r i s e to  e l a s t i c and p l a s t i c deformations r e s u l t i n g i n f i l m rupture, d i s l o c a t i o n precipitate interaction and crack-stress f i e l d i n t e r a c t i o n .  The  - 3 -  environment may  cause i o n i c d i f f u s i o n into oxide layers, d i s s o l u t i o n of  layers, adsorption reactions  at cracktip or fresh metal  surface,  embrittlement by hydrogen and anodic d i s s o l u t i o n by chemical or electrochemical  1.1.  processes.  C r i t i c a l Review of Stress Corrosion Theories A number of mechanisms have been put  stress corrosion cracking  of aluminum base a l l o y s .  theory of stress corrosion of aluminum a l l o y s Brown and Dix  (7) i s as  forward to explain The  described  the  generalized by Mears,  follows:  "Corrosion occurs along l o c a l i z e d paths such as grain boundaries producing f i s s u r e s .  These p r e f e r e n t i a l l y corroded paths may  strata of low inherent resistance to corrosion or they may to the adjoining metal.  represent  be anodic  Components of t e n s i l e stress normal to the path  create a stress concentration at the base of the f i s s u r e .  At s u f f i c i e n t  concentration of stress, the fissures open further, thus exposing unfilmed metal to the corrosion attack.  The  fresh metal i s anodic and there i s  an acceleration of corrosion r e s u l t i n g i n further separations of metal". Thus the effects of corrosion and stress are mutually reinforcingiand result i n an accelerating crack which leads to a reduction  i n area  s u f f i c i e n t to cause mechanical fracture.  ,  The model proposed by Mears et a l . , does not give a detailed mechanistic interpretation of the stress corrosion phenomenon.  Other  models have been proposed, which interpret the observed stress  corrosion  behaviour i n terms of the rupture of protective oxide, the size of the  - 4 -  p r e c i p i t a t e f r e e zone (PFZ) next t o the g r a i n boundary, the i n the PFZ, of  and  the n a t u r e o f p r e c i p i t a t e produced  deformation  during aging.  the r e s u l t s and t h e i r i n t e r p r e t a t i o n s are c o n t r a d i c t o r y .  Some  There i s  no s i n g l e t h e o r y which can c o m p l e t e l y e x p l a i n the s t r e s s c o r r o s i o n c h a r a c t e r i s t i c s o f aluminum base a l l o y s .  1.2.  F i l m Rupture Mechanisms Logan (11) b e l i e v e s t h a t the p r o t e c t i v e oxide f i l m must  r u p t u r e d by p l a s t i c d e f o r m a t i o n proceed.  Corrosion results  i n order f o r c o r r o s i v e attack to  i n r e p a i r o f the f i l m which i s once a g a i n  r u p t u r e d by p l a s t i c d e f o r m a t i o n , to  be  l e a d i n g t o more c o r r o s i o n .  t h i s theory s t r e s s c o r r o s i o n crack i n the a l l o y proceeds  e l e c t r o c h e m i c a l p r o c e s s e s alone and no mechanical  According by  r u p t u r e o f the  alloy  takes p l a c e . F i l m r u p t u r e by i t s e l f i s not s u f f i c i e n t corrosion cracking.  to cause s t r e s s  Pure aluminum does not s t r e s s c o r r o d e even when  the p a s s i v e oxide f i l m i s r u p t u r e d by s l i p s t e p f o r m a t i o n . d u c t i l e aluminum a l l o y s e x h i b i t  Many  excellent stress corrosion resistance.  Another o b j e c t i o n i s r a i s e d by the work of Farmery and Evans They allowed a crack t o propagate  p a r t of the way  through  of  s u s c e p t i b l e Al-7Mg a l l o y and then a r r e s t e d the c r a c k by  of  c a t h o d i c c u r r e n t f o r 30 minutes.  crack d i d not propagate appeared  elsewhere.  When the c u r r e n t was  a g a i n although  I f mechanical  (14).  a specimen application stopped,  25 hours l a t e r f r e s h  the  cracks  f i l m r u p t u r e at the crack t i p was  the d e t e r m i n i n g f a c t o r f o r p r o p a g a t i o n , the c r a c k should have resumed  - 5 -  a f t e r the d i s c o n t i n u a t i o n of the p r o t e c t i v e c u r r e n t , because o f the heavy p l a s t i c d e f o r m a t i o n at the crack t i p . Gruhl process.  The  (15)  first  describes  stress corrosion cracking  stage i s l a r g e l y chemical or  whereas the t r u e crack  formation  metal p h y s i c a l p r o c e s s .  stress corrosion cracking.  corrosion cracking  t o l e r a t e i n i t i a t i o n of  1.3.  a  of  intergranular  yet possess e x c e l l e n t r e s i s t a n c e  to  Another p o i n t t h a t c o n f l i c t s with Gruhl's  appears to be no  and  electrochemical,  i s determined i n the second stage by  f i s s u r e s i n some environments and  i s that there  stage  However, i t f a i l s to e x p l a i n the behaviour  some aluminum a l l o y s t h a t can  theory  as a two  c o r r e l a t i o n between s t r e s s  f r a c t u r e toughness  (12).  Mechanisms I n v o l v i n g P r e c i p i t a t e D i s t r i b u t i o n and P r e c i p i t a t e Free Zone. In c e r t a i n models, the p r e c i p i t a t e - f r e e zone a d j a c e n t ' t o  g r a i n boundary i s b e l i e v e d t o be r e s p o n s i b l e corrosion.  T h i s model i s espoused by  et a l (16)  and Thomas (17) .  the  for intergranular stress  Pugh and  Jones  (13),  McEvily  These workers b e l i e v e t h a t the p r e c i p i t a t e  f r e e zone undergoes p r e f e r e n t i a l p l a s t i c d e f o r m a t i o n on a p p l i c a t i o n o f stress.  As  a r e s u l t the deforming metal w i t h i n the PFZ  t o the r e s t of the g r a i n .  Stress  c o r r o s i o n c r a c k i n g proceeds by  e n t i a l d i s s o l u t i o n o f metal w i t h i n the PFZ suggest t h a t  becomes anodic  (17).  Pugh and  Jones  prefer(13)  l o c a l i z e d p l a s t i c d e f o r m a t i o n i n i t i a t e s a g r a i n boundary  crack which propagates m e c h a n i c a l l y  through the  However, s t u d i e s conducted by H o l l (18) s m a l l a d d i t i o n s o f copper  (0.5%), s i l i c o n and  the p r e c i p i t a t e f r e e zone i n most cases.  PFZ. on s i m i l a r a l l o y s w i t h  i r o n (0.3%) d i d not  It i s s i g n i f i c a n t that  show the  - 6two most susceptible conditions were free of denuded zones but contained Guinier  Preston (G.P.) zones or coherent p r e c i p i t a t e s , whereas the  two resistant conditions did show very narrow (500 A) denuded zones. Dislocations i n highly susceptible condition formed uniformly distributed tangles. Gruhl and Cordier  (19) obtained r e s u l t s on an Al-5Zn-3Mg  (plus 1% other elements) which indicated that stress corrosion was caused mainly by metastable, redissolvable precipitates such as coherent n' phase and improved markedly when these precipitates were replaced by the stable MgZn phase. 2  Jacobs (20) proposes that the i n i t i a t i o n of a  stress corrosion crack can take place at the interface between a p i t t e d p r e c i p i t a t e p a r t i c l e and the aluminum matrix.  The crack propagates  intergranularly v i a a series of corrosion-mechanical steps.  Both  precipitate and aluminum matrix are believed to alternate as anodic phases.  In a recent study H e l f r i c h (21) demonstrated that stress  corrosion cracking does i n i t i a t e at the p a r t i c l e matrix interface i n an Al-2.6Mg-4Zn a l l o y tested i n 6% NaCl solution. The p r e c i p i t a t e free zone hypothesis i s also opposed by , Speidel  (22).  He observed long straight narrow bands of high d i s l o c a t i o n  density, with the dislocations p i l e d up against the grain boundaries in materials of high s u s c e p t i b i l i t y .  In materials aged to a reduced  s u s c e p t i b i l i t y , the s l i p bands contain dislocations of i r r e g u l a r curvature  and d i s l o c a t i o n loops.  Speidel concluded that s u s c e p t i b i l i t y  to stress corrosion cracking i n high strength aluminum a l l o y s i s enchanced by precipitates which are sheared during p l a s t i c deformation  - 7-  such as G.P. zones and intermediate p r e c i p i t a t e s .  The shearing  results  in an increase i n s l i p step height and stress concentrations at the grain boundaries as a r e s u l t of p i l e ups of unpinned dislocations adjacent to the grain boundary.  1.4.  Stress Concentration Hypotheses According to Robertson and Tetelman (8) suitably oriented  dislocation p i l e ups adjacent to grain boundary produce t e n s i l e stress components normal to the grain boundary.  They derive t h e i r results  from the analysis by Stroh (9) who concludes that the maximum normal stress occurs across a plane at 70° to the s l i p plane and diminishes with distance from the b a r r i e r .  O  =  T  m where  e  (V  /2  r  "  (1) v  '  a = maximum normal stress m x  e  = e f f e c t i v e shear stress  L = length of an array of n dislocations r  = distance from the b a r r i e r at which normal stress is calculated  Stroh has put this expression into the G r i f f i t h c r i t e r i o n for fracture leading to the result that a crack w i l l nucleate when  nT  g  >  where G i s the r i g i d i t y modulus.  1.5G  (2)  - 8 -  From e q u a t i o n s t r e s s due  (1) and  (2) a crack w i l l be  to a p i l e up  o f n d i s l o c a t i o n s i s such t h a t  . a / f > 1.5 m L  n  Uhlig  (23)  and Johnson  (24)  G  s t r a i n e d atomic bonds at the Both the shear modulus and  (3) v  et a l b e l i e v e t h a t the a d s o r p t i o n  c o r r o s i v e environment or i t s c o n s t i t u e n t s  by  formed when the maximum t e n s i l e  crack  the  lowers the s t r e n g t h  of of  t i p (stress-sorption cracking).  s u r f a c e energy are thought to be  stress concentration  s u r f a c e i s l i k e l y to be  reduced  Haynie and  Boyd  the g r a i n b o u n d a r i e s .  at a g r a i n boundary on a f r e e  l e s s than i n the  atoms are under a s m a l l e r  i n t e r i o r s i n c e the  surrounding grains however, does not cathodic  and  surface  constraint. (25)  suggest t h a t hydrogen c o n c e n t r a t e s  Under l o c a l i z e d high  thus c r e a t e s u s c e p t i b l e p a t h s .  account f o r the p r e v e n t i o n  to  the  T h e i r model,  of stress corrosion  cracking  charging.  In a r e c e n t  paper Pugh et a l (26)  stress corrosion s u s c e p t i b i l i t y  r e i t r a t e the i d e a that  T h i s seems to c o n t r a d i c t the r e s u l t s o f Beck and  o b t a i n a p o s i t i v e c o r r e l a t i o n between the  corrosion s u s c e p t i b i l i t y  the  of Al-Mg-Zn depends on the width o f  p r e c i p i t a t e f r e e zone, i n c r e a s i n g s i g n i f i c a n t l y w i t h d e c r e a s i n g width.  at  s t r e s s e s these hydrogen  r i c h areas are then assumed to become anodic with r e s p e c t  who  the  the presence o f environment at the crack t i p . The  by  J  f o r Al-7Mg.  Sperry  zone'width and  zone (27)  stress  the  - 9 -  The both by  stress  Pugh (26)  and  corrosion Jacobs  s u r f a c e i s shown to be (10).  This indicates  i s taking  p l a c e e i t h e r d u r i n g or a f t e r the  1.5.  Environmental G i l b e r t and  crack i n 3.5% resulted  Hadden (28)  solution.  p r o p a g a t i o n o f the  crack.  not  absence o f oxygen w h i l e admission o f oxygen  cracking.  During s t r e s s  corrosion  seen to  Al-7Mg would s t r e s s  of t h i s a l l o y a  evolve.  time f o r c r a c k i n g i n c r e a s e s with pH  It seems to decrease with i n c r e a s i n g  a susceptible  dissolution  showed t h a t Al-7Mg a l l o y d i d  stream o f f i n e bubbles o f hydrogen was The  that  Factors  NaCl i n the  in rapid  rough t e x t u r e d ,  i n a 3.5%  NaCl  s a l t concentration  corrode even i n a 50 ppm  NaCl  and  solution  (29). In s p i t e o f the corrosion  tremendous amount o f work done on  b e h a v i o u r o f h i g h s t r e n g t h aluminum a l l o y s and  the  a multitude  o f proposed mechanisms, a r e a s o n a b l y s a t i s f y i n g mechanism can be  s a i d to  1.6.  SCC,  The  Unresolved Points  1.  I f the do  stress 2.  result  hardly  exist.  then why  different  stress  nature o f p r e c i p i t a t e d i s t r i b u t i o n i s important  s i m i l a r p r e c i p i t a t e d i s t r i b u t i o n s r e s u l t i n a widely corrosion  Why  susceptibilities  i s i t , that  i n i d e n t i c a l stress  (12)?  d i f f e r e n t d i s l o c a t i o n arrangements  corrosion  r e l a t e s u s c e p t i b i l i t y to d i s l o c a t i o n 3.  in  characteristics?  corrode i n the  A number o f models  configurations.  While a number o f a l l o y s s u s c e p t i b l e  attack also stress  can  to  intergranular  same medium, t h e r e are  two  other  - 10 classes of aluminum a l l o y s .  One of these undergo intergranular  corrosion  but are immune t o stress corrosion while the others show no i n t e r granular attack but are highly susceptible to stress corrosion (12). 4.  The r o l e o f the p r e c i p i t a t e free zone (PFZ) i s not c l e a r .  Robinson and H o l l have observed widely d i f f e r e n t stress corrosion behaviour i n a l l o y s of s i m i l a r width of PFZ. D u c t i l e crack propagation through the PFZ i s not supported by fractographic evidence on account of an absence of dimpling  i n this  region. 5.  The r o l e of hydrogen i n stress corrosion of aluminum a l l o y s  i s not c l e a r (25). 6.  Aluminum a l l o y s can be stress corroded i n organic l i q u i d s  such as ethyl a l c o h o l , methyl alcohol and carbon t e t r a c h l o r i d e .  It i s  not known whether the mechanism of stress corrosion i n organic l i q u i d s i s the same as i n aqueous s o l u t i o n s .  1.7.  Objectives  of the Present Work  H i s t o r i c a l l y the approach to stress corrosion has been l a r g e l y phenomenological.  Most of the work has been c a r r i e d out t o evaluate  the stress corrosion s u s c e p t i b i l i t y as a function of a l l o y composition and t e s t v a r i a b l e s .  Experimental measurements of the stress  corrosion  s u s c e p t i b i l i t y have involved e i t h e r t i m e - t o - f a i l u r e f o r uncracked or . precracked specimen; or the determination of the rate of crack growth as a function of stress i n t e n s i t y f a c t o r Kj (82). There i s a s c a r c i t y of k i n e t i c data f o r i n i t i a t i o n and propagation stages as a function of temperature.  While Gruhl (73) and  - 11 -  Helfrich  (72) have s t u d i e d t h e temperature  dependence o f s t r e s s  i n aluminum a l l o y s , they have not d i s t i n g u i s h e d between and p r o p a g a t i o n times.  corrosion  initiation  T h e r e f o r e , i t i s not c l e a r whether t h e apparent  a c t i v a t i o n e n e r g i e s they observe a r e f o r t h e i n i t i a t i o n o r p r o p a g a t i o n stage.  Wei (82) has determined  the k i n e t i c s o f f a t i g u e c r a c k growth i n  a h i g h s t r e n g t h aluminum a l l o y immersed i n d i f f e r e n t apparent  a c t i v a t i o n energy  stress concentration factor 1.  The f i r s t  environments.  f o r the p r o c e s s decreases w i t h  The  increasing  K^.  o b j e c t i v e o f t h e p r e s e n t work was t o determine  the e f f e c t  o f i s o t h e r m a l ageing treatment, a l l o y c o m p o s i t i o n , s t r e s s  level,  environment  susceptibility  and specimen geometry on the s t r e s s c o r r o s i o n  o f aluminum a l l o y s . 2.  The next o b j e c t i v e was t o study the k i n e t i c s o f s t r e s s  c r a c k i n g as a f u n c t i o n o f temperature determine  t h e apparent  and environment,  a c t i v a t i o n energy  k i n e t i c observations with  corrosion  i n order t o  and t o c o r r e l a t e the  p r o c e s s e s p o s s i b l y o c c u r r i n g d u r i n g the  i n i t i a t i o n and p r o p a g a t i o n o f s t r e s s c o r r o s i o n c r a c k i n g . 3.  A t h i r d o b j e c t i v e o f t h i s work was t o c a r r y out f r a c t o g r a p h i c  s t u d i e s t o determine  the r e l a t i v e r o l e s o f d i s s o l u t i o n and mechanical  f r a c t u r e processes during s t r e s s  corrosion.  PROCEDURE  2.  2.1.  S t a r t i n g M a t e r i a l s and A l l o y P r e p a r a t i o n High p u r i t y aluminum  and h i g h p u r i t y M e l t i n g was  zinc  (99.99),  h i g h p u r i t y magnesium  (99.95) were used i n the p r e s e n t  investigation.  c a r r i e d out at 700°C i n a c y l i n d r i c a l g r a p h i t e mold,  e n c l o s e d i n a s t a i n l e s s s t e e l bomb and housed i n a v e r t i c a l heated f u r a n c e . The melt was it  (99.99),  The  s t e e l bomb was  resistance  shaken v i g o r o u s l y to s t i r the m e l t .  then s o l i d i f i e d by sharp u n i d i r e c t i o n a l c o o l i n g by  on a l a r g e s t e e l b l o c k w h i l e the upper p a r t o f the mold was  kept i n s i d e the f u r a n c e .  Pipe f o r m a t i o n was  placing still  kept t o a minimum by  t h i s technique. Four types o f a l l o y s were made.  The compositions were checked  by Coast E l d r i d g e Co. and were found t o be c l o s e to nominal. compositions are l i s t e d  The  i n Table (1). Table  (1)  CHEMICAL ANALYSIS OF FABRICATED ALLOYS Nominal Composition Al-9Mg Al-9Mg Al-3Mg-6Zn Al-3Mg-6Zn Al-22Zn Mg-9A1 Mg-9A1  (wt  %)  A c t u a l Composition  (wt  %)  Al-9Mg Al-8.6Mg Al-3Mg-6Zn Al-2.6Mg-6.3Zn Al-21.5Zn Mg-8.6A1 Mg-9.1A1  A l l the a l l o y compositions i n t h i s t h e s i s are s t a t e d i n weight p e r c e n t o f solute.  - 13 -  2.2.  F a b r i c a t i o n by r o l l i n g The c a s t i n g s were a p p r o x i m a t e l y 1" i n diameter and 6" l o n g .  They were homogenised and subsequently hot r o l l e d C o n s i d e r a b l e d i f f i c u l t y was  encountered  which showed hot s h o r t n e s s at temperatures was  therefore rolled  g i v e n between each pass. ground of  water quenched. 0.025".  speed with an  Any  edge c r a c k s formed  When the sheet reached a t h i c k n e s s  then c o l d r o l l e d  0.050" then c o l d r o l l e d  similarly to 0.025".  rolled  at 200°C to a t h i c k n e s s ;,of  The r o l l e d  a l l o y s were used f o r  p r e l i m i n a r y age h a r d e n i n g , s t r e s s c o r r o s i o n and The r o l l e d sheet was  sheared i n t o 2.5"  s t r i p s , which i n t u r n were punched i n t o t e n s i l e a reduced gauge l e n g t h o f 0.8"  2.3.  was  i n the d e s i r e d  u.  Al-3Mg-6Zn was  o f specimen  and  50% t o a f i n a l t h i c k n e s s o f  The heat treatment b e f o r e c o l d r o l l i n g r e s u l t e d  g r a i n s i z e o f 90  were  d u r i n g r o l l i n g were  s o l u t i o n heat t r e a t e d a t 430°C f o r 20 minutes I t was  It  initial  I n t e r m i t t e n t anneals o f 10 minutes  o f f b e f o r e the a n n e a l i n g .  0.05", i t was  i n r o l l i n g Al-9Mg  between 300° to 400°C.  at 150°C at a slow r o l l  r e d u c t i o n o f 10% per pass.  or e x t r u d e d .  parallel  slip line  studies.  long by 0.75"  wide  t e s t specimens h a v i n g  w i t h a width o f 0.21".  The  to the r o l l i n g d i r e c t i o n , f i g .  tensile  axis  (la).  Extrusion 3" long b i l l e t s were machined t o diameter o f 0.96"  at h i g h temperature.  The  and  extruded  f o l l o w i n g t a b l e g i v e s the e x t r u s i o n v a r i a b l e s .  The diameter o f the extruded wire was  0.152".  -  14  -  Figure ( 1 ) . Types of t e n s i l e and stress corrosion specimens, used i n the present work.  - 15 Table (2) EXTRUSION CONDITIONS OF ALLOYS Alloy  Temp. °C  Pressure Initial  Final  Extrusion ratio  Average extrusion speed  Al- 8.6Mg  390  116,000  141,000  38.5:1  5"/min  A l - 2.6Mg-6.3Zn  415  81,500  106,000  38.5:1  5"/min  A l - 22%Zn  380  45,000  60,000  38.5:1  10"/min  Mg- 8.6A1  360370  49,000  62,000  38.5:1  10"/min  Two types of specimens were made from the wire.  They are  shown i n f i g . lb and l c . The notched specimens were machined a f t e r the solution heat treatment while the t e n s i l e specimens were machined before.  This  procedure was made necessary by the fact that specimens notched p r i o r to solution heat treatment bent or developed cracks during quenching. No cracks were found i n specimens notched after solution heat treatment.  2.4.'  Heat Treatment Al-9Mg was solution treated at 450°C f o r 15 minutes i n a i r and  water quenched.  It was aged at 200°C f o r d i f f e r e n t times.  Al-3Mg-6Zn and Al-2.6Mg-6.3Zn were solution treated at 460°C for 4 hours, water quenched, then aged at 160°C. Al-21.5Zn was solution treated at 500°C for 1 hour, step, quenched to 300°C f o r 1/2 hour, water quenched and aged, at room temperature (15).  - 16 -  Figure ( 2 ) .  Stress corrosion c e l l s used i n the present work.  - 17 Mg-8.6A1 was solution treated at 420°C for 70 hours i n Argon, water quenched, then aged at 200°C.  2.5.  Polishing and Degreasing The machined and ground specimens were electropolished i n a  solution containing 45 v o l . %  n i t r i c acid, 45 v o l . %  concentrated hydrochloric acid at 15°C.  methanol and 5 v o l . %  The polishing was carried out 2  at 4 to 6 volts and 2.5 to 3 amperes (2 amp/cm ). A bright surface was obtained after electropolishing.  The  specimens were then washed i n water and u l t r a s o n i c a l l y degreased for 15 minutes i n  chlorothene.  Tensile Tests Tensile tests were done on the Instron machine.  Wedge type  grips were used for f l a t specimens ,and threaded grips f o r round specimens. 2.6.  Stress Corrosion Tests Three types of stress corrosion tests were employed to determine  the  stress corrosion s u s c e p t i b i l i t y (SCC) of a l l o y s .  Constant Load Tests A modified creep stand shown i n f i g . (3) was used.  Vise  type grips were used for f l a t specimens and threaded grips for the round specimens.  The f a i l u r e time was measured with an hourmeter which  could be switched o f f by the f a l l i n g weight. fig.  The stress corrosion c e l l  (2) consisted of a glass tube with t e f l o n or rubber bungs f i t t e d  around the specimen.  The c e l l was sealed with wax or s i l i c o n e grease.  - 18 -  Leads to hour meter  Time switch s i l k thread and spring specimen stress corrosion c e l l grip  guide plate  guide plate turn buckle  b a l l bearing support  load  \\\\\\\ Figure (3). Modified creep stand for constant load tests Tiesenhausen and Lund (103)).  (After  - 19 -  The f a i l u r e time could be measured with an accuracy of ± 1 minute on the hourmeter.  2.7.  A stopwatch was used f o r short f a i l u r e times.  Load Relaxation Tests The major part of stress corrosion testing was carried out on  notched c y l i n d r i c a l specimens i n the Instron.  The load relaxation curve  consisted of three parts, 1.  I n i t i a l load relaxation  2.  Period of constant load  3.  Period of crack propagation r e s u l t i n g i n relaxation of  the load, followed by rapid mechanical fracture. Tests of this type are an improvement over the constant load test f o r several reasons. 1.  The crack can be l o c a l i z e d near the root of the notch.  2.  I n i t i a t i o n and propagation times can be determined with  f a i r accuracy. 3.  Very high stresses can be applied to the notched specimen  because of the t r i a x i a l stress state existing at the root of the notch. As a r e s u l t , the stress corrosion f a i l u r e times can be shortened.  2.8.  Strain Rate Tests Tensile tests were carried out on a quenched and averaged  Al-2.6Mg-6.3Zn, i n deionized water, 3.5% NaCl + 2% K C r 0 2  gel.  4  and s i l i c a  The effect of s t r a i n rate on t o t a l elongation and uniform elonga-  tion was  determined.  - 20 2.9.  Hydrogen Charging Experiments These experiments were conducted to determine whether hydrogen  played any aprt i n the stress corrosion process.  Using 3.5% NaCl +  2% K^CrO^ as the environment, the notched specimen was made the cathode and a platinum wire loop placed around i t , acted as the anode.  A  potential of 4.5 volts was applied for 30 minutes and vigorous hydrogen evolution took place at the specimen surface.  The charging was  continued u n t i l the specimen had been subjected to the required applied stress.  Then the current was switched o f f , so as to prevent the  application of a cathodically protective p o t e n t i a l . A dark f i l m was formed on the specimen surface.  The  p o s s i b i l i t y that such a f i l m may be contributing to stress corrosion was checked by hydrogen charging two specimens for t h i r t y minutes, and holding them for 8 days at 0°C to allow the hydrogen to escape from the notch.  These specimens were retested i n 3.5% NaCl + 2.0%  K^CTO^.  Charging tests were conducted on Al-21.5Zn and Al-2.6Mg^-6.3Zn.  2.10.  Optical Microscopy Optical microscopy was used for grain size determination and  for studying whether crack propagation was i n t e r c r y s t a l l i n e or transcrystalline. A l l fracture surfaces were f i r s t examined under an o p t i c a l microscope at low magnification before a r e p l i c a was taken.  - 21 -  2.11.  Electron Microscopy 1.  Precipitate d i s t r i b u t i o n for d i f f e r e n t aging,  treatments  was studied to a limited extent, using thin f i l m and r e p l i c a techniques. 2.  S l i p line d i s t r i b u t i o n was studied i n peak aged Al-9Mg.  3.  Extensive fractography was carried out on t e n s i l e and  stress corroded specimens.  The extent of the stress corroded area  could be controlled by varying the i n i t i a l stress level i n the load relaxation test.  From the nature of the load relaxation curve i t was  possible to determine the extent of stress corrosion compared to mechanical  fracture. Low magnification (300x) electron probe pictures of the  fracture surface were used to locate areas of the specimen being observed on the r e p l i c a .  3.  RESULTS AND OBSERVATIONS  The isothermal p r e c i p i t a t i o n hardening curves f o r Al-9Mg, Al-3Mg-6Zn and Mg-8.6A1 are shown i n figures (4-6).  The f i r s t  two  alloys were tested i n the r o l l e d condition and the last i n the extruded form.  3.1.  Al-9Mg The a l l o y was solution heat treated at 450°C for 15 minutes i n  a i r , water quenched and aged at 200°C.  The grain size was 90 u .  In figure (4) the 0.2% y i e l d stress i s plotted against the aging  time.  7 hours. aging  The peak y i e l d strength i s reached at an a^ing  time of  The uniform p l a s t i c elongation decreases with increasing  times attaining a minimum value i n seven hours.  3.2.  Al-3Mg-6Zn The r o l l e d alloy was solution heat treated at 460°C for 4 hours  i n a i r , water-quenched and aged at 160°C f o r times ranging from a few hours to 120 hours. treatment.  The specimens were water quenched after the aging  From figure (5) i t i s seen that the maximum y i e l d stress i s  reached i n times between 8 and 10 hours at 160°C.  The uniform elongation  reaches a minimum at the same time and increases again on overaging.  2  5 - 1 0  20 Aging Time Hours  50  Figure (5). Precipitation hardening of Al-3Mg-6Zn  3  100 aged at 160°C.  ^w^-——  1  5  1  1  1  10° .  2 5 Aging Time Hours  1  1  10  1  Figure (6). Precipitation hardening of Mg-8.6A1,'aged at 200°C.  1  2  1  5  10  2  - 26 3.3.  Mg-8.5A1 C y l i n d r i c a l t e n s i l e specimens were solution treated at 420°C  for 70 hours, water quenched and aged at 200°C. 0.2% y i e l d stress i s plotted against ageing time. a maximum value after aging  In figure (6) the The strength reaches  for 24 hours and stays constant up to 50  hours a f t e r which i t starts to f a l l o f f .  The uniform elongation decreases  to a minimum value i n 16 hours and then r i s e s on further aging.  3.4.  Stress Corrosion Testing Techniques Both constant load and load relaxation type of tests were  employed. Constant  load tests are used to evaluate the stress corrosion  s u s c e p t i b i l i t y at different stress levels and d i f f e r e n t heat  treatments.  The drawback of these types of tests i s that once a crack i s i n i t i a t e d , the stress increases at a very rapid rate and the area of stress corrosion crack propagation i s therefore much smaller than the area of mechanical fracture due to overload. ative fractographs  This makes i t very d i f f i c u l t to get represent-  of the stress corroded area.  In addition i t i s  impossible to distinguish between crack i n i t i a t i o n and crack propagation stages.  Another drawback of constant load tests i s that f a i l u r e  may  occur on account of general corrosion more e a s i l y than i n load relaxation tests.  However, since most specimens tested  f a i l e d i n times less than  a day, general corrosion was not a problem. The drawbacks of the constant load tests were overcome by using load relaxation tests on notched specimens. stress to be l o c a l i s e d so that a single crack was  The notch helped the initiated.  The extension  Table (3) SUMMARY OF STRESS CORROSION TESTS AND Type of A l l o y  Type of Test  Al-9Mg Strip  Constant Loading  TESTING CONDITIONS  Purpose of Study  Environment  1. Aging time vs f a i l u r e time at constant stress  3.5% NaCl + 2.0% K C r 0  20  3.5% NaCl + 2.0% K C r 0  20  3.5% NaCl + 2.0% K C r 0  20  2  Temperature  4  2. Aging time vs f a i l u r e time at. constant fraction of y i e l d stress 3. Applied stress v i a f a i l ure time at peak strength Al-9Mg Round notched  Load relaxation test  1. Crack i n i t i a t i o n and propagation times  2  4  2. Nature of load propagat i o n times 3. Activation energy analysis on peak aged specimens Al-3Mg-6Zn Strip  Constant load  1. Applied stress vs f a i l u r e time at peak strength  2  4  2. Aging time vs f a i l u r e time at constant stress 3. Aging time vs f a i l u r e time at constant fraction of y i e l d stress Al-2.6Mg-6.32 Round notched  Load relaxation test  1. I n i t i a t i o n and propagat i o n times  3.5% NaCl + 2.0% K C r 0  2. Nature of load relaxation  Ambient a i r ,  2  4  •  20  °C  Table (3) (Continued) Type o f A l l o y  Type of T e s t  A1.2.6Mg-6.32 (continued)  Al-2.6Mg-6.32Zn Round, t e n s i l e  Tensile  tests  Al-21.5Zn Round, notched  Load r e l a x a t i o n test  Purpose o f Study  Environment  3. E f f e c t o f d i f f e r e n t environments on the l o a d r e l a x a t i o n curve  dry a i r , deionized water, ethanol, methanol  4. A c t i v a t i o n energy a n a l y s i s on overaged specimens  3.5% NaCl + 2.0% K C r 0 , 99.9%  S t r a i n r a t e vs & elongation i n d i f f e r e n t media, f o r as quenched and overaged c o n d i t i o n s  D. w a t e r , 3.5% NaCl + 2.9% K C r 0 , s i l i c a gel  1. Crack i n i t i a t i o n and propagat i o n times a t c o n s t a n t s t r e s s i n d i f f e r e n t media  D. w a t e r , s i l i c a g e l Na d r i e d Kerosene, Mg(0Cl ) , e t h a n o l  2  4  2  3. E f f e c t o f v a r i a t i o n o f s t r e s s 99.9% C H OH on i n i t a t i o n and p r o p a g a t i o n t i m e s 2  Round, notched  Load r e l a x a t i o n test  load  0-85°C 0-45°C 20°C  -20 -30 -65 -70  to to to to  4. Hydrogen c h a r g i n g experiments 3.0% NaCl + 2 . 0 % E f f e c t o f s t a i n r a t e and e n v i r o n - K C r 0 , D.water, ment on m e c h a n i c a l p r o p e r t i e s . silica gel.  20°C 20°C  Aging time v s f a i l u r e time a t a 3.5% NaCl + 2.0% constant f r a c t i o n o f y i e l d s t r e s s K C r 0  20°C  Crack i n i t i a t i o n and p r o p a g a t i o n 3.5% NaCl + 2 . 0 % times K Cr0  20°C  2  Constant  OH  4  2. A c t i v a t i o n energy a n a l y s i s a t 95% C^H^OH different stress levels.  Mg-8.5A1 Round, unnotched  Temperature °C  2  2  4  4  4  20°C 20°C 20°C 20°C  - 29 caused by this crack was registered on the Instron, as a load drop.  In  the absence of a stress r a i s i n g notch, several cracks can be i n i t i a t e d in the specimen.  This can result i n a 'spring l i k e ' accommodtion of the  extension with no v i s i b l e load drop at the moment of i n i t i a t i o n .  Other  advantages of these types of tests are:  1.  The applied load and the time to f a i l u r e can be measured  very accurately on the Instron. 2. measured.  Both crack i n i t i a t i o n and crack propagation times can be  It i s possible to know whether load relaxation i s continuous  or discontinuous. 3.  Very high stresses can be applied at the root of the notch.  In general, the notch t e n s i l e strength i s at least two times the ultimate strength of the smooth specimen.  The times of stress corrosion  i n i t i a t i o n can therefore be shortened d r a s t i c a l l y anywhere from a factor of 3 to 10. 4.  Because of the high t r i a x i a l stress state at the root of  the notch i t i s possible to use very d i l u t e environments and s t i l l obtain f a i l u r e s i n short times. 5.  The stress corrosion test can be stopped and the notched  specimen sectioned longitudinally to observe the mode of crack propagation. 6.  A notched specimen approaches the service  conditions better  than an unnotched specimen because i t takes into account the effect of stress raisers such as screwthreads, sharp f i l l e t s , fatigue and weld cracks and p i t s developed during exposure to a corrosive  environment.  For the above reasons, notched wires were used i n a l l four systems studied.  - 30 Tensile tests at d i f f e r e n t s t r a i n rates were conducted to evaluate the e f f e c t of environment on mechanical properties, p a r t i c u l a r l y ductility.  3.4.1.  Stress Corrosion S u s c e p t i b i l i t y S u s c e p t i b i l i t y to stress corrosion was defined as the inverse  of t o t a l f a i l u r e time i n the constant  load test and the inverse of crack  i n i t i a t i o n time i n the load relaxation t e s t .  The s u s c e p t i b i l i t i e s of  two specimens having d i f f e r e n t y i e l d strengths can be compared either at a constant  f r a c t i o n of the y i e l d stress or at a constant  methods were used.  stress.  Both  From figures (7) and (9) i t appears that both methods  give s i m i l a r variations i n f a i l u r e time vs aging  time curve.  A common  environment namely 3.5 wt.% NaCl + 2.0 wt.% K^CrO^ solution i n deionized water was prepared from reagent grade chemicals. This solution f a c i l i t a t e d an evaluation of the r e l a t i v e s u s c e p t i b i l i t i e s of d i f f e r e n t a l l o y systems.'  The f a i l u r e times were short compared to those i n  alternate immersion tests i n NaCl/H^O^ solutions.  They varied from a  few seconds i n Al-9Mg and A121.5Zn to a few hundred hours i n Al-3Mg-6Zn. If a specimen did not f a i l i n 200 hours the test was discontinued.  Certain  specimens were tested for 600 hours before discontinuing the t e s t .  3.4.2.  Constant Load Tests  3.4.2.1.  Al-9Mg Strip specimens of 2 3/4 inches length and 0.8" gage length  were tested by t o t a l immersion i n the NaCl/K^CrO^ solution under a constant  load.  - 31 -  In figure (7) the f a i l u r e time i s plotted against for isothermal aging at 200°C.  the time  It i s found that the solution treated  and water quenched' specimens are immune to stress corrosion. proceeds, the stress corrosion s u s c e p t i b i l i t y  increases  peak hardness the f a i l u r e time i s only 15 seconds. an improvement i n the stress corrosion  As aging  rapidly.  At  Overaging leads to  resistance.  The minimum i n f a i l u r e time occurs close to 7 hours aging at both stress levels used, namely 20,000 p s i and 40% of the respective y i e l d strengths. There i s a f i v e orders of magnitude decrease i n f a i l u r e time, produced by peak-aging the solution treated  specimen.  In figure (8) the applied stress i s plotted against time on a log scale.  failure  It i s obvious that the f a i l u r e time t i s not  related to the applied stress a by an exponential relationship of the form t  f  i  = A e  i  as suggested by some workers is approached.  (73) . As the threshold  stress of 5800 p s i  The f a i l u r e time changes at a remarkable rate.  At high stress levels the f a i l u r e times did not fluctuate by more than a factor of two and at lower stress levels variations of up to a factor of 4 were observed.  For every stress l e v e l , specimens were  tested i n duplicate and i n some cases i n t r i p l i c a t e . The purpose of the constant load tests was to a r r i v e ati the most susceptible heat treatment.  This heat treatment was subsequently  used i n the stress corrosion tests on the notched specimens.  - 32 -  Aging Time Hours Figure (7). Failure time vs aging time f o r Al-9Mg aged at 200°C.  - 34 3.4.2.2.  Al-3Mg-6Zn In figure (9) fracture time i s plotted against  at a constant applied stress of 22,100 p s i .  aging time  The curve shows a minima  at an aging time of 3 hours at 160°C which corresponds to an underaged condition. When the stress level i s changed to 40% of the 0.2% offset y i e l d stress f o r that p a r t i c u l a r condition and the fracture time plotted against  the aging time, the curve once again has a minima for an aging  time of 3 hours at 160°C. In this system the difference i n cracking as-quenched and peak-hardened conditions  times, between the  i s i n s i g n i f i c a n t compared to  the vast differences between the two conditions  i n Al-9Mg.  In f i g ; (10) the applied stress i s plotted against time for the peak-aged condition of 8 hours at 160°C.  The curve drops  rapidly at high stresses, and then more slowly as the applied i s decreased.  fracture  stress  A s i m i l a r trend i s shown by the overaged specimens.  Once again the points do not l i e on a straight l i n e suggesting that there i s no simple exponential r e l a t i o n s h i p between the rate of , cracking and the applied stress as suggested by Gruhl (73) . H e l f r i c h (72) observed a curve at lower stress levels and a straight l i n e at high stresses.  3.4.2.3. Mg-8.6A1 This a l l o y was used i n the form of a wire because i t was not feasible to r o l l i t . C y l i n d r i c a l t e n s i l e specimens of t h i s a l l o y were stress corroded at a constant load.  The applies stress was 70% of the  - 35 -  Aging time, hours  Figure (9).  Effect of aging time at 160°C on the f a i l u r e time of < Al-3Mg-6Zn i n NaCl/K CrO  100  2  Figure (10)..  5  10  1 2 5 Failure time, Hours  O  Aged 160°C, 8 hours  ©  Aged 160°C, 24 hours  102  Applied stress vs f a i l u r e time for the.aged Al-3Mg-6Zn a l l o y .  - 37 -  14  O  12  I  10J  2  5  —I 10  I 1  2  I  I  I  5 IO 2 Failure Time, Min. 2  I  I  5  10  I 3  2  I •  5  •Figure (12). Effect of applied stress on the f a i l u r e time-of•solution treated and quenched Mg-8,6A1.  - 39 -  y i e l d stress for a given aging time.  The r e s u l t i s shown i n figure  in which the f a i l u r e time i s plotted against the isothermal at 200°C.  (11)  aging time  The unaged a l l o y i s the least susceptible and aging r e s u l t s  in a progressive  increase i n the stress corrosion s u s c e p t i b i l i t y .  Unlike  the previous two systems there i s no observed minimum i n the f a i l u r e time before or at the peak strength.  The minimum f a i l u r e time i s observed  for an aging time of 76 hours when considerable  overaging has  occurred.  The applied stress vs f a i l u r e time i s plotted i n f i g . (12) and the points approximately l i e about a straight l i n e . stress corrosion cracking i s transgranular  The mode of  i n the Mg-8.6A1 a l l o y .  S u s c e p t i b i l i t y tests as a function of aging time were not conducted on Al-22Zn, because i t ages at room temperature and i n such conditions i s very susceptible to stress corrosion  3.4.3.  Strain Rate Tests  3.4.3.1.  Al-2.6Mg-6.3Zn  cracking.  Strain rate tests were carried out over 6 d i f f e r e n t s t r a i n rates i n two d i f f e r e n t media consisting of deionized water and 3.5% 2.0%  NaCl +  K^CrO^ for both the as-quenched condition and the overaged  conditions.  Specimens were also deformed at the lowest s t r a i n rate i n  dessicant dried a i r . The r e s u l t s are shown i n f i g . (13) and f i g . (14).  In both  cases there i s a s i g n i f i c a n t decrease i n uniform elongation with a decreasing s t r a i n rate in both del onized water and NaCl/K^CrO^ solution. Changing the environment to dessicant dried a i r leads to a seven f o l d increase i n d u c t i l i t y at the lowest s t r a i n rate i n figure (13).  ;  Strain rate sec" Figure (13). Strain rate vs uniform and t o t a l elongation f o r solution treated and quenched Al-2.6Mg-6.3Zn Alloy.  Strain rate, sec Strain rate vs uniform and t o t a l elongations f o r an overaged Al-2.6Mg-6.3Zn a l  3.4.3.2.  Al-21.5Zn Table (4)  EFFECT OF STRAIN RATE ON DUCTILITY AND YIELD STRENGTH IN DRY AIR AND DEIONIZED WATER Strain rate  6.0xl0 /sec _6  II  it  II  3  2.7x15 /sec II  3.4.4.  Environment  0.05 psi  0.2 psi  UTS psi  Eln Max  Eln Fract.  %  C  0  s i l i c a gel  38,000  42,300  44,000  0,.9  3 .1  s i l i c a gel  38,600  41,400  43,200  0,.8  2 .5  deionized water  14,250  --  14,250  0,.062  0 .74  deionized water  14,250  14,250  0..045  1 .1  s i l i c a gel  39,400  42,200  49,200  3,.23  5 .0  deionized water  38,500  41,000  43,800  1..17  3 .14  Load Relaxation Tests for Different Alloy-Environment Systems. The crack i n i t i a t i o n and propagation times as determined  load relaxation tests are l i s t e d i n Tables (5-9) and figures  from  (15-16).  - 43 -  Table (5)  CRACK INITIATION AND PROPAGATION TIMES FOR Al-3Mg-6Zn IN 3.5% NaCl + 2% K C r 0 2  4  Applied stress psi  In. time min.  Prop. time min.  Total time min.  460°, 4 hours water quenched, N.T.S.=74,500 + 2000  48,000  208  28  236  50,000  46  49  95  460°, 4 hours, water quenched, 160°C, 24 hours N.T.S.=124,000 ± 5000  55,000  117 194  6 5  123 199  69,000 69,000 .  50 93  2 4  52 97  46,000 46,000  7 10  16 14  23 14  Heat treatment  460°C, 22 hours water quenched, N.T.S.=62,000  Nature of load relaxation min. discontinuous, down to zero load  discontinuous, followed by rapid mechanical fracture j  discontinous, followed by rapid mechanical fracture  N.T.S. = Notch t e n s i l e strength i n p s i .  Grain size The grain size of Al-3Mg-6Zn solution treated for 4 hours at 460°C was 265 microns, whilefor times of 22 hours i t was 570 microns. The grain size for Al-2.6Mg-6.3Zn solution treated for 4 hours at 460°C was 300 microns. grains across the notch.  In other words, there were approximately s i x  - 44 Table (6) CRACK INITIATION AND PROPAGATION TIMES FOR Al-2.6Mg-6.3Zn IN DIFFERENT ENVIRONMENTS Heat treatment  Applied Stress psi  460°C, 4 hours,  57,600  1.6  0..6  2,.2  3.5NaCl + 2.0K CrO  4  water quenched  57,600  1.0  ' 0..5  1..5  3.5NaCl + 2.0K CrO  4  57,600  1.5  0..5  2 .5  Deionized  N.T.S.=72,000  Medium  2  2  water  3..5  13,.5  Ambient a i r  57,600  11  7..5  18,.5  Ambient a i r  57,600  20  6  26..0  Ambient a i r  57,600  940  110  60,000  23  60,000  94.2  127.,4  221,.6  CC1  4  + D. water  68,000  36.6  81..4  118..0  CC1  4  + D. water  68,000  40  71  110  CC1  83,000  Methyl alcohol  34,.10 Methyl alcohol  No f a i l u r e in 26 hours CC1. + Dry a i r (1560' ) 14  480  494  Ambient a i r  20  270  290  Ambient a i r  37.4 89  83,000  11..1  53..4  S i l i c a gel  44.2  83,000  9.,2  1050  60,000  water quenched N.T.S.=104,000  Total time min.  10  water quenched 160°C, 24 hours  Prop. time min.  57,600  68,000 460°C, 4 hours  In. time min.  5.,6  43  Deionized  water  16.,0  105  Deionized  water  No f a i l u r e in 56 hours; S i l i c a gel  I n i t i a t i o n time  No  vssssy P r o p a g a t i o n time  failure  83,000  "0 CD Ml  i i l i c a gel 83,000  83,000  33"  JZ2  rt' ^ '  CD >  D e i o n i z e d water  o  Ambient a i r  SSSSA  17777.  68,000  ft C •H C/) in CD  u  T3 CD •H  CC1  68,000 68,000 68,000  177?  zzz  + Dry a i r  4  CC1 CC1. + D. water  ZL  4  '/Y/yYyY//vxV7/77//\  60,000  Methyl a l c o h o l  Cn  ~V>y S i l i c a g e l  57,600  r—I  -a  PL, PH  •< •57,600 57,600  CD +-> rt CD U H  Ambient a i r  o  D e i o n i z e d water  vyyy**  •H  3  »—i  57,600  3.5% NaCl + 2.0%  Yzza 10 F i g u r e (15).  20  -L  ±  50 100 200 F a i l u r e time, Minutes  X  500  1000  K Cr0 2  O OO  4  JL 2000  5000  E f f e c t o f d i f f e r e n t media on crack i n i t i a t i o n and p r o p a g a t i o n times i n Al-2.6Mg-6.3Zn a l l o y i n s o l u t i o n t r e a t e d and overaged c o n d i t i o n s .  - 46 -  T a b l e (7)  CRACK INITIATION AND PROPAGATION TIMES FOR Al-8.6Mg i n 3.5% NaCl + 2% K C r 0 2  4  Applied stress psi  In. time min.  Prop. time min.  12,500  18.4  1  31,200  8.5  450°C, 30 min., water quenched 200 C, 3 hours, water quenched N.T.S.=71,000  40,800  0.16  20  38,200  0.75  131.2  132.0  450°C, 22 hours water quenched 200°C, 7 hours N.T.S.=48,000  31,000  3  114 .  117  Heat treatment  450°C, 30 min., water quenched 200°C, 7 hours, water quenched N.'T.S.=62,000  9  ..The g r a i n  0.5  s i z e was 90 y f o r the f i r s t  for the t h i r d .  Total time min. 19.4 9.0  20.16  Nature o f load relaxation  Discontinuous, f o l l o w e d by r a p i d fracture  Discontinuous, f o l l o w e d by r a p i d fracture  slow crack propaga t i o n t o zero load  two heat treatments and 250 y  Table (8) CRACK INITIATION AND PROPAGATION TIMES FOR Mg-8.6A1, IN NaCl/K Cr0 2  4  SOLUTION  Heat treatment  420°C, 70 hours water quenched N.T.S.=32,000  Applied stress psi  In. . Prop. time time min. min.  26,000  0.4  1.6  Total time min. 2.0  Type of solution  3.5%NaCl + 2.0% K Cr0 2  26,000  0.2  1.2  3.5%NaCl + 2.0% K Cr0 2  26,000  1.4  11.0  11.4  0.7  10.0  26,000 420°C, 70 hours water quenched  30,000  0.8  1.0  0.035%NaCl + 0.02% K C r 0 no f a i l u r e 2  0.2  0.9  1.1  30,000  0.7  2.0  2.7 70.5 hrs. 0+  4  0.35%NaCl + 2.0% K Cr0 2  26,000  4  3.5%NaCl + 2.0% K Cr0 2  200°C, 1 hour water quenched N.T.S.=36,000  4  3.5%NaCl + 2.0% K Cr0 2  30,000  4  10.7 95 hrs, 0+  0.2  4  0.35%NaCl + 0.2% K C r 0 no steps 2  26,000  4  4  0.35%NaCl + 2.0% 4 failure K  C r 0  n  o  2  The grain size was 120 \i f o r both the solution treated and the aged alloys.  - 48 -  Al-21.5Zn Table (9) EFFECT OF DIFFERENT MEDIA ON THE INITIATION AND PROPAGATION TIMES FOR Al-21.5Zn ALLOY  Heat treatment  500°C, 1 hour, step quenched to 300°C, 1/2 hour water quenched Aged at 20°C 1.5 hours N.T.S.=81,000 (20°C) N.T.S.=84,000 (-60°'C)  Applied stress psi  In. time min.  64,600  --  Total time min.  Medium  0,.2  0,.2  Deionized wate  --  0,.3  0..3  II  it  --  0,.1  0,.1  II  it  6..2  3,.0  9..2  3,.8  1,.0  4,.8  12,.5  1,.5  14  Mg(0Cl )  14,.2  1,.5  15..7  Kerosene + Na  9  3  12..0  Mg(OCl )  0..4  0,.35  0..75  0..5  0,.35  0,.85  0..6  0,.2  0..8  Ambient A i r  1,.6  0,.9  2,.5  Ambient a i r 20  3..8  1,.4  5..2  Ambient a i r 20  1..4  0,.2  1..6  Ethyl alcohol  40,000  0..2  1..9  2..1  3.5NaCl + 2.OK  26,000  0.,5  2,.5  3..0  n  it  0.,7  1..3  2  n  tt  10,000  The grain size was 105 microns.  Prop. time min.  126  s i l i c a gel II  n  4  4  2  2  IMNaOH M  YSSJ\  I n i t i a t i o n time Propagation time  :ezzUMWAM>'ZZA  26,000 p s i  3.5% NaCl + 2.0% K C r 0 2  4  "SL Kerosene dried i n Na  nzzP  CM  O o  Mg(0Cl ) 4  2  vO  32T  T3  S i l i c a gel  T7772  ID  Ambient a i r 20 y  Y///A  •H  r—I  <  Ambient a i r 760 mm 1 M NaOH  SSL  Deionized water 0.1 Figure (16).  X  1.0  1  10 Failure time, Min.  1  100  Effect of different media on i n i t i a t i o n and propagation times i n Al-21.5Zn alloy.  - 50 -  3.5.  The pH o f t h e N a C l / K C r 0 2  The  4  Solution  pH o f t h e s o l u t i o n was 8.6 a t room temperature and d i d not  change s i g n i f i c a n t l y d u r i n g  the t e s t .  Mg-8.6A1 were a d j u s t e d t o t h i s pH.  The d i l u t e s o l u t i o n s used i n  D i l u t e s o l u t i o n s were used i n t h i s  system, because f a i l u r e times i n t h e normal s o l u t i o n were extremely short.  The f a i l u r e times could  l e v e l s , because t h e t h r e s h o l d  3.6.  A c t i v a t i o n Energy The  activated.  not be p r o l o n g e d by u s i n g  s t r e s s was h i g h  (72% y i e l d s t r e s s ) .  Analysis  s t r e s s c o r r o s i o n p r o c e s s i s both s t r e s s and t h e r m a l l y  Data i s s c a r c e l y a v a i l a b l e on t h e k i n e t i c s o f t h e s t r e s s  c o r r o s i o n p r o c e s s i n aluminum a l l o y w s . no  low s t r e s s  Of t h e d a t a t h a t  i s available,  d i s t i n c t i o n i s made between t h e i n i t i a t i o n and p r o p a g a t i o n times. T h i s study was undertaken with a view t o determine t h e  apparent a c t i v a t i o n energy o f t h e r a t e c o n t r o l l i n g p r o c e s s i n the s t r e s s corrosion  o f aluminum a l l o y s .  The e f f e c t o f temperature on t h e r a t e  o f i n i t i a t i o n and p r o p a g a t i o n was s t u d i e d ethyl alcohol. 1. rates  This  environment was s e l e c t e d  I t was p o s s i b l e  o f extremely s u s c e p t i b l e  temperatures.  to overage d u r i n g  was  found t h a t  f o r two r e a s o n s .  t o study the c r a c k i n i t i a t i o n and p r o p a g a t i o n a l l o y s , such as Al-21.5Zn a t sub-zero  The f a i l u r e times i n t h e s a l t  f o r a meaningful a n a l y s i s .  2.  i n a common environment namely  s o l u t i o n were t o o s h o r t  Low temperatures a l s o reduced t h e tendency  t h e t e s t and extended the range o f t h e A r r h e n i u s p l o t .  With e t h y l a l c o h o l , h i g h s t r e s s the s e p a r a t i o n  igore a c c u r a t e a t h i g h s t r e s s e s  l e v e l s could  be used.  It  o f i n i t i a t i o n and p r o p a g a t i o n times was than a t low s t r e s s e s .  - 51 -  Generally the i n i t i a t i o n times varied from a few minutes to a maximum of 66 hours.  Temperature was regulated within +1°C by  periodic addition of l i q u i d nitrogen to the ethyl alcohol bath around the c e l l . For the Al-Mg-Zn a l l o y , d i s t i l l e d water, and NaCl/K CrC> 2  solutions were also used i n the temperature  range of 0-85°C.  duplicate specimens were tested at extreme temperatures specimens at intermediate  4  Generally  and single  temperatures.  Before testing, the specimens were electropolished, degreased in chloroethane and c a r e f u l l y examined for any pre-existing cracks i n the notch.  According to the Arrhenius rate law  Rate = A e  RT  (1)  where A = Frequency factor Q = Apparent  activation energy of the rate c o n t r o l l i n g process.  In the present case, rates of cracking were defined as the inverse of i n i t i a t i o n and propagation times i n seconds.  The logarithm  of the rate was plotted against the reciprocal of absolute  temperature.  It i s assumed that the crack i n i t i a t i o n and propagation distances remain constant at a given stress.  Therefore the rates of  i n i t i a t i o n and propagation can be defined as the inverse of i n i t i a t i o n and propagation times, respectively.  However, as i n i t i a l a p p l i e d s t r e s s  decreases, the crack propagation distance increases and the above assumption  i s no longer v a l i d .  ;  - 52 -  In a l l cases, the Arrhenius plots were straight l i n e s .  The  apparent a c t i v a t i o n energy was calculated from the slopes of the l i n e s . The lines were the best f i t s as calculated from regression In addition,the  analysis.  error i n the slope f o r a 95% confidence l i m i t was  calculated f o r each set of data. The results are described  3.6.1.  i n more d e t a i l below.  A1-21.5ZH This a l l o y was exposed to 99.9% ethyl alcohol as well as 95%  ethyl alcohol, the remainder was water.  The temperature ranged from  20° to 70°C and the applied stress from 58,000 p s i to 72,000 p s i .  The  Arrhenius plot i s shown i n f i g . (17). The l i n e s are p a r a l l e l to each other within the experimental error. The apparent activation energies are summarised following  i n the  table. Table (10)  ACTIVATION ENERGIES OF CRACK INITIATION FOR Al-21.5Zn IN ETHYL ALCOHOLWATER MIXTURES Alloy Composition wt % Zn  Applied stress psi  Environment  Temperature range,°C  Apparent Activation energy kcals/gm.mole  Al-21.5Zn  64,600  95%C H 0H  20-70  11.1 ± 1.3  "  58,000  95%C H 0H + 5% water  20-45  11.0 +1.6  "  72,000  99.9%C H 0H +0.1% water  20-30  11.0 ± 3.9  "  58,000  99.9%C H 0H +0.1% water  20-20  12.0 ± 2.3  2  2  5  5  2  2  5  5  — 0 ~ - 0 ~  Figure (17).  Logarithm of crack i n i t i a t i o n rate vs 1/T f o r Al-21.5Zn  9 5 %  Et  hanol  i n ethyl alcohol.  T°K  Al-22%Zn 95% Alcohol  —O—O-  99.9%  Alcohol  - 55 When the rate of the crack i n i t i a t i o n was divided by the absolute temperature and the logarithm of this number plotted against the inverse of absolute temperature, the plot was once again a straight l i n e i n a l l four cases.  The apparent a c t i v a t i o n energies as calculated from  the slopes were s l i g h t l y d i f f e r e n t from the values shown i n the last table (10).  From f i g .  (18) i t i s seen that  Q_  Rate = T • A e  RT  (2)  This equation incorporates the effect of temperature on the s t r a i n induced d i f f u s i o n of hydrogen i n the l a t t i c e (90). Table (11) ACTIVATION ENERGIES FOR Al-21.5Zn FROM LOG RATE/T vs 1/T PLOT Alloy composition wt % Zn  Applied stress psi  Temperature range  Al-21.5Zn  64,000  95%C H OH + 5% water  20 to -70  10,.4 + 1,.2  58,000  95%C H OH + 5% water  20 to -45  9,.4 + 1..4  72,000  99.9%C H 0H + 0.1%H 0  20 to -30  10,.5 + 3,.7  58,000  99.9%C H 0H + 0.1%H 0  20 to -20  11..6 + 2..2  Environment  2  2  5  5  2  2  5  5  2  2  Activation energy kcal/gmmole  Therefore the given data f i t both equations 1 and 2.  3.6.2.  Applied Stress vs. Failure Time In f i g . (19) the applied stress i s plotted against logarithm  of f a i l u r e time f o r notched Al-21.5Zn i n 99.9% ethyl alcohol.  A  straight line f i t i s obtained which means that f o r constant temperature f a i l u r e time i s related to the exponential of the applied stress as follows:  O  I n i t i a t i o n time  60 A  50  —  40  —  Total  time  \CA  O^pA  IT) If)  <u  u  T3 <D •H I—I  CA \ .  30  OH  CA  <  20  1 io  l  Figure  1  2 (19).  5  1 ,1 10" 2 F a i l u r e time, Min.  The e f f e c t o f a p p l i e d ethyl alcohol.  l 10'  s t r e s s on l o g f a i l u r e time o f Al-21.5Zn  a l l o y i n 99  - 57 -  t  f  = K^e"*  (3)  2 0  where t ^ = f a i l u r e time a = applied stress K^K^ = constants  It w i l l be shown i n the discussion that f a i l u r e time t ^ can also be related to the applied stress a by a power law of the form  t  3.6.3.  f  = Kf  (4)  n  Al-2.6Mg-6.3Zn This a l l o y was stress corroded i n three environments. 1.  99.9% ethyl alcohol +0.1% water  2.  3.5% NaCl + 2.0% K_CrO. 1 4  3.  Double d i s t i l l e d water.  The a l l o y was used i n the overaged condition.  In f i g . (20) the  logarithm of crack i n i t i a t i o n rate i s plotted against the inverse of absolute temperature f o r 99.9% ethyl alcohol environment.  The apparent  activation energy of the rate c o n t r o l l i n g process i s 14.7 kcals/gmmole. When a s i m i l a r analysis i s carried out i n 3.5% NaCl I^CrO^ solution, a straight line i s obtained.  +2.0%  But the apparent a c t i v a t i o n  energy i s now 26 kcals/gmmole, as shown i n f i g . (21a). The activation energy i s 26 * 6kcals/gm mole i n double d i s t i l l e d water.  The large deviation r e s u l t s from the fewer data points.  Since  double d i s t i l l e d water and s a l t solution were studied at the same applied  Al-2.6Mg-6.3Zn Ethanol Q = 14.7 kcals/raole  1  2.7 Figure (20).  3.0 Logarithm of crack i n i t i a t i o n r a t e v c i /T -P^ in ethyl alcohol ° 7  f  r&  n  o v e r a  S  e d  Al-2.6Mg-6  2.7  3.0  3.5 10 /T 3  4.0 °K  _1  Figure (21). Logarithm of crack i n i t i a t i o n rate vs 1/T for Al-2.6Mg-6.3Zn i n NaCl/K CrO solution and water. 4 2  Al-8.6Mg 1  Ethanol Q = 9.6  kcals/mole  - 61 -  stress, the temperature range giving reasonable i n i t i a t i o n than 24 hours) was narrower i n water.  times (less  Consequently, fewer points were  obtained i n double d i s t i l l e d water. Therefore, the same process seems to operate during stress corrosion of Al-2.6Mg-6.3Zn i n d i s t i l l e d water and NaCl/K^CrO^, whereas a different process appears to be rate c o n t r o l l i n g i n the ethyl alcohol.  3.6.4.  Al-9Mg 3  In f i g . (22) the log i n i t i a t i o n  rate i s plotted against 10 /T  in 99.9% ethyl alcohol-0.1% water, environment. The apparent activation energy as calculated from the slope i s 9.6 kcals/gm mole.  The temperature range used was between 39°C and -11°C.  The error i n activation energy i s _ 2.9 kcals/gm mole. 3.6.5.  Apparent Activation Energy During Propagation The propagation rate was assumed to be constant during  cracking and was defined as the inverse of propagation time as measured from the load relaxation curve.  According to Kraft and Mulherin (77),  during crack propagation the stress intensity factor K^ increases to a l i m i t i n g value K^, when mechanical fracture takes place at an i n d e f i n i t e l y fast crack v e l o c i t y . Because the propagation rate i s measured under an unsteady state condition, the Arrhenius plot i s d i f f i c u l t to interpret.  Wei (82)  obtained apparent activation energies for corrosion fatigue a l aluminum alloys under constant stress intensity factor and found them to be inversely proportional to K .  - 62 The Arrhenius plots for propagation rates are presented i n figures  (23-16).  Only two systems show a s i g n i f i c a n t c o r r e l a t i o n  between log rate and 1/T.  They are Al-2.6Mg-6.3Zn i n NaCl/K Cr0 2  4  solution, f i g . (23) and Al-21.5Zn i n 99.9% ethanol f i g . (26) which y i e l d apparent activation energies of 14.7 and 3.8 respectively. plots i n figures  The  (24,25) show considerabel scatter and cannot y i e l d a  meaningful activation energy.  However, i t appears that crack propagation  in aluminum alloys i n alcohol i s r e l a t i v e l y temperature  insensitive.  There are several reasons f o r the wide scatter i n propagation rates. 1. only  At the stress levels used, the propagation times  5% of the t o t a l time.  Small errors i n separation  constituted  of i n i t i a t i o n  and propagation stages caused large errors i n propagation time. 2.  The stress intensity factor was constantly  increasing  during propagation and this caused errors i n the measured rate. The value of Q l i e s i n the range of 14 kcals/mole for Al-Mg-Zn in double d i s t i l l e d water and i n the range 1-4 kcal/mole for aluminum alloys i n alcohol.  3.7.  Hydrogen Charging  Experiments  The following tables  (12,13) l i s t the type of charging and the  f a i l u r e time of Al-21.5Zn and Al-2.6Mg-6.3Zn i n NaCl/K Cr0 2  4  solution.  Notched specimens were tested at a constant stress of 10,000 p s i i n Al-21.5Zn and 76,500 p s i i n Al-2.6Mg-6.3Zn. carried u n t i l the moment of load application.  Hydrogen charging was It was not possible .to  separate i n t i a t i o n and propagation times i n the Al-Zn a l l o y .  The charg-  ing time i n table (12) was 30 minutes at 4.5 v o l t s , 0.5 amps with a current  - 63 -  Figure (23). Logarithm of propagation rate vs 1/T f o r Al-2.6Mg-6.3Zni NaCl/K.CrO. and D. water \  Al-2.6Mg-6.3Zn 99.9% Ethanol  Al-8.6Mg 99.9% Ethanol  Poor correlation  Poor correlat  0 0 0 1 3.1  3.2  J 3.3 10 /T 3  iO  3.4 °K  _1  •  3.5  i 3-  1  • 3.2  ' 3.3  Q  i 3.4  J — I  3.5  10 /T 3  0  3.6 K  L 3.7  _ 1  Logarithm of propagation rate vs 1/T for Al-Mg-Zn and Al-Mg alloys i n ethyl  alcohol  3.  0 Al-21.5Zn  95% ethanol  Q = 1.2 kcals/gm mole  0  .  0  0 $  Poor Correlation  0  2  OH  o u  0  0  Cu  DO O —1  J 3.7  I  I 4.0  I  I  I 10 /T 3  Figure (25).  - —0  1  I 4.5  L  _  j  L 5.0  PK"  1  Logarithm of propagation rate vs 1/T for Al-21.5Zn i n 95% ethyl alcohol  0 Al-21.5Zn 99.9% E t h a n o l - Q = 3.86 kcals/gm  mole  OO  2  ..  /  log Prop. Rate  1  0 3  1 3 .4  3.5  1  I  1  1  I  10 /T (26).  Logarithm  1  4.0 3  Figure  0  i  r  i  1 4.5  °K  1  o f p r o p a g a t i o n r a t e vs 1/T f o r Al-21.5Zn i n 99% e t h y l  alcohol  - 67 2 density of 0.6 amps/cm . The f a i l u r e times were shortened by a factor of three by p r i o r hydrogen charging of the specimen.  Table (12) HYDROGEN CHARGING EXPERIMENTS ON Al-21.5Zn IN 3.5% NaCl + 2.0% K C r 0 2  4  AT 10,000 p s i Experimental condition  Avg. time  Hydrogen charged Hydrogen charged  37.2 37.4  Hydrogen charged Uncharged Uncharged  Total time minutes  Voltage § current  Current density  0.5A, 4.5 Volt  0.6 Amp/cm^  33.0  "  "  41.9  "  »  126 103  Uncharged  123 60  Hydrogen charged held 8 days tested uncharged  110 121.5  Hydrogen charged held 8 days tested uncharged  133.1  Cathodic protection  228.6  0.5A, 4.5 Volt  0.6 Amp/cm'  - 68 -  Table  (13)  HYDROGEN CHARGING EXPERIMENTS ON AGED Al-2.6Mg-6.3Zn ALLOY IN 3.5% 2.0%  Experimental condition  K Cr0 2  4  AT 76,500 p s i  Failure time minutes  Hydrogen charged  Avg. min.  3.5 11.6  6.4  4.1 Uncharged  NaCl +  Voltage Current  Current density  4,.5V,  1 .5 A  1.8  4,.5V,  2 A  2.4  4,,5V,  2 A  2.4  48.9 25.5  --  37  --  35.75 Cathodic  protection  Cathodic proection removed  No f a i l u r e 151 min.  4.,5V,  2 Amp.  Failure i n 1 min.  3.8.  Electron Microscopy  3.8.1.  Precipitate D i s t r i b u t i o n After Aging The p r e c i p i t a t e d i s t r i b u t i o n was studied by taking  indirect  carbon replicas from etched specimens.  The replicas were shadowed with  chromium.  for certain a l l o y systems.  Thin f o i l s were also studied  The  microstructure of aged Al-9Mg as revealed by thin f o i l and r e p l i c a technique i s shown i n figures  (27-31) . The microstructure i s  characterised  by an almost continuous network of grain boundary precipitates surrounded by a p r e c i p i t a t e free zone having a width between 2 to 10 u depending on  - 69 -  the aging time.  The grain i n t e r i o r consists of rod l i k e precipitates  believed to be the p a r t i a l l y coherent phase B ' (31,32,33). The i n t e r p a r t i c l e spacing as well as the width of the precipitate free zone decreases with increasing time of aging at 200°C u n t i l i t reaches a minimum a f t e r aging f o r 7 hours.  The minimum spacing  of precipitates as determined by the method of Kocks (68) i s Overaging leads to an increase i n the spacing to <v 4 u.  /A/N= lu  The grain  boundary p r e c i p i t a t e has been shown to be the phase 3' by d i f f e r e n t workers (32,46). The microstructure of the aged Al-3Mg-6Zn a l l o y i s shown i n figure (32,33).  Once again grain boundary p r e c i p i t a t e s , a p r e c i p i t a t e  free zone and matrix precipitates are present.  The grain boundary  precipitate i s presumably MgZi^ while the p r e c i p i t a t e i n the grain i n t e r i o r i s either the h.c.p. t r a n s i t i o n phase n (52) or the phase x (53). Both the i n t e r p a r t i c l e spacing and the width of the p r e c i p i t a t e free zone are smaller i n the aged Al-3Mg-6Zn a l l o y compared to the aged Al-9Mg a l l o y . The present results agree with the electron microscopy of Al-3Mg-6Zn and Al-9Mg alloys carried out by other workers (31-39,52,53). Extensive information i s available on the electron microscopy and X-ray d i f f r a c t i o n study of these two systems as well as Mg-8A1 and Al-22Zn alloys (41-45).  For this reason, only a limited amount of electron  microscopic study of these alloys was attempted i n the present work.  - 70 -  Figure  Figure  (27). Al-9Mg aged 7 hours a t 200°C. Matrix p r e c i p i t a t e d i s t r i b u t i o n  (28). Al-9Mg aged 7 hours a t 200°C. g.b.ppt. 10,000 x.  10,000 x.  (a) 2000 x F i g u r e 30.  (b) 3000 x  Al-9Mg, aged a t 200°C. Showing the p r e c i p i t a t e f r e e zone, (a) Aged 4 hours  (b) Aged 5 h o u r s .  - 72 -  (a)  2000 x  c (1400 x) Figure (31).  d (1400 x) Al-9Mg aged at 200°C. Precipitate d i s t r i b u t i o n . (a) aged 6 hours, (b) 7 hours, (c) 24 hours, (d) 48 hours.  - 73 -  (b) 20,000 x Figure 32.  Shows the difference i n precipitate size between the aged Al-9Mg and Al-3Mg-6Zn. (a) Al-9Mg aged for 7 hours at 200°C. (b) Al-3Mg-6Zn aged for 30 minutes at 160°C.  - 74 -  - 75 -  (c) 2000 x Figure 34.  Al-9Mg aged 5 hours at 200°C. Strained (a) 1.3 % (b) 4 %  (c) 4 %.  Figure  (36). Mg-8.6A1, solution treated and water quenched. granular stress corrosion crack 80 x  Trans-  - 77 -  S l i p Line Study Age hardened specimens of Al-9Mg were deformed 0.7%, 3%, and 4% i n an Instron.  1.3%,  Single stage s i l i c o n monoxide replicas were  taken from the surface of the t e n s i l e specimens which were etched deformation.  The results are shown i n figure (34).  observed up to 2% s t r a i n .  2%,  No s l i p  before  was  Some s l i p was v i s i b l e at 3% s t r a i n while  extensive s l i p and c r o s s - s l i p were seen at 4% s t r a i n .  The s l i p lines  were present throughout the grain and not just confined to the p r e c i p i t a t e free zone as reported by Thomas (32). Stress Corrosion Crack Path The aluminum base alloys exhibited intergranular stress corrosion cracking. fractographs  The intergranular nature of the cracking i s evident from  shown i n figures (37-41) .  The crack path i n Mg-8.5A1 was figures (35,36).  t r a n s c r y s t a l l i n e as shown i n  It follows a d i r e c t i o n perpendicular to the d i r e c t i o n  of the applied s t r e s s .  3.8.2.  Fractography A detailed fractographic study was  carried out using the electron  microscope andthe electron microprobe analyzer.  The  l a t t e r , used as a  low magnification scanning microscope was useful i n locating the area of the fracture surface from which the r e p l i c a was pictures were of two  types:  a)  topography image  b)  absorbed electron image  taken.  The microprobe  - 78 -  For each specimen, the r e p l i c a micrographs are shown a f t e r the scanning electron micrographs.  The fractographs shown i n figures (37-  40) are taken from specimens which exhibited complete stress corrosion failure.  The c r i t e r i a for this type of f a i l u r e was that the load should  gradually decrease to zero i n the Instron load relaxation test.  In  other words, there was no t e n s i l e overload component r e s u l t i n g i n an abrupt load drop.  Therefore, the features shown i n figures (37-40) can  be considered part of a ' t y p i c a l ' stress corrosion surface. Fractographs of notched specimens pulled to fracture were also studied i n order to distinguish the stress corrosion surface from the t e n s i l e overload fracture surface.  These are discussed i n appendix  (1). The p a r t i a l l y stress corroded specimens showed the most complex fracture features because they contained the stress corroded as well as the mechanically f a i l e d areas.  In such cases i t was very d i f f i c u l t to  distinguish between the two types of areas, p a r t i c u l a r l y on a r e p l i c a . An example i s shown i n figure (41) .  Stress corrosion area i n this  specimen was present near the edge while the mechanically f a i l e d area was near the centre. mechanical  However, part of the edge showed t r a n s c r y s t a l l i n e  fracture. A t y p i c a l 'dry' t e n s i l e fracture of the notch showed the  following features. 1.  Dimples of cusps were present on the intergranular facets.  2.  S l i p steps were seen on intergranular facets.  3.  Elongated dimples were present on the grain facets i n  some cases.  - 79 4.  A narrow band of transgranular fracture containing 'rim  l i n e s ' was present at the edge i n Al-Zn and Al-Mg-Zn a l l o y s . 5.  The 'dry' fracture was predominantly intergranular.  A l l the above features were r a r e l y observed together i n the completely  stress corroded specimen.  S l i p steps were present i n such  specimens and a small f r a c t i o n of the edge showed rim lines but most of the surface appeared smooth and bright at low magnification.  At  high magnification the grain facets appeared to have a rough surface. The edge could be i d e n t i f i e d from the machine marks.  Similar to those  seen i n figure(37a,b). The stress corroded surface can be seen i n figure (37) . The intergranular cracking has occurred over 95% of the edge shown i n figure (37a,b).  Only 5% i s covered by transgranular rim l i n e s .  Near  the centre, s l i p steps can be seen i n figure (37c) . The rough nature of the surface i s evident from figures (37d,e).  The specimen was an  Al-22Zn a l l o y , stress corroded i n NaCl/K^CrO^ at low s t r e s s . The same alloy was hydrogen charged and stress corroded, at a low stress.  Again 95% of the edge showed intergranular cracking  while 5% had transgranular rim lines as seen i n figure (38a,b). product can also be seen.  Corrosion  The facets shown i n figure (38c) appeared to  have surfaces shown i n figures (38d-h) when examined at high magnifications.  While corrosion product i s present i n figures (38e,f) the  roughness of the suface can be seen i n areas free of corrosion product such as f i g . (38d,f,g,h). of steps.  In figure (38h) the surface appears to consist  37(c) 240 x  37(d) 2500 x  37 (e)  10,000 x  Figure (37). Aged Al-22Zn a l l o y , stress corroded i n NaCl/K Cr0 2  4  at  10,000 p s i . Complete stress corrosion fracture. (37a). Stress corrosion surface near the edge of the notch with machine marks. (37b). Transgranular fracture near the end. (37c). S l i p lines near the centre showing p l a s t i c deformation during stress corrosion. (37d,e)Stress corrosion surface.  38(c) 240x  38(d) 5000 x.  - 83 -  38(g) 7000 x  38(h) 20,000 x  - 84 -  Figure  38.  Al-22Zn, hydrogen charged f o r 30 minutes and  s t r e s s corroded  i n NaCl/K^CrO^, a t 10,000 p s i .  38(a)  Edge, showing r i m l i n e s and  corrosion product.  38(b)  Edge, s t r e s s c o r r o s i o n s u r f a c e  38(c)  Centre, s t r e s s c o r r o s i o n surface, i n t e r fracture.  38(d)  T y p i c a l s t r e s s c o r r o s i o n s u r f a c e w i t h a rough texture.  and  slip  lines. granular  3 8 ( e , f ) C o r r o s i o n product i n the s t r e s s corroded a r e a . Dark l i n e s i n 38(f) are f o l d s i n the r e p l i c a . 38(g,h) Another type of s t r e s s c o r r o s i o n surface a t two d i f f e r e n t m a g n i f i c a t i o n s . The rough area appears to c o n s i s t of s t e p s .  - 85 -  When Al-2.6Mg-6.3Zn i s s t r e s s corroded g r a i n f a c e t s show s l i p observed  over t h e f r a c t u r e s u r f a c e .  branching o f f i n f i g . in f i g . of  s t e p s , f i g . (39a,b).  (39c-e).  (39b).  in fig.  2  4  Intergranular cracking i s  I n t e r g r a n u l a r c r a c k s can be seen  The rough t e x t u r e o f t h e f a c e t s i s shown  Some evidence o f d u c t i l i t y  corroded dimples  i n NaCl/K Cr0 , the  (39e).  i s shown by t h e presence  T h i s specimen showed a stepwise  load  relaxation. When Al-9Mg i s s t r e s s c o r r o d e d i n NaCl/I^CrO^ s o l u t i o n , the f r a c t u r e i s completely slip  i n t e r g r a n u l a r w i t h no evidence o f r i m l i n e s o r  steps as seen i n f i g .  (40a).  The s u r f a c e has a c o r r o d e d  when viewed a t h i g h e r m a g n i f i c a t i o n , f i g s .  (40b-d).  appearance  The lower  right  area i s p r o b a b l y caused by t h e t e a r i n g o f r e p l i c a . The  complex n a t u r e o f t h e f r a c t u r e s u r f a c e o f a p a r t i a l l y  s t r e s s corroded specimen i s b e s t i l l u s t r a t e d by f i g . was  s t r e s s corroded  i n a l c o h o l a t 30,000 p s i .  (41).  The specimen  The a r e a o f s t r e s s  c o r r o s i o n was a s m a l l f r a c t i o n o f t h e t o t a l a r e a .  More than h a l f t h e  edge showed a t r a n s g r a n u l a r f r a c t u r e with r i m l i n e s , w h i l e t h e t e n s i l e f r a c t u r e near the c e n t r e was i n t e r g r a n u l a r . are p r e s e n t i n f i g . (41a,b).  Rim l i n e s and s l i p  A m a g n i f i e d view o f t h e t r a n s g r a n u l a r  f r a c t u r e s u r f a c e w i t h r i m l i n e s i s shown i n f i g .  (41c).  the s m a l l area o f s t r e s s c o r r o s i o n s u r f a c e near t h e edge. shear dimples in f i g .  and s l i p  (41e,f).  environment.  steps  F i g . (41d)  is  Dimples,  l i n e s associated with d u c t i l e f r a c t u r e are present  The dimples  i n F i g . (41f) have been a t t a c k e d by t h e  39 (c) 6000 x  39(d) 7500 x  - 87 -  39 (e) 9500 x  Figure  (39). Al-2.6Mg-6.3Zn,  s o l u t i o n t r e a t e d and water quenched, s t r e s s  corroded i n NaCl/K^CrO^ a t 55,000 p s i . Predominantly corrosion  surface.  39(a).  Edge, i n t e r g r a n u l a r f r a c t u r e s u r f a c e w i t h s l i p during s t r e s s c o r r o s i o n .  39(b)  Centre, i n t e r g r a n u l a r corroded a r e a .  extensive  f a c e t s , deformation i n s t r e s s  39(c,d). C h a r a c t e r i s t i c s t r e s s c o r r o s i o n surface with product or t e a r marks. 39(e).  stress  Dimples w i t h rough s u r f a c e region.  i n the s t r e s s  corrosion  corrosion  -  40(c) 3300 x  88 -  40(d) 3300  x.  - 89 Figure  40.  Aged Al-9Mg, s t r e s s c o r r o d e d i n NaCl/K^CrO^ Predominantly s t r e s s c o r r o s i o n 40(a)  a t 7000 p s i .  surface.  Edge, i n t e r g r a n u l a r f a c e t s w i t h a b r i g h t appearance.  40(b,c) S t r e s s c o r r o s i o n s u r f a c e showing a rough 40(d)  texture.  S t r e s s c o r r o s i o n s u r f a c e , the bottom r i g h t a r e a i s presumably an a r t i f a c t caused by the t e a r i n g of r e p l i c a .  - 90 -  91  41(e) 9,000 x  Figure (41).  41(f) 9,000 x  Al-22Zn stress corroded i n 99.9% ethyl alcohol at 30,000 p s i . Stress corrosion and t e n s i l e overload modes of failure.  (41a).  Near edge, transgranular mechanical fracture showing rim lines.  (41b).  Near centre, deformation overload f a i l u r e .  (41c).  Near edge, magnified view of rim lines shown i n (41a).  (41d).  Near edge, stress corrosion surface.  on grain facets during t e n s i l e  (41e,f).Near centre, t e n s i l e overload region with dimples and s l i p l i n e s . Note the roughening of the dimples by alcohol.  - 92 -  Figure (42).  Phase diagrams of Al-Mg and Al-Zn systems. The Composition and heat treatment temperatures of the binary alloys used are indicated (112) .  4.  DISCUSSION  4.1.  P r e c i p i t a t i o n Hardening and M e c h a n i c a l  Properties  In t h i s s e c t i o n the p r e c i p i t a t i o n hardening r e a c t i o n s and the nature o f r e s u l t i n g p r e c i p i t a t e s i s c o n s i d e r e d . i s then c o r r e l a t e d t o mechanical p r o p e r t i e s ductility.  The  such as y i e l d  microstructure strength  and  The phase diagrams Al-Mg and A l - Z n systems a r e shown i n  f i g u r e 42.  4.1.1.  Al-9Mg aged at 200°C The aging  sequence o f t h i s a l l o y can be r e p r e s e n t e d  supersaturated  solid solution  G e i s s l e r , B a r r e t and Mehl  -»• g'  ->-  as  BCN^Al^)  (46) have e s t a b l i s h e d t h a t the  t r a n s i t i o n s t r u c t u r e 6' i s a p a r t i a l l y coherent phase which forms as p l a t e l e t s on [120] o f the matrix and t h a t g phase i s incoherent the m a t r i x .  with  The s t r u c t u r e o f 3' i s u n c e r t a i n but g phase i s complex f . c . c .  with 1168 atoms p e r u n i t c e l l (62). P r e c i p i t a t i o n at g r a i n boundaries can be observed m i c r o s c o p i c a l l y before  aging  has p r o g r e s s e d f a r enough t o change y i e l d s t r e n g t h  e l o n g a t i o n measureably.  and  Changes i n y i e l d s t r e n g t h and e l o n g a t i o n a r e  approximately c o i n c i d e n t with appearance o f p r e c i p i t a t i o n w i t h i n the  - 94 grains  (31).  increase 31)  Grain boundary precipitates do not contribute  in strength.  to the  The p r e c i p i t a t e free zone observed i n f i g s .  (27-  i s denuded of solute and vacancies. Nicholson and Embury (47) have shown that nucleation  3'  of  requires vacancies because i t has a larger atomic volume than the matrix. Apparently the vacancy deficiency cannot be relieved by the boundary acting as vacancy source.  grain  Therefore the rod l i k e 3'  cannot nucleate i n the zone adjacent to the grain boundary. other hand ppt. nucleation  precipitates On  the  at the grain boundary i s favoured because  of a lowering i n the value of free energy change.  The  grain  boundary  precipitate i s then i n a lower energy state than the intermediate precipi t a t e formed within the grain and i t w i l l tend to increase  i n size  during aging at the expense of intragranular p r e c i p i t a t e s . the grain boundary p r e c i p i t a t e i s frequently granular p r e c i p i t a t e and  In addition  nucleated before the  i s able to draw solute atoms from within  intrathe  grains. The nucleation of precipitates i s d i f f i c u l t i n Al-9Mg a l l o y and occurs heterogeneously.  A comparison of f i g . (27) with f i g . (33)  indicates that the precipitates i n Al-9Mg are coarse and the density p r e c i p i t a t i o n i s low.  of  This i s the explanation of the poor age-hardening  c h a r a c t e r i s t i c s of binary Al-Mg alloys compared to ternary Al-Mg-Zn a l l o y s . Most of the p r e c i p i t a t i n g phase at aging times employed i n this work i s the t r a n s i t i o n p r e c i p i t a t e 3 '  (31).  Fink and Smith obtained  only 3' up to 136 hours at 200°C for an Al-lOMg a l l o y (31). the p r e c i p i t a t e becomes coarser.  On overaging  Results of Perryman and Brooks  as well as Thomas and Nutting (32) show that the equilibrium  (67)  precipitate  - 95 on aging at 200°C does not appear on peak aging but appears well into the overaging stage.  4.1.2.  Al-3Mg-6Zn Aged at 160°C The aging sequence i n this a l l o y has been generally accepted  as (51). Solid solution ->- Spherical G.P. Zones -»• n' ^ been suggested by Mondolfo et a l (52) that  nCMgZn^).  It has  n' i s a h.c.p. t r a n s i t i o n  phase which forms with the following orientation relationship with the matrix:  (001) n  f  II ( H l ) A l  Further i t i s proposed that as aging progresses the n  1  phase  gradually  transforms into n phase by a steady change i n l a t t i c e parameters. The n phase i s also hexagonal and occurs with at least six d i f f e r e n t orientation relationships with the matrix (53). The paths which are most frequent have an orientation r e l a t i o n s h i p :  (0001) n II (HO)Al  According to Thackery (53), specimens aged at 165°C for 1, 5, 16, and 120 hours a f t e r a cold water quench from 465°C did not show any sign of n'.  The p r e c i p i t a t e consisted  phase X and equilibrium phase uncertain,  n•  of a hitherto unreported  The c r y s t a l structure of X phase i s  though probably hexagonal, with Zn as the predominant element. The aging treatment i n the present work was very close to that  - 96 used by Thackery and therefore  the structure probably consists of disper-  sion of X and n phases, along with spherical G.P. zones. The probable aging sequence i s :  Spherical G.P. zones  -> n (MgZn ) 2  From f i g . (32-33) i t i s evident that the p r e c i p i t a t e size and i n t e r p a r t i c l e spacings i n Al-Mg-Zn alloys are considerably in Al-9Mg a l l o y .  less than  The p r e c i p i t a t e free zone varies from 0.2 to 0.5 u  i n width compared to 2-10 u f o r Al-9Mg.  The p r e c i p i t a t e i n f i g . (27)  i s at least an order of magnitude larger than the p r e c i p i t a t e i n f i g s . (32-33).  The width of the p r e c i p i t a t e free zone obtained here agrees  well, with that obtained by Embury and Nicholson under comparably aging ;  conditions  (54).  However, these authors contend that the predominant  phase after aging at 180°C for 3 hours i s n 1 . Thackery (53) suggests that the so-called actually be a type of n p r e c i p i t a t e .  n  1  plates may  It i s also possible that  n' may  in fact be the X phase.  4.1.3.  Al-21.5Zn Aged at Room Temperature Garwood and Davies (55) have plotted aging curves f o r 'step-  quenched' Al-Zn alloys containing  17% and 25% Zn aged at room temperatures.  From t h e i r data i t i s possible to calculate the time required 21.5Zn a l l o y to a t t a i n maximum hardness at room temperature. was approximately 1.5 hours. least 11 days.  for an A l This time  The maximum hardness was maintained for at  -  The aging sequence  97  -  at elevated  temperatures as proposed by  Carpenter and Garwood (56) i s : Supersaturated s o l i d  solution  s p h e r i c a l G.P. zones ->  coherent p l a t e l e t s o f rhombohedral s o l i d solution of equilibrium  a -»- f c c a ' ->• hep z i n c  rich  composition.  A c c o r d i n g t o N i k l e w s k i (57) e t a l the coherent p r e c i p i t a t e s are s t a b l e below  110°C.  In a specimen aged t o peak hardness a t room temperature t h e r e o  i s a mixed  structure  o f s p h e r i c a l G.P. zones ^ 35 A i n diameter and  p l a t e l e t s o f a' c o n j u g a t e s o l i d s o l u t i o n c o n t a i n i n g o  Zn  (57).  on  [111] h a b i t  a p p r o x i m a t e l y 80% o  These p l a t e l e t s a r e 100 A i n d i a and ^ 20 A t h i c k . p l a n e s o f the m a t r i x .  The i n c r e a s e  They form  i n hardness on a g i n g  at room temperature can be a t t r i b u t e d t o these two coherent phases. 4.1.4.  Mg-8.5A1 aged a t 200°C The p r e c i p i t a t i o n r e a c t i o n Supersaturated s o l i d  phase has a complex c u b i c  i n Mg-8.5A1 a l l o y i s (58-61):  s o l u t i o n ->• E q u i l i b r i u m  Mg^Al^  '^17^12  structure.  No i n d i c a t i o n o f a t r a n s i t i o n l a t t i c e form o f Mg^ A l ^ G.P. zone f o r m a t i o n were observed even a t low temperatures d u r i n g  o  r  early  stages o f p r e c i p i t a t i o n (58-61). It has been shown (61) t h a t  the M g ^ A l ^ p l a t e s  form p a r a l l e l  to the b a s a l  p l a n e and a r e e l o n g a t e d i n t h r e e c l o s e packed d i r e c t i o n s o f  the m a t r i x .  Under optimum hardening c o n d i t i o n s  the i n t e r p l a t e s p a c i n g  o  was r e l a t i v e l y  l a r g e about  around the p r e c i p i t a t e .  2000 A and no coherency s t r a i n s were observed  - 98 -  Both intragranular and intergranular p r e c i p i t a t i o n occurs competitively.  C e l l u l a r p r e c i p i t a t i o n takes place a high angle or disordered  grain boundaries.  C e l l u l a r p r e c i p i t a t i o n i n this Mg-8.5A1 a l l o y i s  believed to occur by the mechanism proposed by Smith (61) i . e . a grain boundary migrates into the adjacent  grain with lamellae of p r e c i p i t a t e  and depleted matrix forming behind the advancing grain boundary.  4.1.5.  Age Hardening Mechanism of Mg-Al According  to Clark (61), the r e s u l t s of trace analysis on large  grained specimens of Mg-9A1 show that d i f f e r e n t modes of deformation are operative i n the solution treated and f u l l y age hardened condition. Basal s l i p and  [10l2] twins are the p r i n c i p a l deformation modes  in the solution treated condition.  In the f u l l y aged condition, basal  s l i p and prismatic s l i p predominate. absent.  [1012] twins are almost completely  Fine basal s l i p i n a solution treated specimen i s replaced  by  wavy s l i p i n a f u l l y hardened specimen. Transmission microscopy shows that as aging proceeds, planar arrays of basal d i s l o c a t i o n decrease and areas of complex tangles increase.  These tangles are responsible for hardening the alloy., The  [1012] twins i n Mg-Al a l l o y do not shear the M g ^ A l ^  plates and as the size and number of M g ^ A l ^ plates increase the amount of [1012] twinning proportionately decreases. In view of the simple p r e c i p i t a t i o n process i n Mg-Al i t i s not surprising that the agehardening response i s small.  The 0.2% o f f s e t  y i e l d stress increases from 27,000 p s i to 40,000 p s i e.g. 50%. form p l a s t i c elongation drops: from 4.5%  The uni-  to 1% and r i s e s s l i g h t l y on  - 99  Figure (43).  Plates of equilibrium M g A l 1 7  1 2  phase with basal plane habit,  Arrows denote [1120] d i r e c t i o n s . Mg-9A1 aged 8 hours at 260°C.  After Clark (61) .  - 100 overaging as shown i n f i g . (6). The  interplate spacing i n Mg-Al alloys i s not small enough to  require shearing of the p r e c i p i t a t e for d i s l o c a t i o n motion and result in a large increase i n strength.  The role of M g ^ A l ^ precipitates i s to  break up basal g l i d e , cause cross s l i p and thus produce d i s l o c a t i o n tangles which i n effect strainhardens the matrix.  However, the  increased  strength produced by s t r a i n hardening i s not as large as i n systems such as Al-Cu and Al-Mg-Zn.  4.1.6.  Effect of Aging on Mechanical Properties The  increase of nearly  200%  of Aluminum Alloys  i n the y i e l d strength of aged  Al-Mg-Zn i s remarkable when compared with a mere 30% alloys.  increase in Al-Mg  On the other hand the uniform elongation on aging, decreases  by a factor of 4 i n in Al-Mg, f i g s .  Al-Mg-Zn compared to a decrease by a factor of  24  (4,5).  In Al-21.5Zn a l l o y the increase i n hardness for the given heat treatment i s about 100%  (55) .  The strength should increase by a pro-  portional amount,(48).  The following q u a l i t a t i v e description can  be  given of the v a r i a t i o n i n strength of age hardening aluminum alloys (63) . I n i t i a l l y the strength of the a l l o y i s that of the supersaturated matrix. If the i n i t i a l p a r t i c l e s are very small i n at,least one dimension, the strength i s also very temperature dependent.  The y i e l d stress at this  stage i s governed by the stress necessary to force d i s l o c a t i o n through the p r e c i p i t a t e s . precipitates may  The work done i n forcing dislocations through the be governed by factors l i k e coherency stresses,  order of precipitate and  interface e f f e c t s .  internal  - 101 -  As the p r e c i p i t a t e particles, increase i n size and possibly change their internal structure or the nature of t h e i r interface with the matrix, the work done i n cutting each p a r t i c l e increases. Dislocations are eventually force between the p a r t i c l e s , instead of through them.  According to Thomas et a l (32) the p a r t i a l l y coherent precipitates  such as g' i n Al-Mg deform with matrix. Further increase i n p a r t i c l e spacing then leads to a decrease in strength, the y i e l d strength varying as 1/d where d i s the mean planar i n t e r p a r t i c l e spacing, or according to the equation:  '  Gb 1  =  T  s  +  47  , , l n  {  d-2r, 1 ^b~ (d-2r)/2  i f Orowan c r i t e r i o n i s met In this equation: x x  r  }  ( 1 )  (102)  = i n i t i a l y i e l d stress (shear) of the hardened a l l o y . s  = y i e l d stress of the matrix material 1  d = i n t e r p a r t i c l e spacing r = radius of p a r t i c l e s G = shear modulus • .=  1 / 2 C 1  +  T  ^ )  v = Poisson's r a t i o  This sequence i s i l l u s t r a t e d schematically i n f i g . times the applied stress  (44).  At short aging  x^ necessary to force d i s l o c a t i o n between the  p a r t i c l e s i s much greater than the measured y i e l d stress.  The v a r i a t i o n  of x^ with aging time i s denoted by the dotted curve i n F i g . (44). After longer aging times  f  controls the y i e l d stress.  If the p a r t i c l e s  - 102 -  \  - 103 are s u f f i c i e n t l y strong, the flow becomes very turbulent since either d i s location loops with matrix Burger's vector or interface dislocations are l e f t around p a r t i c l e s . This produces a high density of dislocations which r e s i s t s further s l i p . In p o l y c r y s t a l l i n e alloys of aluminum the grain boundary precipitate exerts l i t t l e influence on y i e l d strength.  Usually  the  d u c t i l i t y of the alloys i s decreased by the presence of grain boundary precpitate. The poor age hardening response of Al-Mg a l l o y compared to Al-Mg-Zn and Al-Zn alloys can be explained by the following reasons:  1.  A requirement of high strength i n an aged a l l o y i s that  the p r e c i p i t a t e dispersion should be fine (49-50).  In Al-Mg a l l o y the  dispersion i s coarse compared to Al-Mg-Zn and Al-Zn a l l o y s . In Al-Mg-Zn and Al-Zn the dispersion i s f i n e r by at least an order of magnitude and hence the strengthening  i s more pronounced f i g s .  (27-33). 2.  The nucleation of precipitates i s d i f f i c u l t  i n the Al-Mg  system owing to the large unit c e l l of the equilibrium precpitate .Mg^Al^. This p r e c i p i t a t e exhibits an ct-Mn complex f.c.c. structure with atoms per unit c e l l . Al-Mg-Zn and G.P.  The nucleation of intermediate  1168  precpitate n' i n  zones i n Al-Zn i s r e l a t i v e l y easy.  Further evidence of this fact can be seen i n the age hardening response of these alloys at room temperature.  Both Al-Mg-Zn and Al-Zn  show p r e c i p i t a t i o n within a few hours of solution treatment (55,66).  - 104 Al-9Mg alloys on the other hand can r e t a i n nearly a l l magnesium i n the supersaturated s o l i d solution and microscopically v i s i b l e p r e c i p i t a t e may take years to form (64) . The decrease i n uniform elongation on aging i n Al-Mg, Al-Zn and Al-Mg-Zn alloys can be attributed to the presence of grain boundary p r e c i p i t a t e s .  From figures i n Appendix I,  i t appears that the t e n s i l e fracture i s predominantly i n t e r c r y s t a l l i n e . Intergranular precipitates can adversely a f f e c t d u c t i l i t y by nucleating voids.  Extensive void formation can be observed i n figures shown i n  Appendix  I. I n t e r f a c i a l energy would appear to be an important para-  meter and the stress concentrating effect of the precipitates may be of special importance (65). In the circumferentially notched rods the voids are nucleated on the intergranular p r e c i p i t a t e s . The crack travels towards the surface. The i n i t i a l mode of fracture i s intergranular.  When i t reaches near  the surface the stresses may be high enough f o r i t to proceed i n a transgranular manner. as i n Al-3Mg-6Zn. granular.  The mode can then be mixed i n t e r - and transgranular  In al-21.5Zn and Al-9Mg i t i s predominantely i n t e r -  In certain cases the f i n a l transgranular fracture i s accompanied  by the presence of rim lines shown i n f i g s . (37-41).  These linear  features on the edge are found i n notched and unnotched Al-Mg-Zn alloys and notched Al-Zn a l l o y s .  They are absent i n unnotched Al-Zn and Al-Mg  alloys. It appears that the predominance of intergranular fracture i s associated with the ease of void nucleation as well as the a v a i l a b i l i t y of suitably oriented grain boundary paths.  In notched specimens,the  final  - 105 -  fracture has to pass through the root of the notch even i f no suitably oriented grain boundary i s present.  Therefore the crack goes through  the grain and leaves linear markings or rim lines on the edge of the notch.  4.2.  1  Parameters Affecting Stress Corrosion Cracking i n Aluminum Alloys  4.2.1.  Stress Corrosion S u s c e p t i b i l i t y and Microstructure i n Al-9Mg Alloy Stress corrosion s u s c e p t i b i l i t y increases remarkably on iso-  thermal aging at 200°C as seen i n figure (7).  The as quenched a l l o y d i d  not f a i l even i n 600 hours whereas the peak aged a l l o y f a i l e d i n just 15 seconds.  On comparing figures (7) and (4) i t i s seen that the  stress corrosion s u s c e p t i b i l i t y varies i n the same fashion as the 0.2% offset y i e l d stress.  Since the y i e l d stress i s largely determined by  the nature of precipitates and t h e i r d i s t r i b u t i o n i t i s reasonable to believe that these factors influence the times to f a i l u r e .  The as  quenched Al-9Mg a l l o y can r e t a i n magnesium i n the s o l i d solution for a long time (66).  Isothermal aging results i n p r e c i p i t a t i o n of rod l i k e  6' phase i n the grains and grain boundaries.  3' phase corrodes i n  NaCl/K^CrO^ solution even when i t i s not coupled to a cathodic phase (69). The maximum s u s c e p t i b i l i t y to stress corrosion occurs close to peak hardness i n Al-9Mg.  During pre-aging and over-aging stages the  f a i l u r e time changes r a p i d l y .  The corresponding variations i n y i e l d  strength, i n t e r p a r t i c l e spacing and width of the PFZ are modest. satisfactory stress corrosion theory must explain:  A i  - 106 1.  The reason for the maximum stress corrosion s u s c e p t i b i l i t y  close to peak hardness i n Al-9Mg a l l o y s . 2.  The several orders of magnitude increase i n stress corrosion  s u s c e p t i b i l i t y corresponding to r e l a t i v e l y minor changes i n y i e l d strength, p r e c i p i t a t e d i s t r i b u t i o n and the width of the PFZ. The Al-9Mg alloy i s subject to i n t e r c r y s t a l l i n e corrosion i n s a l t solution due to the presence of the grain boundary p r e c i p i t a t e g or g' (69) .  Perryman found that i f an Al-7Mg a l l o y was slowly  cooled  from the solution treatment temperature, i t contained a discontinuous grain boundary p r e c i p i t a t e and was resistant to stress corrosion cracking.  Edeleanu (70) stated that stress corrosion occurred even when no  second phase was observed at the grain boundary and that hydrogen was evolved from the cracks p r i o r to f a i l u r e .  Farmery and Evans (14)  attributed cracking to an anodic reaction at the crack t i p that required the maintenance of an acid condition there.  Erdmann-Jesnitzer and  Hotzsch (69) reported that stress corrosion cracks i n Al-7Mg followed grain boundaries of large mis-orientation which were normal to the d i r e c t i o n of the applied stress.  This i s i n agreement with present  results i n which the d i r e c t i o n of crack propagation was along grain boundaries perpendicular  4.2.2.  to the applied stress.  Applied Stress and Stress Corrosion S u s c e p t i b i l i t y A threshold stress equal to 20% of the y i e l d stress was;  obtained  for peak aged Al-9Mg a l l o y , f i g . (8).  It i s seen that a change  i n stress level from 20% to 30% of the y i e l d stress r e s u l t s i n the specimen f a i l i n g i n one hour.  This go-no-go type of test, with the  - 107  -  specimen f a i l i n g i n one hour or not f a i l i n g i n 200 hours i s believed to be a feature of the chloride/chromate environment (71). level has very l i t t l e influence  upto 75% of the y i e l d stress.  after large differences  i n f a i l u r e times r e s u l t from small  on the applied  The  stress.  There-  variations  shape of t h i s curve i s very similar to  that obtained i n tests of susceptible,  4.2.3.  Stress  Nature of S l i p and  wrought Al-Mg-Zn alloys  (69).  Stress Corrosion S u s c e p t i b i l i t y  Thomas and Nutting (32) observed p r e f e r e n t i a l s l i p i n the denuded zone i n Al-7Mg a l l o y overaged at 200°C.  To test this observation,  f l a t underaged and peak aged specimens were strained d i f f e r e n t amounts in .tension and both carbon-chromium r e p l i c a s and  single stage  silicon  monoxide replicas were taken from t h e i r surfaces. No s l i p could be observed i n the peak aged specimen up i t s fracture s t r a i n of denuded zone was  1%.  visible.  to  Certainly no p r e f e r e n t i a l s l i p i n the  S i m i l a r l y , there was  in the underaged specimen up to a s t r a i n of 2%,  no p r e f e r e n t i a l  slip  f i g . (34).  Transgranular s l i p could be observed in the underaged specimen at 3% s t r a i n .  S l i p bands and  of the grain, f i g . (34b). lines and bands were seen.  s l i p lines were present i n the i n t e r i o r  At the fracture s t r a i n of 4% extensive Cross s l i p was  slip  also present, f i g . (34c).  Similar results have been obtained on Al-Mg-Cu-Zn by Jacobs (20), Speidel (22) and other workers.  They indicate that no  preferential  p l a s t i c deformation occurs i n the p r e c i p i t a t e freezone as postulated by Pugh (13) and Thomas (17) . Transmission  microscopy work of Speidel  i  (22)  - 108 reveals no evidence for a large density of dislocations i n the p r e c i p i tate free zone.  However Sedriks et a l (87) have observed d i s l o c a t i o n  tangles i n the PFZ.  They postulate that dislocations are generated at  precipitate p a r t i c l e s and move across the soft s o l i d solution i n the PFZ u n t i l they are stopped by fine p a r t i c l e s at PFZ-grain interface. At higher stresses the dislocations overcome the b a r r i e r s and move into the grains r e s u l t i n g i n the formation of bands.  4.2.4.  Stress Corrosion S u s c e p t i b i l i t y and Microstructure i n Al-3Mg6Zn A l l o y The maximum s u s c e p t i b i l i t y occurs at an aging time of 3 hours  at 160°C.  This time corresponds to a y i e l d strength of about 80% of  the peak y i e l d strength.  The occurrence of maximum s u s c e p t i b i l i t y well  before peak aging i n Ai-Mg-Zn alloys has been observed by other workers (22).  This i s i n contrast to the behaviour of Al-Mg a l l o y s , which are  most susceptible at peak strength.  The structure consists of a fine  dispersion of p a r t i a l l y coherent n ' or X phase (53).  The grain boundary  p r e c i p i t a t e s , presumably equilibrium phase MgZn are coarser than the 2  matrix p r e c i p i t a t e s and about 0.1 \i long. 0.25  p i n width (26).  The PFZ i s approximately  The i n t e r p a r t i c l e spacing appears to be about  0.1 u. Another difference from Al-Mg system, i s the fact that the as-quenched a l l o y i s susceptible to stress corrosion.  This can be  explained i n terms of the ease of nucleation of G.P. zones and subsequent phases even at room temperature.  Speidel (22) observed grain  boundary and matrix p r e c i p i t a t i o n i n Al-Mg-Zn-Cu alloy at room temperature  - 109 -  a f t e r aging f o r several hours. The width of PFZ does not vary much at the same aging temperature, (38) yet the f a i l u r e times between 3 hours and 24 hours at 160°C, vary by a factor of 10. The two heat treatments produce similar y i e l d strengths and % elongations. Sedriks et a l (26) working with a similar a l l o y , obtained three d i f f e r e n t sizes of PFZ at the same strength l e v e l , by varying the aging times and temperature.  They  observed an increase i n f a i l u r e time or a decrease i n s u s c e p t i b i l i t y with increasing width of PFZ. The present r e s u l t s indicate that the width of PFZ alone cannot account f o r the large differences i n susceptib i l i t i e s between underaged and overaged specimens.  The specimens used  by Pugh et a l (26) were i n d i f f e r e n t stages of p r e c i p i t a t i o n hardening which means that there were differences, i n the types of p r e c i p i t a t e s , i n t e r p a r t i c l e spacings, as well as PFZ. Not only do structures with similar PFZ show very d i f f e r e n t s u s c e p t i b i l i t i e s , but also specimens with same y i e l d strength and d u c t i l i t y can have widely different s u s c e p t i b i l i t i e s to stress corrosion. This can be seen from figures (5,9). It appears that the phases formed during the early stages of aging, such as G.P. zones and intermediate precipitates have a marked influence on the stress corrosion s u s c e p t i b i l i t y of Al-3Mg-6Zn a l l o y s . As aging proceeds, the proportion of intermediate precipitates decreases and that of equilibrium phase MgZn  2  increases.  s u s c e p t i b i l i t y to stress corrosion decreases.  As a r e s u l t , the Gruhl (19) observed that  heat treatments which caused the metastable coherent phases to be replaced by stable MgZn„ phase also resulted i n improved resistance to  . -  stress corrosion.  Speidel  110  -  (22) proposed that coherent precipitates are  sheared by dislocations which therefore p i l e up adjacent to the grain boundary and cause high stress concentrations.  Such precipitates also  result i n coarse s l i p because shearing of obstacles of the flow stress i n the activated s l i p plane.  leads to a reduction  The following d i s l o c a -  tions tend to remain and to move i n the same plane, thus reducing the cross section of the p a r t i c l e s even more and further s l i p i n the same plane i s f a c i l i t a t e d . Jacobs (20) and other workers have observed that  susceptible  alloys tend to show coarse s l i p and non-susceptible alloys show f i n e slip. When aging i s continued, the volume f r a c t i o n of precipitates which cannot be sheared by dislocations increases.  Such precipitates  are by-passed by dislocations and loops are l e f t around them which interact with mobile dislocations and increase  the f r i c t i o n stress.  Very overaged alloys show bands of high d i s l o c a t i o n loop density  (22).  Stress concentration at the t i p of such bands and the associated  slip  step i s very small. fig.  Such alloys are very resistant to SCC as seen from  (9).  According to Speidel  (22) the most susceptible alloys are those  which can sustain high stress concentration at the grain boundary because of d i s l o c a t i o n p i l e ups, exhibit coarse s l i p and thereby expose fresh metal surface to the environments. Al-Mg-Zn alloys before peak strength.  These conditions  are met by  Other important factors are the  potential difference between grain boundaries and i n t e r i o r of the grain, nature of the grain boundary p r e c i p i t a t e and the rate of chemical  - Ill and electrochemical  d i s s o l u t i o n of the grain boundary and localized  plastic  deformation i n the v i c i n i t y of the crack.  4.2.5.  Applied Stress and Stress Corrosion S u s c e p t i b i l i t y When the applied stress i s decreased from 90% of the y i e l d  stress to 60% the f a i l u r e time increases by a factor of two.  A  further decrease i n the stress to 30% of the y i e l d stress increases the f a i l u r e time by a factor of 5 as seen i n F i g . (10). Thus, Al-3Mg-6Zn in the peak aged condition shows a change i n f a i l u r e time with applied stress which i s very s i m i l a r to that observed i n Al-9Mg i n peak aged condition.  However, below 60% of the y i e l d stress, the rate of increase  of f a i l u r e time i s much slower i n Al-3Mg-6Zn than i n Al-9Mg.  For example,  a change i n stress level from 60% to 30% produces a f i v e - f o l d increase i n f a i l u r e time i n Al-3Mg-6Zn and a 170 f o l d increase i n Al-9Mg. Therefore, the s u s c e p t i b i l i t y  of peak aged Al-9Mg i s highly  stress sensitive compared to that of the peak aged Al-3Mg-6Zn. Gruhl (73) postulates that the rate of cracking i s related to the applied stress as  Z = K* exp (c )  where Z = rate of cracking - inverse of f a i l u r e time K* = k.exp. (- =2. ) c,k = are  constants  (2)  - 112 -  Q = activation energy of the rate c o n t r o l l i n g process Thus the rate of cracking  i s given by  Z = K  exp(  c a  - Q/RT)  (3)  The present results indicate that a plot of log. f a i l u r e time vs applied stress i s not a straight line but rather a curve.  Helfrich  (72) obtained a straight line at higher stresses turning into a curve at lower stresses.  Gruhl obtained a constant threshold  stress at  different temperatures, whereas H e l f r i c h observed a decreasing stress l e v e l with increasing temperature.  threshold  These differences could be  due to the d i f f e r e n t compositions of alloys as well as d i f f e r e n t stress corrosion testing techniques used i n the above works. stress i s plotted against  When log applied  log f a i l u r e time f o r Al-9Mg and Al-3Mg-6Zn i n  salt solution and a least squares f i t applied to the points, straight lines with correlation c o e f f i c i e n t s of 0.92 and 0.98 are obtained. are shown i n figures (45,46).  They  This indicates that applied stress (a) i s  related to the f a i l u r e time (t^) by an equation of the form  t  f  = A • a"" • exp (JL.)  (4)  where A i s a constant f o r a system. The values of n are 6.0 f o r Al-9Mg and 2.4 f o r Al-3Mg-6Zn.  In other words  t ^ i s more stress sensitive i n the case of Al-9Mg a l l o y than i n the Al-3Mg-6Zn a l l o y . It i s important to note that the above equation holds f o r  Q  JL  1  5  10  3L  10  10'  J  2  ** No f a i l u r e i n 200 hours  10  10'  F a i l u r e time, Sec. Figure  (45).  Log-log p l o t o f a p p l i e d  s t r e s s vs f a i l u r e time f o r  Al-9Mg  i n NaCl/K CrC> 2  4  - 115 -  stresses considerably higher than the threshold stress.  Near the threshold  stress the f a i l u r e times are longer than predicted by equation (4). The power law relationship between applied stress and i n i t i a t i o n time holds f o r notched Al-21.5Zn i n 99.9% ethanol tested i n the notched form.  The value of n i s 6.3 and the correlation c o e f f i c i e n t i s 0.90.  The plot i s shown i n f i g . (47). Thus f o r Al-21.5Zn i n ethanol the values of a and t ^ can f i t  t  f  an exponential equation of the form  = B • exp (-go) • exp (2-)  (5)  It should be noted that the Al-21.5Zn i n ethanol was tested i n the notched form, while other alloys were tested i n the constant load apparatus as unnotched s t r i p s .  Both the equations are v a l i d only at  stress levels higher than the threshold.  The exponential f i t i s  marginally better i n Al-Zn while the powerlaw f i t i s better i n Al-Mg and Al-Mg-Zn.  I f Gruhl's (73) results are plotted on a log-log scale they  can f i t both equations (4) and (5) giving n values lying between 2 and 3 compared to value of 2.6 i n the present  work.  The choice of an  equation describing the stress dependence of f a i l u r e times cannot be made solely on the basis of the least squares f i t .  Rather i t should  be compatible with a t y p i c a l process taking place during stress corrosion. The crack propagation rate f o r Al-Zn i n ethanol i s also stress sensitive and can f i t either a log rate vs log a plot or a log rate vs a  plot.  The value of n obtained from a log-log plot shown i n figure  (48) i s 4.5.  1  n = 4.5  - 118 4.3.  Stress Corrosion S u s c e p t i b i l i t y of Mg-8.5A1 This system was studied because unlike the Al base a l l o y s , the  second phase i n aged Mg-8.5A1 alloys i s cathodic to the matrix by approximately 0.4 volts i n aqueous s a l t solution.  The crack path was  transgranular and normal to the d i r e c t i o n of the applied stress. i s evident from figures  This  (35,36).  When s u s c e p t i b i l i t y i s plotted against aging time, there i s a steady increase i n stress corrosion s u s c e p t i b i l i t y with aging, f i g .  (11).  Unlike A l base alloys the resistance to stress corrosion does not improve on overaging.  Between as-quenched and peak aged condition there  i s a 50% increase i n y i e l d strength and a 200 f o l d decrease i n f a i l u r e time.  After the maximum strength i s reached, overaging results i n 10%  drop i n y i e l d strength, while the f a i l u r e time decreases by a factor of 4.  Obviously, the p r e c i p i t a t i o n of M g ^ A l ^ phase leads to the  increased s u s c e p t i b i l i t y to stress corrosion cracking. According to Priest (74) , cracking i s t r a n s c r y s t a l l i n e i f the specimens are water quenched and i n t e r c r y s t a l l i n e i f they are furnace cooled.  Logan (75) observed threshold stresses between 70 and 90% f o r  annealed Mg-6.5Al-lZn a l l o y .  In the present work the threshold stress  was not obtained although i t was less than 50% of the y i e l d stress.  The  higher s u s c e p t i b i l i t y of this alloy i s probably due to the higher aluminum content.  Loose and Barbian (75) observed an improvement i n  stress corrosion resistance with a decrease i n aluminum concentration. According to Timanova (76) the nucleus of a microcrack i s created i n a Mg-8.5A1 a l l o y as a result of the selective corrosion of the intermetallic phase or of the supersaturated s o l i d solution.  The  - 119 -  stresses destroy the protective f i l m , at the crack t i p , which then corrodes at a higher rate than the crack walls. see why  It i s d i f f i c u l t to  selective corrosion should occur at the intermetallic phase which  is cathodic to the surrounding  matrix.  Heidenreich et a l (75) propose that the matrix adjacent to the i n t e r m e t a l l i c anodic to the i n t e r m e t a l l i c and goes into solution by an electrochemical process. Logan (75) observed an increase i n p o t e n t i a l i n the more active d i r e c t i o n when the stress was of the protective f i l m .  applied and associated i t with the rupturing  The specimen was  seen to be strained at a  diminishing rate a f t e r f u l l load had been applied. rate was  If the average s t r a i n  above a c r i t i c a l value, the stress concentration at few  sites  were believed to deform the specimen surface at a s u f f i c i e n t rate to prevent the f i l m from reforming.  The specimen would then f a i l i n a  short time. In the present work, a similar extension-time when notched Mg-8.5A1 alloy was  curve was  stress corroded i n the Instron.  observed The  extension was not discontinuous, as observed by Logan. Logan further proposes that i f the f i l m free area i s very small and i s surrounded by a large f i l m covered area, rapid corrosion of the anode w i l l take place.  Such corrosion w i l l form a sharp trench  preventing f i l m from reforming at i t s apex and w i l l lead to the i n i t i a t i o n of stress corrosion crack.  Cracking could be stopped by the a p p l i c a t i o n  of cathodic current and r e i n i t i a t e d i f the current was removed (74); suggesting that crack propagation was  an electrochemical process.  However,  2  Logan calculated that a current density of 14 amps/cm  would be required  - 120 to produce cracking at the rate determined i n a Mg-Al a l l o y . excessively high current density and i t i s more probable propagation  i s an electrochemical-mechanical  This i s an  that crack  process.  The fact that no discontinuous extension was observed i n load relaxation tests i n the present work, does not mean that the crack propagation i s a continuous  process.  The series of steps constituting the  fracture process may be too short i n time to be resolved by this  technique.  The increase i n stress corrosion s u s c e p t i b i l i t y with increasing aging time i s due to the reduction i n the area of the amtrix solution r e l a t i v e to the intermetallic phase M g ^ A l ^ -  solid  As a consequence,  the anodic current density increases and d i s s o l u t i o n takes place at an increased rate. Clark (61) observed planar arrays of d i s l o c a t i o n i n the solution treated condition and d i s l o c a t i o n tangles i n the oberaged condition. Planar arrays should lead to coarse s l i p and increased s u s c e p t i b i l i t y , but such i s not the case.  It appears that the role of deformation  i s to  build up stress concentration by d i s l o c a t i o n p i l e ups at the b a r r i e r s . As a r e s u l t , s l i p occurs i n grains that were previously undeformed. The protective f i l m i s ruptured and stress corrosion i s i n i t i a t e d .  The  delayed s t r a i n i n these r e l a t i v e l y resistant grains could account for the i n i t i a l decrease i n the load relaxation curve.  4.4.  Variation i n Stress Corrosion S u s c e p t i b i l i t y with the A l l o y System In this section a comparison w i l l be made between the stress  corrosion s u s c e p t i b i l i t i e s of Al-9Mg, Al-3Mg-6Zn, and Mg-8.6A1 a l l o y s .  - 121 -  It i s d i f f i c u l t to compare f a i l u r e times i n systems having widely d i f f e r e n t isothermal  aging curves.  The problem i s further complicated  by the fact that the maximum s u s c e p t i b i l i t y  to stress corrosion i s  reached before peak strength i n Al-3Mg-6Zn, at peak strength i n Al-9Mg and well after peak strength i n Mg-8.6A1.  The comparisons are made on  the basis of f a i l u r e times observed i n constant load t e s t s .  In a l l  cases, a 3.5% NaCl, + 2.0% K^CrO^ solution was used. At 40% of the y i e l d stress the solution treated and quenched Al-9Mg i s immune.  From e x t r a p o l a t i o n of the l i n e i n f i g . (12), i t  appears that Mg-8.6A1 w i l l also be immune at 40% of the y i e l d stress. Only Al-3Mg-6Zn i s susceptible and f a i l s i n 50 hours. During preaging Al-9Mg i s the most susceptible, Al-3Mg-6Zn i s next and Mg-8.6A1 i s s t i l l immune. In the peak aged condition Al-9Mg f a i l s i n 1/2000 of the f a i l u r e time of Al-3Mg-6Zn.  Mg-8.6A1 i s s t i l l immune.  In the overaged condition the order of s u s c e p t i b i l i t y  remains  the same as before, but the differences i n f a i l u r e times are reduced. There i s only a two-fold difference i n s u s c e p t i b i l i t y  of Al-9Mg and  Al-3Mg-6Zn, while Mg-8.6A1 i s expected to be immune. If f a i l u r e times are compared at 40% of the y i e l d stress f o r the most susceptible condition of each system, i t i s found that Al-9Mg is  200 fold more susceptible than Al-3Mg-6Zn and nearly 7000 f o l d more  susceptible than Mg-8.6A1 a l l o y . In figure (49), the stress vs log f a i l u r e time for Al-9Mg, Al-3Mg-6Zn and Mg-8.6A1 i s plotted.  It w i l l be noticed that Mg-8.6A1  and Al-9Mg show f a i l u r e times which are highly stress sensitive while  Failure time seconds Figure (49). Applied stress vs f a i l u r e time for three systems studied under constant in NaCl/K Cr0 2  4  load,  - 123 -  f a i l u r e times i n Al-3Mg-6Zn are comparatively the applied stress.  i n s e n s i t i v e to changes i n  The reason why Mg-8.6A1 appears to be immune at 40%  of the y i e l d stress, whereas Al-3Mg-6Zn i s s t i l l subject to stress corrosion cracking, i s that the stress level at 40% y i e l d stress f o r Mg-8.6A1 l i e s below the cross over of the two curves.  Therefore,, the  r e l a t i v e s u s c e p t i b i l i t i e s to stress corrosion depend on the f r a c t i o n of the y i e l d stress at which comparisons are made.  For example at 90%  of the y i e l d stress Al-3Mg-6Zn i s the least susceptible.  If failure  times are compared at the same absolute stress l e v e l , the stronger a l l o y would appear to have longer f a i l u r e time above the cross over point, because that stress level would correspond to a r e l a t i v e l y lower f r a c t i o n of i t s  y i e l d stress.  Then again, comparisons are d i f f i c u l t to  make at stress levels between the y i e l d stresses of two a l l o y s .  For  instance at 30,000 p s i , the y i e l d stresses of Al-9Mg and Mg-8.6A1 are exceeded and their s u s c e p t i b i l i t i e s cannot be assessed r e l a t i v e to Al-3Mg6Zn or to each other.  Therefore  i t seems that the best approach i s  to make comparisons of stress corrosion behaviour at a constant f r a c t i o n of the y i e l d stress.  In this work the fractions of y i e l d stress were  40% for Al-9Mg and Al-3Mg-6Zn and 70% for Mg-8.6A1 i n t h e i r aged conditions.  4.5.1.  I n i t i a t i o n and Propagation of Stress Corrosion Cracks i n NaCl/ K Cr0 2  4  Solution  The i n i t i a t i o n time was defined as the elapsed time between the application of f u l l load and the f i r s t detectable drop i n the load. It i s possible that cracks could have been i n i t i a t e d before the load drop  - 124 -  was noticeable.  However, specimens unloaded before the anticipated  i n i t i a t i o n time, f a i l e d to reveal any cracks.  Allowing f o r a factor  of two variations i n the anticipated i n i t i a t i o n time, i t was observed that cracks were seen only after the drop i n the load.  Initiation  times were generally reproducible within a factor of two. Three types of load-relaxation curves were observed as shown in f i g .  (50). The solution heat treated Al-3Mg-6Zn exhibited the type  of curve shown i n f i g .  (50a) at stresses up to 55,000 p s i .  In this  the load relaxation occurs i n small steps down to zero load.  This  type of behaviour suggests that crack propagation may be a discontinuous process, so that each step represents an electrochemical-mechanical advance of the stress corrosion crack.  However, aluminum alloys show  Portevin-Le Chatelier effect due to the pinning and unpinning of d i s l o c a tions by solute atoms, and this effect produces a discontinuous y i e l d i n g . During the notch t e n s i l e tests the load f a l l s i n a discontinuous manner much l i k e i n stress corrosion tests. 1.  This gives r i s e to two p o s s i b i l i t i e s .  Stress corrosion crack propagation i s continuous but the  p l a s t i c flow at the crack t i p i s discontinuous as a result of locking and unlocking of d i s l o c a t i o n s . 2.  Both crack propagation and p l a s t i c flow are discontinuous,  but only the discontinuity of the p l a s t i c flow i s detectable by the present technique.  The second type of curve i s shown i n f i g .  the stress corrosion crack propagates specimen showing steps.  part of the way through the  Then at D, the stress at the t i p of the crack  exceeds the ultimate strength of the a l l o y and mechanical occurs.  (50b);  fracture  This type of behaviour was observed i n most aged alloys of  - 125 -  Figure (50).  Three types of load relaxation curves observed i n the present work.  - 126 -  aluminum.  A t h i r d type of load-relaxation curve i s shown i n f i g .  (49c).  This behaviour was noticed i n Mg-8.6A1 alloys under a l l conditions. The crack propagation i s e n t i r e l y continuous and no steps are detected i n the load-relaxation curve. Thus, solution treated Al-3Mg-6Zn showed the f i r s t type of curve at low stresses, Al-2.6Mg-6.3Zn and Al-8.6Mg showed the second type of curve (49b) i n the aged condition; Al-21.5Zn and Mg-8.6A1 showed the t h i r d type of curve although the former a l l o y exhibited small discontinuities during the propagation of the stress corrosion crack. The crack propagation v e l o c i t i e s were calculated for alloys in which the stress corrosion crack propagated through the entire cross section.  This was done by dividing the radius of the notch by  the propagation time as measured on the load-relaxation curve.  As  mentioned e a r l i e r , through stress corrosion cracks were observed i n r e l a t i v e l y few specimens and therefore the propagation v e l o c i t i e s are approximations at best. quenched Al-Mg-Zn i s  The propagation rate i n s a l t solution f o r as1 mm/hr at 55,000 p s i ; f o r aged Al-Zn i t i s 0.8  mm/hr at 10,000 p s i ; for aged Al-Mg i t i s  4 mm/hr at 30,000 p s i . The  as-quenched Mg-Al showed a crack propagation rate of 4.8 mm/hr at 26,000 psi  i n a s a l t solution ten times more d i l u t e than the one used for  aluminum base a l l o y s .  Pugh observed a propagation rate of 0.35 mm/hr  in an Al-Mg-Zn alloy tested i n 3.5% NaCl solution at 40,000.  The higher  rate o f crack propagation i n the present work i s probably caused by the high t r i a x i a l stress state at the root of the notch.  The rate of crack  propagation i s influenced by the i n i t i a l applied stress. I  n  the Al-Zn  alloy i n alcohol at room temperature the rate of propagation i s proportional  - 127 to the logarithm of applied stress according to expression of the form  n  rate «  a  where n - 4.5 These rates were calculated assuming that crack progresses a distance L over a c i r c u l a r front u n t i l the stress on the remaining area exceeds the notch t e n s i l e strength. Therefore i f R = radius of the notch P = load at fracture p  then stress on the remaining area = TT  Equating this to the notch t e n s i l e strength  u(R-L)  Rearranging  knowing  2  ~  °  — (R-L) we get  n  P L = R- y — —  (6)  the cracklength' and propagation time, the propagation rate can  be calculated. Because the radius of curvature of the machined notch i s considerably greater than that of the crack, the above expression would over estimate the crack length.  During crack propagation, the  load f a l l s but the stress concentration at the crack t i p increases because of a reduction i n the load bearing area.  The behaviour of  s u b - c r i t i c a l crack propagation i s i l l u s t r a t e d i n figure (50) where the stress i n t e n s i t y factor time (77).  i s plotted against crack v e l o c i t y and fracture  - 128 -  Figure  (51).  V a r i a t i o n i n s t r e s s i n t e n s i t y f a c t o r Kj d u r i n g c r i t i c a l c r a c k p r o p a g a t i o n (77).  sub-  - 129 Kj =  where  a/~~c~ x constant  a = applied stress c = i n i t i a l crack length  Stress corrosion cracks can propagate above a minimum threshold of stress intensity factor = K^g^. ^ISCC ^  S  e x c e e  d e d u n t i l near  mechanical fracture occurs.  The v e l o c i t y i s nearly constant once where i t becomes i n d e f i n i t e l y fast and  The shape of this curve resembles  load relaxation curves shown i n f i g .  (49b,c).  constant for a given metal-environment system.  the  Both K j g ^ and  are  The level of crack  v e l o c i t y i s several orders of magnitude higher than can be inferred from models i n which the crack  dissolves or i s the s i t e of  alternate formation and rupture of the oxide f i l m (78.79).  According  to Pugh, (26) crack v e l o c i t i e s of 0.35 mm/hr observed i n an Al-Mg-Zn 2 alloy correspond to current densities of 0.3 amp/cm  and these are not  considered to be unreasonably high for a corrosion process. 4.5.2.  The Relative Lengths of I n i t i a t i o n and Propagation Times i n NaCl/K Cr0 Solution 2  4  In Al-Mg-Zn a l l o y i n i t a t i o n times i s greater than propagation time except for the large grained a l l o y .  From Table (17) i t i s seen  that lowering the stress results i n an increase i n the f r a c t i o n of propagation time f o r the same heat treatment.  The heat treatment given  to a specimen remarkably alters the f r a c t i o n of propagation time r e l a t i v e to i n i t i a t i o n time.  Propagation time comprises  the greater part of  the t o t a l f a i l u r e time i n the large grained (570 p) specimens.  In the  - 130 solution treated specimens of smaller grain size (256  u ) propagation  time i s 12 to 50% of the t o t a l time, whereas i n the overaged i t constitutes only 3-5% of the t o t a l time.  specimens  Therefore the duration of  crack propagation i s favored by low stress, large grain size and a lack of p r e c i p i t a t e . In Al-8.6Mg a l l o y the propagation time consists of 92-99% of the t o t a l time for underaged alloys and only 5-16% f o r peak aged a l l o y s . The large grained specimens also show propagation times between 95-97% of the t o t a l time.  In aged Al-21.5Zn alloythere i s no d e f i n i t e effect  of stress level on the propagation part of stress corrosion, although i t consitutes the major part of the f a i l u r e time.  The propagation  part varies from 65 to 95% of the t o t a l f a i l u r e time. In solution treated and quenched Mg-8.6A1 a l l o y the propagation time accounts f o r 74 to 96% of the t o t a l time.  D i l u t i o n of the solution  results i n an increase i n the f r a c t i o n of propagation time while varying the stress l e v e l does not have any e f f e c t .  The a l l o y aged for a  short time (1 hour at 200°C), shows a s l i g h t decrease i n the proportion of propagation time to i n i t i a t i o n time.  4.6.  The Effect of Different Environments on the Stress Corrosion Behaviour of Aluminum Alloys Aluminum alloys were found to stress corrode i n a number of  environments as seen i n Table (9-10) and figures (15) and (16).  The  total f a i l u r e time decreased as the amount of water i n the environment increased.  Both aqueous and organic liquids caused cracking provided  the applied stresses were high enough.  The purpose of t h i s set of  - 131 -  Table (14) RELATIVE FRACTIONS OF INITIATION AND PROPAGATION TIMES IN NaCl/K Cr0 2  Alloy Heat Treatment  Al-3Mg-6Zn as-quenched  Applied Stress psi  4  SOLUTION  % Initiation time  % Propagation time  48,000 55,000 55,000 55,000 69,000 69,000  88 50 95 97.5 96 95.5  12 50 5 2.5 4 4.5  as-quenched, large grain size  46,000 46,000 .  33 41  67 59  Al-2.6Mg-6.3Zn, as-quenched  57,600 57,600  71 67  29 33  Al-8.6Mg, underaged  40,800 38,200  8 0.6  92 99.4  peak aged  12,500 31,200  95 94.5  5 5.5  Al-8.6Mg, peak aged, large grain s i z e  31,000 27,000  2.6 4.5  97.4 95.5  Al-21.5Zn, peak aged  40,000 26,000 26,000 10,000  5 17 35 30  95 83 65 70  26,000 26,000  20 17  80 83  II  over-aged II  n n  11  II  n  Mg-8.6A1, as-quenched II  n  it II  Aged 200°C, 1 hour  26,000 26,000  3.5 6.5  Remarks  Discontinuous load relaxation in Al-Mg-Zn system.  Discontinuous load relaxation in Al-8.6Mg system  Continuous load relaxation in Al-21.5Zn  96.5 93.5  solution dilutee by 10:1 Continuous load relaxation i n Mg-8.6A1  30,000 30,000  20 18  80 82  30,000  26  74  - 132 Table (15) RELATIVE LENGTHS OF INITIATION AND PROPAGATION; TIMES IN VARIOUS ALLOY-ENVIRONMENT SYSTEMS Alloy, heat treatment  Environment  Applied stress psi  % Initiat i o n time  % Prop time  Al-2.6Mg-6.3Zn Solution treat and quenched  3.5% NaCl + 2.0% K.CrO. L 4 D. water Ambient a i r  57,600 57,600  71 67 •  29 33  57,600 57,600 57,600 57,600 57,600  60 74 60 77 89.5  40 26 40 23 10.5  60,000 60,000 60,000 68,000 68,000  83 68 42.5 30 36  17 32 47.5 70 64  83,000 83,000 83,000 83,000 76,500 86,000 86,000 40,000 40,000 64,600 64,600  3 7 87 85 98 88.5 81 87.5 78 0 75  97 93 13 15 2 11.5 19 12.5 22 100 25  64,600  69  31  64,600 64,600 64,600 64,600 64,600  79 67.5 85.5 75 87.5  21 32.5 14.5 25 12.5  64,600 64,600 64,600 58,000 58,000 58,000 58,000  58 53 87.5 87 95 71 82  42 47 12.5 13 5 29 18  II  f  II  S i l i c a Gel  ti  Methyl alcohol  II  CC1  II  CCI4 CCI4  n  Overaged  4  + D. water + D. water  Ambient a i r  n  D. wat er  11  Ethyl alcohol  Al-9Mg  Ethyl alcohol  Al-21.5Zn peak aged  D. water Ambient a i r (760 mm) Ambient a i r (20 u ) S i l i c a gel Mg(0Cl ) 4  2  Sodium dried, Kerosene, 1 M, NaOH Ethyl alcohol, 99.9% Ethyl alcohol 95% -  - 133 experiments, was to establish the effect of d i f f e r e n t environments on the i n i t i a t i o n and propagation of stress corrosion cracks.  Preliminary  tests were carried out on solution treated Al-2.6Mg-6.3Zn but later work was confined to the peak aged and overaged conditions. necessary  This was  i n view of the fact that supersaturated aluminum alloys  undergo p r e c i p i t a t i o n during p l a s t i c deformation.  Incipient p r e c i p i t a -  tion at the t i p of a stress corrosion crack i s l i k e l y to change the rate of chemical or electrochemical attack as well as mechanical properties of the a l l o y .  In the peak aged or overaged condition the  above processes would not occur as most of the solute would be present as p r e c i p i t a t e and not i n the s o l i d solution. As the amount of water i n the environment i s reduced, the f a i l u r e time increases s i g n i f i c a n t l y for Al-2.6Mg-6.3Zn a l l o y i n both the solution treated and overaged conditions.  This can be seen from  figure (15) . The f a i l u r e times increase by factors of 6 to 8 when the environment i s changed from deionized water to ambient a i r . A further 50 f o l d increase takes place when a dessicant such as s i l i c a gel i s used.  This effect shows the importance of water vapour i n the stress  corrosion of aluminum alloys provided a s u f f i c i e n t stress concentration exists at the root of the notch. also i l l u s t r a t e d by  The importance of water vapour i s  figures (13) and (14) .  A s i g n i f i c a n t increase  in elongation i s observed when the regular environment i s replaced by a dessicant. In the as-quenched a l l o y , propagation between 10 to 40% of the t o t a l f a i l u r e time.  times constitutes  In the overaged a l l o y ,  most of the f a i l u r e time i n ambient a i r i s taken up by the propagation of the stress corrosion crack, whereas i n deionized water, i n i t i a t i o n  - 134 time seems to constitute the major part of the time to f a i l u r e . The load relaxation curve i s discontinuous up to a certain point beyond which mechanical fracture takes place. Organic liquids such as methyl alcohol, ethyl alcohol, and carbon tetrachloride also caused stress corrosion cracking i n aluminum alloys.  The cracking times were an order of magnitude longer i n methyl  alcohol than i n water.  The major part of stress corrosion f a i l u r e was  s t i l l taken up by the crack i n i t i a t i o n stage.  In carbon tetrachloride  saturated with water, f a i l u r e occurred i n about the same time as i n water free carbon tetrachloride. 1.  There are two p o s s i b i l i t i e s .  The cracking was caused by the minute quantity of water  present i n the organic l i q u i d and further increase i n water concentration did not have a s i g n i f i c a n t effect on the stress corrosion rate. Hence similar times were observed for water saturated and water free carbon tetrachloride. 2.  Water played no part i n the process and therefore no  difference was detected i n f a i l u r e times. A further experiment was carried out to resolve the r o l e of water i n the process.  The c e l l containing carbon tetrachloride was  surrounded by a polythene bag containing a dessicant.  Therefore the  organic l i q u i d was i n equilibrium with dry a i r rather than ambient a i r . The specimen did not f a i l i n 100 hours whereas previously i t f a i l e d i n less than two hours.  This demonstrates that water plays an important  role i n the process. In ethyl alcohol the propagation time was less than twenty percent of the t o t a l time to f a i l u r e both i n the Al-Mg-Zn and Al-Mg a l l o y .  - 135 Al-21.5Zn a l l o y was  extremely susceptible to stress corrosion  and f a i l e d 350 times faster than aged Al-2.6Mg-6.3Zn i n water; at the same f r a c t i o n of t h e i r respective notch strengths. in s i l i c a gel was  also 1000  times shorter.  The f a i l u r e time  Different media were t r i e d  in an attempt to prolong the l i f e of specimens.  These included  magnesium perchlorate, sodium dried kerosene, and sodium hydroxide and a i r at reduced pressure.  However, the l i f e times of specimens could  not be prolonged beyond f i f t e e n minutes.  The f a i l u r e times were  increased 35 to 70 f o l d compared to deionized water, by using dessicants.  various  This compares with an increase of 80 times i n the overaged  Al-2.6Mg-6.3Zn a l l o y . Except i n deionized water, the major part of f a i l u r e times consisted of the i n i t i a t i o n of the stress corrosion crack.  The  i n i t i a t i o n stage constituted between 65 to 95% of the t o t a l time for d i f f e r e n t dessicants and organic l i q u i d s . i n i t i a t i o n was crack was  almost instantaneous.  This suggests that perhaps the  i n i t i a t e d during the loading of the specimen. The load relaxation curve was  The  load drop was  fracture.  However, i n water the  continuous as i n figure (46c) .  15 to 20% of the maximum load, followed by mechanical  The stress corrosion as well as mechanical fracture was  intercrystalline.  The fracture surface was bright i n aqueous mediums  and d u l l i n ethyl alcohol. Aluminum a l l o y s  can react with alcohol and produce hydrogen  once the protective oxide f i l m i s ruptured.  This was  demonstrated by  the following experiment. A piece of aluminum was  rubbed on coarse emery paper and  - 136 -  immediately dipped i n a mercurous chloride solution. aluminum amalgam was  formed wherever the oxide f i l m was broken and  prevented the repair of the f i l m . alcohol.  A mercury-  Then the specimen was  A combustible gas, presumably hydrogen, was  placed i n ethyl  seen to evolve  from the specimen as a r e s u l t of the action of fresh aluminum surface with the alcohol. and water.  A similar effect was  noticed i n carbon tetrachloride  The evolution of the gas was most rapid i n water.  This  shows that without i t s protective oxide f i l m , aluminum surface i s highly reactive in| ynjeous and organic l i q u i d s .  It i s therefore not surprising  that traces of water i n the a i r , or i n dessicants are s u f f i c i e n t to cause stress corrosion cracking i n the  4.6.1.  aluminum a l l o y s .  Stress Corrosion i n Alcohols Stress corrosion cracking i n ethyl alcohol can be caused either  by the trace of water present i n i t or by a reaction between the a l l o y and ethyl alcohol at points where the protective oxide f i l m has been ruptured. A l l three aluminum base a l l o y s were susceptible to stress corrosion i n 99.4%  ethyl alcohol.  Al-2.6Mg-6.3Zn a l l o y was  also  susceptible i n methyl alcohol as well as i n carbon tetrachloride as seen i n f i g . (15). The  importance of residual water i n determining the rate of  cracking i s i l l u s t r a t e d in f i g . (16).  For the same stress l e v e l , the  rate of i n i t i a t i o n i s between 6 to 8 times faster i n 95% ethyl than i n 99.9%  ethyl alcohol.  alcohol  However,the apparent activation energy of  the process of stress corrosion cracking  i n alcohol containing  small  amounts of water, i s substantially d i f f e r e n t from the apparent a c t i v a t i o n  - 137 -  energy of the process responsible f o r stress corrosion i n d i s t i l l e d water,  This i s evident from figures (20) and (21) . The a c t i v a t i o n  energy i n ethyl alcohol i s 14.7 kcals/gm mole whereas i n d i s t i l l e d water and NaCl/K^CrO^ solution, i t i s 26 kcals/gm mole.  Therefore i t  appears that while addition of water enhances the rate of crack i n i t i a t i o n i t does not change the apparent activation energy of the process responsible for cracking i n alcohol.  Furthermore, the apparent  activation energy of this process i s between 9.6 and 14.7 kcals/gm mole i n the three aluminum base a l l o y s , compared to an apparent activation energy of 26 kcals/gm mole i n d i s t i l l e d water and NaCl/ K^CrO^.  The nature of the possible processes consistent with the above  activation energies w i l l be considered l a t e r .  4.6.2.  Effect of Environment on the Mechanical Properties of Al-2.6Mg6.3Zn Environments such as deionized water and NaCl/K^CrO^ solution  affected the t e n s i l e d u c t i l i t y of the heat treated Al-Mg-Zn and Al-Zn a l l o y s , but had no s i g n i f i c a n t effect on the y i e l d stress or ultimate t e n s i l e strength of the Al-Mg-Zn a l l o y .  Both environments markedly  reduced the uniform p l a s t i c elongation as shown i n f i g s . (13,14) and table (7).  D'Antonio et a l (101) observed that the flow stress of  Al-Mg-Zn alloys was s t r a i n rate insensitive at room temperature between s t r a i n rates employed i n the present work.  They concluded that  discontinuous flow should be observed at slow s t r a i n rates at room temperature.  Discontinuous flow was observed both i n deionized water  - 138 -4 and NaCl/K^CrO^ up to s t r a i n rates of 2 x 10  -1 sec  i n the present  work. In the solution treated Al-2.6Mg-6.3Zn a l l o y the uniform p l a s t i c elongation increases with increasing s t r a i n rate i n both deionized water and NaCl/K^CrO^.  The elongation i n deionized water  is always greater than that i n NaCl/K^CrO^.  The elongation increases  by a factor of 9 as the s t r a i n rate i s increased from 4 x 10 ^ sec ^ -3 -1 to 2 x 10  sec The low d u c t i l i t y at low s t r a i n rates i s a result of long  exposure to the environment.  As the s t r a i n rate goes up, the time of  exposure under stress goes down.  Consequently, the embrittling  effect of the environment i s reduced and the d u c t i l i t y increases.  At  very high s t r a i n rates the environment is p r a c t i c a l l y i n e f f e c t i v e i n reducing the d u c t i l i t y and the l i m i t of elongation i s set by the mechanical properties of the a l l o y .  Therefore the elongation at the  highest s t r a i n rate i n both deionized water and NaCl/K^CrO^ i s the same as elongation i n ' a i r at a lower s t r a i n rate; which i n turn i s nearly the same as elongation i n dessicant dried a i r at even a lower s t r a i n rate. Therefore more aggressive environments require high s t r a i n rates i n order to produce the same elongation as i n the less  aggressive  environments. When a dry a i r ( s i l i c a gel) environment i s used at a s t r a i n rate of 2 x 10 ^ sec  the uniform elongation increases by more than  a factor of 7 while the flow stress remains unchanged compared to t h e i r respective magnitudes i n the more aggressive environments. The change i n elongation, due to the change i n s t r a i n rate  - 139 -  alone, was about 20% of the t o t a l difference i n elongation, between —6 -1 3 1 s t r a i n rates of 2 x 10 sec and 2 x 10 sec . The environment was _  d i r e c t l y responsible for the remaining 80% increase or decrease i n the uniform p l a s t i c elongation. Similarly i n the overaged Al-Mg-Zn a l l o y , both uniform elongation and d u c t i l i t y increase with increasing s t r a i n rate. The uniform elongation i n s i l i c a gel-dried a i r was 4 times that i n NaCl/^CrO^ and deionized water at comparable s t r a i n rates f o r the overaged Al-Mg-Zn.  In a l l environments at lower s t r a i n rates the  load relaxation was discontinuous, or steplike.  This suggests that  the steps observed i n stress corrosion tests were also probably the result of discontinuous y i e l d i n g of the a l l o y at low s t r a i n rates, rather than caused by the arrest and propagation of the stress corrosion crack. 10  Discontinuous y i e l d i n g was observed up to s t r a i n rates of 2 x  - 4 - 1 sec  . Therefore, i t i s reasonable to i n f e r , that when steps are  observed i n a stress corrosion test, the crack propagates at a s t r a i n - 4 - 1 rate of not more than 2 x 10  sec  , whereas when no steps are observed  the crack propagation takes place at a higher s t r a i n rate. In the overaged a l l o y discontinuous y i e l d i n g took place at a lower s t r a i n rate  10  3  sec ^.  There are two  reasons why steps  are not observed i n the stress corrosion test of overaged Al-2.6Mg-6.3Zn alloy under s t a t i c load. 1.  The uniform elongation or d u c t i l i t y of the overaged a l l o y  is less than 50% of the d u c t i l i t y of the solution treated a l l o y . Consequently the region of crack propagation i n solution treated a l l o y 'is larger than i n the overaged a l l o y as i s the p r o b a b i l i t y of observing a stepwise decrease i n the load.  There i s also less solute i n the overaged  - 140 -  2.  S t r a i n rates at which discontinuous y i e l d i n g i s observed  are one order of magnitude less i n the overaged alloys compared to the solution treated a l l o y s .  In the environments such as deionized  water and NaCl/K^CrO^ the crack propagation results i n a s t r a i n rate higher than 10 ^ sec * and no steps are observed. The extent of step formation can be increased by changing from an aggressive environment to a less aggressive environment.  This  effect was noticed i n solution treated Al-2.6Mg-6.3Zn a l l o y , when dessicant dried a i r was used instead of deionized water or ambient air.  Stress corrosion crack propagated through the specimen and the  load relaxation was discontinuous  i n dry a i r , while i n other environments  the crack propagated p a r t i a l l y through the specimen before mechanical fracture due to t e n s i l e overload  4.6.3.  occurred.  Effect of Environment and Strain Rate on Mechanical Properties of Al-21.5Zn Deionized water d r a s t i c a l l y reduces the d u c t i l i t y of aged  Al-21.5Zn a l l o y .  The uniform elongation can be increased 16 f o l d and  the t o t a l elongation three times when deionized water i s replaced by dessicant dried a i r , at s t r a i n rates of 6 x 10 ^ sec  The y i e l d  stress i s increased by a factor of 2.5, table (7). -3 At s t r a i n rates of 2.7 x 10  -1 sec  there i s a 3 fold increase  in uniform elongation on using dry a i r i n place of deionized water and no s i g n i f i c a n t change i n the y i e l d stress. On increasing the s t r a i n rate from 6 x 10 ^ sec ^ to 2.7 x 10  - 3 - 1 sec  the elongation i n deionized water increases 20 f o l d and that  - 141 in s i l i c a gel 4 f o l d . Al-21.5Zn stress corrodes even i n dessicant dried a i r and that i s the reason for observing the l a t t e r increase i n uniform tion.  elonga-  The values of elongation and y i e l d strength depend on the time  the a l l o y has been exposed to the environment.  These times are small  at higher s t r a i n rates and therefore the y i e l d strength or elongation is not affected s i g n i f i c a n t l y .  At low s t r a i n rates the y i e l d strength  as well as elongation i s reduced i n the dessicant dried a i r and deionized water.  This i s so, because s u f f i c i e n t time i s available for stress  corrosion cracks to be nucleated when the s t r a i n rate i s low. t r u l y inert environment was  4.6.4.  No  found for Al-21.5Zn.  Effect of Strain Rate on Serrated Yielding Serrated y i e l d i n g was most often observed i n the solution  treated Al-Mg-Zn a l l o y and was  altogether absent i n the Al-Zn a l l o y .  In the solution treated a l l o y serrated y i e l d i n g was observed up to a - 4 - 1 s t r a i n rate of 2 x 10  sec  , whereas i n the overaged a l l o y i t was  present only below a s t r a i n rate of 10 ^ sec  1  .  The reason serrations are present only below a c r i t i c a l  strain  rate i s that at low s t r a i n rates i t i s easier for solute atoms present in the solution to segregate to d i s l o c a t i o n s .  In addition the  tion of solute atoms is higher i n the supersaturated in the overaged a l l o y .  concentra-  s o l i d solution than  As a result longer d i f f u s i o n times are required  for a s u f f i c i e n t number of solute atoms to segregate to a d i s l o c a t i o n and cause discontinuous y i e l d i n g i n the case of the overaged a l l o y . Therefore i n an overaged a l l o y , discontinuous y i e l d i n g can be observed  - 142 only at s t r a i n rates lower than i n the solution treated a l l o y s . No serrations are observed i n the overaged Al-21.5Zn a l l o y because i t i s highly b r i t t l e . When pure aluminum (99.995) was tested i n two d i f f e r e n t environments at s t r a i n rates of 2 x 10 ^ sec ence was observed. water.  1  no s i g n i f i c a n t d i f f e r -  The environments were s i l i c a gel and deionized  S i m i l a r l y at higher s t r a i n rates there was no change i n y i e l d  strength or uniform elongation.  This i s so, because pure aluminum i s  not susceptible to stress corrosion.  Thomas (81) reported that pure  (five9's) aluminum was susceptible to stress corrosion i f quenched from an annealing temperature of 600°C.  However, i n the present  work both smooth and notched specimens of 99.99% aluminum were found to be immune i n deionized water and NaCl/K^CrO^ solution. Tensile tests carried out using  water as the heating medium,  under the assumption that water i s an inert medium can lead to erroneous r e s u l t s .  This i s p a r t i c u l a r l y true i n the case of aged -4  Al-21.5Zn and Al-2.6Mg-6.3Zn alloys at s t r a i n rates less than 10  -1 sec  The same results are to be expected i n the aged Al-8.6Mg and Mg-8.6A1 alloys.  Uniform and t o t a l elongations are the properties most affected  by water and NaCl/K^CrO^.  Y i e l d strength and work hardening rate are  not changed at d i f f e r e n t s t r a i n rates (80) except i n the highly suscept i b l e Al-21.5Zn a l l o y . The change i n uniform elongation f o r a given a l l o y , due to the change i n environment depends on several factors. (i) environments.  These are:  The r e l a t i v e s u s c e p t i b i l i t y of the a l l o y i n the two Higher the difference i n the s u s c e p t i b i l i t i e s i n the two  - 143 -  environments greater the difference i n uniform elongation.  For example,  in the solution treated Al-2.6Mg-6.3Zn a l l o y , the stress corrosion s u s c e p t i b i l i t y i n deionized water and NaCl/K^CrO^ i s almost the same and so i s the uniform elongation. i s 500 times that i n NaCl/K Cr0 2  4  But the f a i l u r e time i n s i l i c a gel  and the uniform elongation i s 7 f o l d .  Similar results are observed f o r Al-21.5Zn a l l o y . (ii)  The aging treatment of the a l l o y .  The absolute value  of the uniform elongation i s decreased i n the overaged Al-2.6Mg-6.3Zn alloy.  The s t r a i n rate at which environment induced  embrittlement  occurs i s reduced by at least an order of magnitude r e l a t i v e to the solution treated a l l o y . ( i i i ) The value of the s t r a i n rate. occurs at the lowest s t r a i n rate.  Maximum  embrittlement  Very low s t r a i n rate tests approximate  a constant load stress corrosion test and therefore produce the maximum embrittlement or the minimum uniform elongation.  4.7.  Temperature Dependence of the Stress Corrosion Process The stress corrosion cracking of aluminum alloys tested i n  alcohol, d i s t i l l e d water and NaCl/K^CrO^ solution, i s a thermally activated process.  The apparent activation energy of the rate control-  l i n g process varies with the alloy-environment system.  It also changes  during i n i t i a t i o n and propagation stages. The temperature  dependence of f a i l u r e times of Al-Mg-Zn alloys  in s a l t solutions has also been studied by Gruhl (73) and H e l f r i c h (72). Gruhl obtains a value of 12.6 kcals/gm mole at 75% of the y i e l d stress.  - 144 -  At a  stress level close to the y i e l d strength the Arrhenius plot  i s not a straight l i n e , but the gradient increases between 50 and 70°C. The apparent a c t i v a t i o n energy between these two temperatures i s 22.2 kcals/gm mole.  At other stress levels there i s no s i g n i f i c a n t  change i n the a c t i v a t i o n energy.  H e l f r i c h (72) obtains activation  energy values between 17 and 20 kcals/mole.  The observed a c t i v -  ation energy increases with decreasing applied stress and has a value of 20 kcals/mole i n the absence of stress. Before comparing  the results of the present work with  those of Gruhl and H e l f r i c h , i t should be remembered that their composition of alloys as well as testing techniques were d i f f e r e n t . In the present work Al-2.6Mg-6.3Zn was tested by continuous immersion of notched c y l i n d r i c a l specimens i n 3.5NaCl/2.OK^CrO^ solution.  H e l f r i c h used C rings of Al-4.2Zn-5Mg immersed i n  aerated 5.8 %NaCl solution.  Gruhl tested an Al-4.9Zn-3Mg a l l o y i n  a 3 %NaCl solution which was circulated around the specimen. A d i s t i n c t i o n was made between i n i t i a t i o n and propogation times in the present work, whereas t o t a l times to f a i l u r e were measured in theother two works.  The c r i t e r i o n for f a i l u r e i n Helfrich's  work (72) was the f i r s t v i s i b l e evidence of cracking at f i v e diameters magnification. Not a l l the aluminium alloys are s i m i l a r l y affected by temperature. The stress corrosion f a i l u r e times of agehardened 2024 aluminum alloy were found to be independent of testing temperature between 3 and 60° by Rostoker and Nicholas (84).  - 145 Wei  (82) studied moisture enhanced faituge crack propagation  in a high strength aluminum a l l o y and "interpreted" i t as a thermally activated process with apparent activation energies that depend strongly on the crack t i p stress i n t e n s i t y factor and vary i n magnitude from 1-4 kcals/gm mole. It should be emphasized that the value of apparent a c t i v a t i o n energy determined by experiment can correspond to a number of rate controlling processes.  Therefore additional information  to eliminate some of the processes with overlapping  i s necessary  a c t i v a t i o n energies.  In addition the values of the apparent a c t i v a t i o n energies of some of the possible processes are not available. w i l l be discussed  Several possible models  i n the light of the above r e s u l t s .  The values of  apparent a c t i v a t i o n energies of some of the possible rate c o n t r o l l i n g processes are l i s t e d i n table (16).  4.8.  The Hydrogen Mechanism The essential steps i n the proposed hydrogen mechanism of stress  corrosion are as follows: 1.  The f i r s t step i n the i n i t i a t i o n of a stress corrosion  crack i s the electrochemical boundary p r e c i p i t a t e .  d i s s o l u t i o n of the equilibrium grain  The attack takes place at the partial-matrix  interface because the protective oxide f i l m i s broken at the interface. H e l f r i c h (83) observed d i s s o l u t i o n at the constituent particle-matrix interface and Jacobes (20) observed d i s s o l u t i o n of the equilibrium precipitate.  However, Hunter (12) found that the MgZn  p a r t i c l e s were  - 146 Table (16) THE APPARENT ACTIVATION ENERGIES FOR CRACK INITIATION AND PROPAGATION AND THE POSSIBLE RATE CONTROLLING PROCESSES Observed Q  26  Alloy-Env. system  Initiation • or Propagation  Possible Mechanism  Al-Mg-Zn of s a l t solution  Initiation  1. Electrochemical d i s s o l u tion of MgZn or aluminum. Q = up to 25 kcals/mole (85) 2  2. Creep process controlled by cross s l i p . Q = up to 27.5 kcals/mole (97) 12 (9.6-14.7)  1. Al-Zn,Al-Mg, Al-Mg-Zn i n alcohol  Initiation  1. Diffusion of hydrogen through aluminum s o l i d solution Q = 10.9 kcal/mole (97)  2. Al-Mg-Zn i n s a l t solution  Propagation  2. Creep process controlled by intersection of d i s l o c a tions Q = 11 ± 8 kcals/mole (89) 3. Movement of single vacancies through aluminum s o l i d solution Q = 12 kcals/mole (73) 4. Escape of d i s l o c a t i o n from segregated atmosphere (108)  1-4  Al-Zn, Al-Mg,  1. Transport processes i n the l i q u i d Q = 2-6 (101) .  - 147 -  l e f t unattacked by an a c i d i f i e d NaCl/AlCl^ solution whereas the grain boundary was attacked.  It i s possible that the nature of attack  depends on the pH of the solution.  In fact H e l f r i c h (83) observed that  when Al-Mg-Cu and Al-Cu alloys were stress corroded i n an acid (pH = 0.5) salt-dichromate solution cracking i n i t i a t e d at the grain boundaries. But when these alloys were stress corroded i n a solution of pH = 4 the electrochemical attack started at the constituent p a r t i c l e s .  Therefore,  i t appears that at acid pH's the protective oxide f i l m i s destroyed and fresh aluminum i s exposed which i s attacked i n preference to MgZn  2>  From the potential -pH diagram of Al-H^O system i t i s seen that the oxide i s stable only between pH's of 4 to 8.8 (88) which i s therefore the range of passivation of aluminum.  In this range MgZn  to aluminum and i s therefore corroded by the environment. pH i n the present work was  2  i s anodic Since the  8.6 i t i s reasonable to suggest that  corrosion i n i t i a t e d at the grain boundary precipitates i n a l l the aluminum a l l o y s . 2.  While dissolution takes place at the anode, atomic hydrogen  i s liberated at the aluminum s o l i d solution which acts as the cathode. The hydrogen atoms diffuse to the region of the notch.  Dissolution  assists i n maintaining the notch while deformation at the t i p tends to blunt i t .  Hydrogen probably diffuses to the blocked d i s l o c a t i o n arrays  at the region ahead of the notch (90).  These arrays can be produced by  the blocking of dislocations at grain boundaries and by formation of a shear stress f i e l d i n the t r i a x i a l region beneath a notch (9). A l t e r n a t i v e l y , hydrogen can nucleate at the incoherent interface of the grain boundary p r e c i p i t a t e .  It has been shown (91) that the absence  - 148 -  of large voids and mobile  dislocations i n iron whiskers i s related to  t h e i r resistance to hydrogen  embrittlement.  The hydrogen precipitates at the defects i n the molecular form and an expansion due to the gas pressure takes place as proposed by Zapffe ( 9 2 ) , Tetelman and Robertson  (8) and. others.  pressure P lowers the applied stress growth.  The internal  aF, necessary to cause crack  It i s proposed that observable crack propagation takes place  a f t e r the crack has grown to a c r i t i c a l length 2 C . q  During the  i n i t i a t i o n time, the crack grows to this c r i t i c a l length 2 C by a q  linking up of microvoids formed by recombination of atomic hydrogen at incoherent interfaces.  The hydrogen pressure i s addivtive to the  t r i a x i a l stress state and a s s i s t s i n the separation of the matrix from incoherent p a r t i c l e s . Once a crack of a c r i t i c a l length i s i n i t i a t e d slow crack growth occurs.  This process consists of void formation ahead of the  crack by the action of t r i a x i a l stress state and hydrogen pressure and d u c t i l e linking up to the voids i n the presence of applied s t r e s s . When the crack reaches a c r i t i c a l length 2 C „ i t s a t i s f i e s the c r i t e r i a f o r unstable fracture. From G r i f f i t h ' s c r i t e r i a , the fracture stress, f o r plane s t r a i n condition i n an e l a s t i c s o l i d i s given by  a  p r  = (^ ) (1-v )irC p  2C  C  = internal crack length  2  (7) .  - 149 -  From Orowan's correction (102)  a  2EP -  = (  ,1/2 )  (1-V )TTC  (8)  Z  F  where P = y + P ' w Y  =  surface energy  P = p l a s t i c work w r  Also G = 2P where 2P = t o t a l energy associated with generation of two surfaces of unit areas.  2 (1-v )  cA*C  )  -»- 1 f o r most metals  1  /  2  = (EG)  1 / 2  (9) / EG or  The term OpCp  =  OpF  =  / -nC„  i s known as the stress intensity parameter at  fracture stress  K  2 C  =• EG  (10)  It w i l l be noticed that processes occurring during i n i t i a t i o n are the same as those taking place during the propagation of the stress corrosion crack with one important difference. l i e s i n the value of the stress intensity factor K . T  The difference During  initiation  - 150 -  the stress i n t e n s i t y factor remains e s s e n t i a l l y constant at a value of K = K-j-g^j-..  Once crack propagation starts there i s a sharp increase  in the stress intensity factor even i f the actual crack v e l o c i t y remains rather constant (77).  When Kj increases to a value K^,  unstable  fracture takes place at an i n d e f i n i t e l y fast crack v e l o c i t y . Thus, the propagation stage i s marked by a rapidly increasing stress intensity factor and a r e l a t i v e l y constant crack v e l o c i t y u n t i l i n s t a b i l i t y i s reached.  The high stress concentration influences  the kinetics i n such a way that a d i f f e r e n t process becomes rate controlling.  Consequently,  the activation energies of the i n i t i a t i o n  and propagation stages are s i g n i f i c a n t l y d i f f e r e n t i n the same a l l o y environment system. Further support of the hydrogen hypothesis i s furnished by the work of E l l s and Evans (96) .  They observed hydrogen  agglomerates  at the grain boundaries i n samples of proton i r r a d i a t e d aluminum at temperatures  less than 100°C.  at higher temperatures.  Grain boundary cracking was  observed  According to them, the main driving force  for nucleation and i n i t i a l growth of hydrogen f i l l e d bubles i s due to a supersaturation of gas i n the metal and bubbles can form at 100°C with gas concentrations as low as 1 ppm.  Their work shows that micro-  scopic c a v i t i e s are not necessary f o r bubble-formation even though a cavity i s a highly preferred point of nucleation of hydrogen bubbles. 3  Bubble growth i s possible when the pressure i s  10  atmospheres i n  a bubble of O.l^uradius.  This i s equivalent to a hydrogen concentration  of  The concentration i s calculated from Sievert's  law  10  ppm  at 100°C.  - 151 -  S=  1/2  K •P  3  S = 2 x 10  -19400 exp (• RT  The d i f f u s i v i t y of hydrogen at 100°C i s 8 x 10 charging experiments  (11)  cm /sec.  Hydrogen  on Al-Mg-Zn and Al-Zn alloys i n the present work  showed that there i s a s i g n i f i c a n t reduction i n the f a i l u r e time i f cathodic charging i s continued up to the application of load and then stopped.  This enables the d i s s o l u t i o n to occur i n the absence of a  protective cathodic p o t e n t i a l . a notch.  Dissolution results i n the formation of  The hydrogen forms voids ahead of the crack which link up by  ductile tearing. In the Al-Zn a l l o y , the crack propagation started as soon as maximum load was applied and p r i o r hydrogen charging s i g n i f i c a n t l y shortened the propagation time.  In the Al-Mg-Zn a l l o y both crack  i n i t i a t i o n and propagation times were shortened s i g n i f i c a n t l y by hydrogen charging.  In both the alloys i f charging was continued after maximum  load, f a i l u r e times could be prolonged considerably.  Therefore i t may  be surmised that anodic dissolution accompanied by the maintenance of a sharp notch i s an essential step i n the stress corrosion process. This step determines the flux of hydrogen available for d i f f u s i o n as well as the stress concentration existing at the root of the notch. Hydrogen charging results i n an increase i n the flux of hydrogen and a decrease i n time required for d i f f u s i o n to form the microvoids.  Con-  sequently, both the i n i t i a t i o n and propagation times should be reduced.  - 152 This i s observed i n the Al-Mg-Zn and Al-Zn a l l o y s . Haynie and Boyd (95) used an autoradiographic  technique to  determine whether hydrogen concentrates at the grain boundaries, using 1 N Na^SO^ containing t r i t i a t e d water. present i n higher concentrations  at the grain boundaries of stressed  Al-Mg-Zn alloy than i n the unstressed stress apparently boundaries.  They found that hydrogen i s  material.  They concluded that  increases the s o l u b i l i t y of hydrogen at the grain  When a s t r i p specimen i s anodically polarised on one  side  and cathodically/ polarised on the other, immediate stress corrosion cracking takes place.  There i s no cracking when the specimen i s polarised  only anodically or cathodically. Haynie and Boyd conclude that hydrogen may stress corrosion process i n two ways.  It may  be involved i n the  increase  localised  attack at the grain boundaries by increasing the cation gradient through the strained oxide f i l m .  The r e l a t i v e d i s t r i b u t i o n of protons i n the  oxide at the grain boundary and grains may boundary attack.  cause p r e f e r e n t i a l grain  Both anodic potential and applied t e n s i l e stress  increase the d i f f u s i v i t y of the aluminum ions through the oxide f i l m . It has also been shown (96) that alloys susceptible to stress corrosion show higher rates of hydrogen evolution than the alloys.  non-susceptible  In the present work, when pure aluminum samples were amalgamated  to break up the oxide f i l m and exposed to ethyl alcohol, carbon t e t r a chloride and s a l t solutions, evolution of hydrogen was rate of evolution was  observed.  The  higher i n the s a l t solution than i n alcohol.  Normally, the d i s t i n c t i o n between stress corrosion  cracking  and hydrogen embrittlement i s made by applying anodic and cathodic  current.  - 153 -  If embrittlement i s enhanced by an anodic current the process i s stress corrosion cracking and i f i t i s enhanced by a cathodic current then the process i s considered hydrogen embrittlement.  But this kind of  d i s t i n c t i o n i s inadequate i f anodic d i s s o l u t i o n i s an essential step for hydrogen embrittlement as postulated i n the present model.  The  function of anodic dissolution i s not clear.  Drawbacks of the hydrogen model 1.  I f hydrogen d i f f u s i o n was the rate c o n t r o l l i n g process  cathodic charging during stress corrosion should lead to a reduction in stress corrosion l i f e t i m e .  However, cathodic protection i s possible  in aluminum a l l o y s . 2.  Only the crack i n i t i a t i o n i n alcohol and crack propagation  in s a l t solution have apparent activation energies equal to activation energy for d i f f u s i o n of hydrogen.  The i n i t i a t i o n rate i n s a l t solution  and the propagation rate i n alcohol have activation energies which are very different from that f o r hydrogen d i f f u s i o n . mechanism cannot convincingly explain (a)  A hydrogen d i f f u s i o n  why  The apparent activation energy of propagation i s less than that of i n i t i a t i o n .  (b)  The differences i n respective a c t i v a t i o n energies i n s a l t solution and alcohol.  3.  The application of anodic potential seems to be necessary  for hydrogen to have any e f f e c t . by the present hydrogen model.  The reason for this can not be explained  - 154 -  4.  The d i f f u s i v i t y of hydrogen at -70°C i n alcohol i s too  low to account for the observed propagation times. 5.  A hydrogen model cannot explain why oxygen i s necessary  for stress corrosion.  4.9.  The Requirements of a Satisfactory Model A s a t i s f a c t o r y stress corrosion model for aluminum alloys must  explain the following observations: 1.  The stress corrosion cracking i s intergranular i n both  s a l t solution and ethanol.  The stress corroded area appears to have a  rough surface. 2.  The presence of coherent precipitates enhances the  s u s c e p t i b i l i t y to cracking s i g n i f i c a n t l y .  The presence of grain boundary  precipitates and a narrow PFZ i s also associated with high stress corrosion s u s c e p t i b i l i t y .  On the other hand, solution treated and  p r e c i p i t a t e free Al-Mg alloy as well as pure aluminum are immune to stress corrosion. 3.  A change i n environment, at a given stress level changes  the crack i n i t i a t i o n and crack propagation rates. 4.  A change i n the i n i t i a l applied stress always changes the  i n i t i a t i o n and propagation rates for a given alloy-environment system. 5.  The temperature  s e n s i t i v i t i e s of crack i n i t i a t i o n and pro-  pagation rates are d i f f e r e n t for a given alloy-environment system. temperature  The  s e n s i t i v i t y also changes with the environment.  S p e c i f i c a l l y the model should explain the apparent a c t i v a t i o n energies l i s t e d i n the table (17).  - 155 -  Table (17) APPARENT ACTIVATION ENERGIES OF CRACK INITIATION AND PROPAGATION FOR ALUMINUM ALLOYS  Alloy  Environment  Al-2.6Mg-6.3Zn  Average ' Q in kcals/mole during Initiation Propagation 1  Salt solution  26  14  D.D. water  26  --  11  Ethanol  12  1-4  Al-8.6Mg  Ethanol  12  1-4  Al-22Zn  Ethanol  12  1-4  II  4.9.1.  Models Involving Either Dissolution or Deformation and Their Drawbacks (a)  According to H e l f r i c h (72) the rate c o n t r o l l i n g step i n  the stress corrosion of Al-Mg-Zn a l l o y i n s a l t solution i s the anodic dissolution of MgZn p a r t i c l e s at the grain boundary. 2  The model i s  based on the equation of H i t l i g and Charles (100), that activation free energy for a corrosion reaction under stress can be expressed as  Q*(s)  = Q*(o)  +  s(|21) + . . . . s = 0  Q*(o)  = a c t i v a t i o n energy at zero stress s = stress  aO* (•?—) ^9s }  = V* = a c t i v a t i o n volume  - 156 -  V* i s equated to the molar volume of the phase MgZr^. Two types of anodic sites have been observed i n Al-Mg-Zn alloys.  Hunter (12) found that an a c i d i f i e d (pH = 1) NaCl/AlCl^ solution  attacked the grain boundaries but not the MgZn^ p r e c i p i t a t e .  On the  other hand Jacobs (20) and H e l f r i c h (21) observed attack at the constituent p a r t i c l e interface.  Pugh (26) observed that stress corrosion surface  had a corroded appearance.  In the present work the t y p i c a l stress  corrosion surface of figures (37-40) was observed.  The surfaces exhibited  a rough appearance i n both s a l t solution and ethanol. H e l f r i c h (72) did not distinguish between i n i t i a t i o n and propagation rates.  His model postulates d i s s o l u t i o n of grain boundary  precipitates leading to b r i t t l e fracture between p r e c i p i t a t e s but the d e t a i l s of the mechanism are not c l e a r .  It does not explain why maximum  s u s c e p t i b i l i t y i s associated with coherent p r e c i p i t a t e s and why the temperature s e n s i t i v i t y i s d i f f e r e n t during i n i t i a t i o n and propagation stages.  It i s d i f f i c u l t to explain the present  results on the basis  of Helfrich's model. (b)  Sodriks et a l (26) believe that l o c a l i z e d p l a s t i c deforma-  tion i n the PFZ causes p r e f e r e n t i a l anodic d i s s o l u t i o n of the deforming metal within the PFZ.  The crack propagation rate of 0.35 mm/hr i s 2  equivalent to a current density of 0.6 amp/cm  which i s believed to be  reasonable for a corrosion process (26). This model again does not explain the r o l e of coherent p r e c i p i tates, the stress s e n s i t i v i t y and the temperature dependence of the stress corrosion process.  It does not account for the i n i t i a t i o n stage  nor does i t describe i n d e t a i l mechanics of crack propagation.  - 157 -  (c)  Jacobs  (20) has proposed that both the p r e c i p i t a t e and f i l m  free aluminum alternate as anodes with respect to the f i l m covered crack walls.  The i n t e r p a r t i c l e distance i s bridged by mechanical  fracture.  However no evidence of d u c t i l i t y has been found i n the stress corroded area i n the present work. (d)  Krafft and Mulherin (77) believe that microvoid coalescence  takes place adjacent to a p r e c i p i t a t e or an i n c l u s i o n . link up with the crack t i p and the crack advances.  The microvoids  The fractography  i n t h i s work does not rule out the p o s s i b i l i t y that microvoid formation takes place which i s subsequently obliterated by general corrosion. Once again this theory does not explain i n d e t a i l the role of heat treatment, stress and temperature during i n i t i a t i o n and propagation stages. Models based on d i s s o l u t i o n cannot explain some of the observations l i s t e d i n section (4.8). why pre-immersion  In addition such models do not  explain  i n stress corroding solutions did not reduce the  i n i t i a t i o n or propagation times. Similarly models based on deformation processes a f f e c t i n g the rate of s l i p step emergence, s l i p step height and spacing cannot account for the fact that a change i n environment  always changes the rate of  i n i t i a t i o n and propagation i n the same a l l o y at the same stress  level.  Therefore while apparent activation energies of 26 kcals/mole and 12 kcals/ mole are close to the apparent activation energies such deformation processes as c r o s s - s l i p of dislocations and intersection of forest dislocations respectively; s t i l l the deformation process alone cannot explain a l l the r e s u l t s .  - 158 Therefore one can only conclude that both d i s s o l u t i o n and deformation processes contribute to the rate of i n i t i a t i o n and propagation of stress corrosion crack.  A model based on t h i s conclusion i s now  proposed.  4.10  The Proposed Model The application of stress results i n generation of dislocations  at the grain boundary precipitates as proposed by Sedriks et a l (26). The dislocations move across the p r e c i p i t a t e free zone u n t i l t h e i r motion i s impeded by the precipitates adjacent to the PFZ.  The  precipitates are sheared by dislocations and attacked by the environment as shown i n f i g . (52a).  Corrosion ceases, when the p r e c i p i t a t e i s  completely dissolved and the aluminum s o l i d solution i s covered by a thin f i l m of i t ' s oxide.  This p o s i t i o n i s shown i n f i g . (52a).  Further  advance of the crack must await the a r r i v a l of the next sheared precipitate.  After a certain time the sheared p r e c i p i t a t e arrives at  the crack t i p and i s dissolved as before, f i g . (52b).  The process of  cooperative dissolution and deformation continues and the crack advances along a s l i p plane u n t i l i t reaches an intersecting s l i p plane. the crack changes d i r e c t i o n turning towards the PFZ.  Then  This change of  d i r e c t i o n i s favoured because the shearing i s most pronounced adjacent to the PFZ.  The position i s shown i n f i g . (52c).  The crack reaches  the PFZ and then turns away from i t and the sequence s t a r t i n g from fi.g  (52a) i s repeated.  The result i s a zig-zag crack path confined  close to the PFZ and produced by l i n k i n g of tunnels. During the i n i t i a t i o n stage the crack reaches a c r i t i c a l length u n t i l K  T  > K „ . T  r r  Then crack propagation starts with the same  Cc)  (d)  Roughened surface Figure (52). I l l u s t r a t i o n of the proposed model.  - 160 sequence o f events increasing rate.  except t h a t deformation  T h e r e f o r e , t h e r a t e o f crack p r o p a g a t i o n i s much  g r e a t e r than t h a t o f i n i t i a t i o n . no l o n g e r support  When the reduced  the a p p l i e d l o a d , the s t r e s s  and f a i l u r e o c c u r s .  c r o s s s e c t i o n can  intensity factor  -  The mode o f f a i l u r e i s g e n e r a l l y i n t e r c r y s t a l l i n e  and i s accompanied by v o i d The  takes p l a c e a t an ever-  temperature  formation. and s t r e s s dependence d u r i n g i n i t i a t i o n and  p r o p a g a t i o n stages can be expressed as  Rate = A  a " exp (-Q/RT)  q  (12)  or under c e r t a i n c o n d i t i o n s  Rate = B  where A  q  and B  q  q  e  p o  exp (-Q/RT)  (13)  a r e c o n s t a n t s depending on t h e a l l o y - e n v i r o n m e n t  thermo-mechanical treatment, It i s proposed  system,  orientation of grains etc.  t h a t the apparent  a c t i v a t i o n energy  observed  i n t h e p r e s e n t work i s composed o f two terms, one f o r a d e f o r m a t i o n process and the other f o r a d i s s o l u t i o n p r o c e s s . initiation  T r a n s i t i o n from  :  t o p r o p a g a t i o n i s accompanied by a change i n t h e r a t e c o n t r o l -  l i n g p r o c e s s o f d e f o r m a t i o n , but no change i n t h e p r o c e s s o f d i s s o l u t i o n .  rSo t h a t  n ^ A Rate = A  N  q  a  t ^Def. , ^ D i s s . e x p ( — — ) ( — ^ — )  .... (14)  - 161 -  The  equations  I f creep alone was  from  (13) are s i m i l a r t o creep equations  (89).  c o n t r o l l i n g the r a t e o f c r a c k i n g , then a change i n  environment from s a l t not a f f e c t  (12) and  s o l u t i o n t o e t h a n o l , u s i n g the same a l l o y ,  the apparent  a c t i v a t i o n energy.  26 k c a l s / m o l e t o 12 k c a l s / m o l e .  should  But the v a l u e o f Q changes  T h e r e f o r e d i s s o l u t i o n as w e l l  as  creep must be i n v o l v e d i n the i n i t i a t i o n o f s t r e s s c o r r o s i o n c r a c k s . S i m i l a r l y i f d i s s o l u t i o n alone was to e x p l a i n the s t r e s s s e n s i t i v i t y change i n apparent The  sum  r a t e c o n t r o l l i n g , then i t i s d i f f i c u l t o f the c r a c k i n g r a t e s -as w e l l as  a c t i v a t i o n energy between i n i t i a t i o n and o f apparent  the  propagation.  a c t i v a t i o n energies during crack  initiation  i s g i v e n by Q_ „ Def  +  Qr, * Def  +  x  T h e r e f o r e Q„ cannot Def  x  x  Q_. Diss  =  26 i n s a l t  Qr, • Diss  =  12 i n e t h a n o l  12  kcals/mole.  exceed  11  solution  (15) (16) • J  During c r a c k p r o p a g a t i o n the d i s s o l u t i o n p r o c e s s i s b e l i e v e d to remain the same, but s i n c e the observed decreases  the creep p r o c e s s must change. Q'^ + Def  x  Q'n x: Der  x  x  +  apparent  activation  energy  Therefore,  Q^Diss  =  14 i n s a l t  solution  QnDiss  =  2 i n ethanol  (17) J  (18)  I f the creep process d u r i n g i n i t i a t i o n i s c o n t r o l l e d by o f d i s l o c a t i o n s from segregated atmosphere as proposed Parkins  (108) then Q  D g  £  - H  escape  by Congleton  and  kcals/mole.  S u b s t i t u t i n g i n equations i s o l u t i o n = 14 k c a l s / m o l e , CL. Diss  (15) and  (16) we  o b t a i n CL. Diss x  in- s a l t  i n ethanol = 1 k c a l / m o l e .  x  S u b s t i t u t i n g these v a l u e s i n equations Q'pgf d u r i n g p r o p a g a t i o n - 0 ^ 1  kcal/mole.  (17) and  (18) we  get •  - 162 Summarising: Q„. i n s a l t. s o l u t i o n and D .D. water Diss  14 k c a l s / m o l e  x  Q,. i n ethanol diss  1 kcals/mole  Q  D e f  during i n i t i a t i o n  Q  D e f  during propagation  Then t h e apparent  11 k c a l s / m o l e 1 kcals/mole  a c t i v a t i o n e n e r g i e s can be added as f o l l o w s :  Table (18) PROPOSED ACTIVATION ENERGIES Alloy  Env.  Initiation Q  Diss  Q  Def  Propagation Q  Total  Q  Diss  ^Def  Q  Total  14  11  25  14  1  15  14  11  25  14  1  15  Ethanol  1  11  12  1  1  ' 2  Al-Mg  Ethanol  1  11  12  1  1  2  Al-Zn  Ethanol  1  11  12  1  1  2  Al-Mg-Zn  Salt D.D.  1!  water  The v a l u e s o f Cv. and add up t o t h e observed apparent Diss Def x  r  r  r  r  activation  energies l i s t e d i n Table (17).  The a c t i v a t i o n e n e r g i e s o f e l e c t r o c h e m i c a l d i s s o l u t i o n o f t h e coherent or p a r t i a l l y coherent p r e c i p i t a t e s available i n literature.  i n salt solution oralcohol  a r e not  I t i s not known whether d i s s o l u t i o n i s  c o n t r o l l e d by a n o d i c o r c a t h o d i c p r o c e s s e s .  A c c o r d i n g t o West (85)  metals can have apparent a c t i v a t i o n e n e r g i e s o f e l e c t r o c h e m i c a l . d i s s o l u t i o n r a n g i n g from 5-25 k c a l s / m o l e which a r e s t r o n g l y on the a n i o n p r e s e n t .  dependent  - 163 -  It  i s p o s s i b l e t h a t the r a t e c o n t r o l l i n g step  hydrogen a t the cathode by any o f the f o l l o w i n g  a)  M-H  +  M-H  -  H  i s desorption of  reactions:  2  (adsorbed hydrogen on t h e metal) b)  (H 0)  +  3  +  M-H  +  e  -  H  ( s o l v a t e d hydrogen) c)  D e p o l a r i s a t i o n i n the presence o f oxygen 4(M-H)  If the l a s t observation  +  0  2  -  2 H0 2  step was r a t e c o n t r o l l i n g , t h a t would e x p l a i n the  o f some workers  (12) t h a t oxygen i s e s s e n t i a l f o r s t r e s s  corrosion. The  d i s s o l u t i o n process i n a l c o h o l i s assigned  energy df 1 kcal/mole.  Such low v a l u e s  with t r a n s p o r t p r o c e s s e s i n the l i q u i d Boreskov  The  of Q are generally (101).  (105) have observed the f o r m a t i o n  by a l c o h o l and p o s t u l a t e  an a c t i v a t i o n  Greenler  of surface  associated  (104) and  e t h o x i d e on alumina  t h a t hydrogen ions a r e evolved  i n the r e a c t i o n .  r a t e c o n t r o l l i n g . step c o u l d be the d i f f u s i o n o f r e a c t i o n p r o d u c t s  i n the crack. The  d i s s o l u t i o n p r o c e s s e s a r e b e l i e v e d t o remain e s s e n t i a l l y  unchanged d u r i n g The  t h e t r a n s i t i o n from i n i t i a t i o n t o p r o p a g a t i o n .  deformation process during  by d i s l o c a t i o n - s o l u t e i n t e r a c t i o n . d u r i n g p r e - a g i n g can cause p i n n i n g  crack  i n i t i a t i o n can be c o n t r o l l e d  The s u p e r s a t u r a t i o n of d i s l o c a t i o n s .  of solute  Congleton and  - 164  Parkins  (108) have s t u d i e d the s u b s t i t u t i o n a l s t r a i n aging i n A l - Z n  and Al-Mg a l l o y s d u r i n g creep. a c t i v a t i o n energy i s between 10-11  They determined  t h a t the  apparent  o f escape o f d i s l o c a t i o n s from the segregated kcals/mole.  j-  where  -  =  40 exp  altmosphere  They a r r i v e d at an e x p r e s s i o n of the  (-U/RT)  form  (16)  e = strain rate e = plastic U = apparent  strain activation  energy  During crack p r o p a g a t i o n the i n c r e a s i n g s t r e s s at the c r a c k t i p causes  the r a t e c o n t r o l l i n g process of deformation  Deformation The  d u r i n g crack p r o p a g a t i o n appears  to change.  to be temperature  insensitive.  s h e a r i n g of p r e c i p i t a t e s takes p l a c e at an i n c r e a s e d r a t e l i m i t e d  not by the r a t e of escape o f d i s l o c a t i o n s from a segregated but r a t h e r by some o t h e r p r o c e s s .  The  atmosphere  l i m i t i n g process c o u l d be the long  range i n t e r a c t i o n o f d i s l o c a t i o n s which i s temperature  insensitive.  The s t r e s s s e n s i t i v i t y o f i n i t i a t i o n and p r o p a g a t i o n r a t e s c o u l d be e x p l a i n e d by steady s t a t e creep.  I f the time dependent s t r a i n d u r i n g  s t r a i n aging i s p l o t t e d a g a i n s t s t r e s s on a l o g - l o g p l o t , v a l u e o f n between 3 and 9 are o b t a i n e d f o r two Al-Zn a l l o y s used by Congleton Parkins  (108).  The  'n' v a l u e s o f Al-Zn i n a l c o h o l based  v a r i e s between 6 and 4.5 The  f o r i n i t i a t i o n and p r o p a g a t i o n  on equation  (89,109).  (12)  respectively.  r a t e equations f o r s t r e s s c o r r o s i o n i n i t i a t i o n and  t i o n are s i m i l a r i n form to creep equations  and  propaga-  But the parameters  - 165  -  are based on models i n v o l v i n g d i s l o c a t i o n climb  and  contain a c t i v a t i o n  energy f o r s e l f d i f f u s i o n , which i n aluminum i s 34 k c a l s / m o l e . value  i s too h i g h to f i t i n the proposed s t r e s s c o r r o s i o n model. Although the a s s i g n e d  are somewhat a r b i t r a r y they  values  o f apparent a c t i v a t i o n e n e r g i e s  are not unreasonable f o r the  b e l i e v e d t o be t a k i n g p l a c e d u r i n g s t r e s s c o r r o s i o n .  processes  Further  ments i n the model must await b e t t e r knowledge o f the r a t e processes and  This  deformation  Furthermore, more knowledge i s r e q u i r e d about process  the  taking place during stress corrosion.  The  Observations  The  model e x p l a i n s a l l the o b s e r v a t i o n o u t l i n e d e a r l i e r  the evidence generation  E x p l a i n e d by the Proposed Model  for i t i s largely circumstantial.  of a continuous  path  a c t i o n o f s y n e r g e t i c processes 1. l i n k up  controlling  i n v o l v e d i n the d i s s o l u t i o n of p r e c i p i t a t e s i n s a l t s o l u t i o n  ethanol.  4.10.1.  develop-  Cracking  o f sheared  o f chemical  It explains  heterogeneity,  o f d i s s o l u t i o n and  seems to be  but  the by  the  deformation.  i n t e r g r a n u l a r because the r a t e o f  p r e c i p i t a t e s w i l l be h i g h c l o s e to the PFZ.  The f i n  roughened surface seen i n f r a c t o g r a p h y i s c o n s i s t e n t with the model. • I t i s a l s o p o s s i b l e t h a t the c r a c k i n g i s i n i t i a t e d at rim l i n e s shown i n figs.  (37a,  38b,  41a)  which have a z i g - z a g appearance.  Although  the  'dry' t e n s i l e f r a c t u r e showed rim l i n e s such specimens a l s o s t r e s s corrosion  cracked. 2.  The  model e x p l a i n s the h i g h s u s c e p t i b i l i t y o f pre-aged  a l l o y s which have h i g h e r p r o p o r t i o n o f coherent  or p a r t i a l l y  coherent  - 166  -  p r e c i p i t a t e capable of being sheared.  A narrow PFZ would increase  the  rate of shearing of p r e c i p i t a t e s by moving them c l o s e r to sources of dislocations.  The grain boundary p r e c i p i t a t e s are e s s e n t i a l as a  source of d i s l o c a t i o n s , and to cross grain boundary by mechanism of K r a f f t et a l (77).  In  highly overaged a l l o y s with no p r e c i p i t a t e s that can  be  sheared, i t i s p o s s i b l e that the g r a i n boundary p r e c i p i t a t e s c o n s t i t u t e the path of stress corrosion by nucleating microvoids i n the manner suggested by K r a f f t and Mulherin.  The present model explains why  stress  corrosion does not occur i n s o l u t i o n treated Al-Mg and pure aluminum.. These two have no p r e c i p i t a t e s and therefore are immune to stress corrosion. 3.  I t explains why  a change i n environment or a change i n  stress l e v e l changes both the i n i t i a t i o n and propagation r a t e s .  This  i s so because d i s s o l u t i o n and deformation r e - i n f o r c e each other during stress corrosion. 4.  I t explains the temperature dependence of i n i t i a t i o n  and  propagation rates since temperature influences the d i s s o l u t i o n as w e l l as deformation r a t e s . 5.  I t may  account for the f a c t that oxygen i s e s s e n t i a l f o r  stress corrosion since oxygen may  a s s i s t i n cathodic d e p o l a r i s a t i o n of  hydrogen and increase the rate of d i s s o l u t i o n .  I t may  also explain  NaCl i s a more aggressive environment than NaBr and Nal. of Cl  at the anode may  be easier than that of Br~ or I ~ .  The  why  desorption  - 167  4.11.  -  Conclusions 1.  S t r e s s c o r r o s i o n c r a c k i n g i n aged Al-9Mg, Al-3Mg-6Zn and  Al-21.5Zn i s i n t e r g r a n u l a r i n both aqueous and and  ethanol environments  i s i n f l u e n c e d by the magnitude o f the a p p l i e d t e n s i l e  c o m p o s i t i o n , heat treatment  stress,  and m i c r o s t r u c t u r e o f the a l l o y ,  composition  o f the environment, geometry o f the specimen, o r i e n t a t i o n o f g r a i n boundaries  with r e s p e c t to the t e n s i l e a x i s and 2.  temperature.  S t r e s s c o r r o s i o n c r a c k i n g occurs by i n i t i a t i o n of c r a c k s  at s u i t a b l y o r i e n t e d g r a i n boundaries with an i n c r e a s i n g  and p r o p a g a t i o n o f the c r a c k s  s t r e s s i n t e n s i t y f a c t o r , l e a d i n g to mechanical  o f the remaining m a t e r i a l .  The  fracture  i n i t i a t i o n time c o n s t i t u t e s the major  p o r t i o n o f the t o t a l time t o f a i l u r e .  The  relative  l e n g t h s of  initiation  and p r o p a g a t i o n time as w e l l as the r e l a t i v e areas o f s t r e s s c o r r o s i o n and t e n s i l e f r a c t u r e can be v a r i e d by changing 3.  the s t r e s s  I n i t i a t i o n and p r o p a g a t i o n times  i n aqueous s o l u t i o n s  and e t h a n o l show an A r r h e n i u s type o f temperature The  apparent  a c t i v a t i o n energy  t i o n i s b e l i e v e d t o c o n s i s t of two sheared p r e c i p i t a t e s and  dependence.  d u r i n g i n t i a t i o n and  terms one  propaga-  f o r d i s s o l u t i o n o f the  the other f o r a d e f o r m a t i o n p r o c e s s r e s p o n s i b l e  f o r l i n k i n g up of the p r e c i p i t a t e s . step i n d e f o r m a t i o n  level.  Furthermore  the r a t e c o n t r o l l i n g  i s b e l i e v e d t o change d u r i n g the t r a n s i t i o n  i n i t i a t i o n to propagation.  The  from  synergetic processes of d i s s o l u t i o n  and  deformation a r e t h e r e f o r e r e s p o n s i b l e f o r s t r e s s c o r r o s i o n c r a c k i n g i n aluminum a l l o y s . 4.  The  \  f r a c t o g r a p h y and o t h e r evidence  model i n v o l v i n g s h e a r i n g and  i s consistent with a  l i n k up o f p r e c i p i t a t e s a d j a c e n t to the  boundary to form a path of chemical h e t e r o g e n e i t y and  subsequent  grain  dissolution  - 168  of t h i s path i n the 5. suffers  The  -  environment.  hydrogen embrittlement model i s p l a u s i b l e ,  from a number o f drawbacks.  p l a y s an  I t i s more l i k e l y t h a t  i n d i r e c t r o l e i n d i s s o l u t i o n by  c o n t r o l l i n g the  but hydrogen  cathodic  process.  4.12.  Suggestions f o r F u r t h e r Work 1.  I t would be u s e f u l  to extend the  study of i n i t i a t i o n  p r o p a g a t i o n times to Al-Mg-Zn-Cu and  Al-Cu a l l o y s  o r g a n i c environments and  stress  with the  present .2.  be  The  stress  o p e r a t i n g i n two  corrosion  i f the  types o f  sensitivity  The  k i n e t i c s o f notched specimens s h o u l d and  same r a t e c o n t r o l l i n g p r o c e s s i s  specimens.  k i n e t i c s o f p r o p a g a t i o n should be  same s t r e s s - i n t e n s i t y f a c t o r be  temperature  of smooth specimens, both d u r i n g i n i t i a t i o n  p r o p a g a t i o n stage to see  can  and  and  results.  compared t o that  3.  compare the  i n aqueous  and  so that  studied  at  a meaningful a c t i v a t i o n  the  energy  derived. 4.  More i n f o r m a t i o n i s r e q u i r e d  of p r e c i p i t a t e s then the  and  on the  aluminum s o l i d s o l u t i o n s  models p o s t u l a t i n g  dissolution  i n d i f f e r e n t media.  anodic or c a t h o d i c c o n t r o l l e d  However, the  kinetics  reactions  can  be  tested.  condition  may  be  s i g n i f i c a n t l y d i f f e r e n t from that under  Only  dissolution  k i n e t i c s under zero  stress  stress.  - 169 Appendix I  Tensile  Fractography o f Aluminum A l l o y s The  intergranular  mode o f f r a c t u r e  i n aluminum a l l o y s was predominantly  i n both smooth and notched t e n s i l e specimens.  In n o t c h e d '  specimens o f Al-Mg-Zn and Al-Zn t h e mode o f f r a c t u r e was t r a n s g r a n u l a r near the edge, r e s u l t i n g i n a c i r c u l a r band with s t r i a t e d referred  t o as " r i m l i n e s " i n t h i s work.  A l - Z n and ^ 150-300 y i n Al-Mg-Zn. typical tensile fracture  surface  fractures  The band width was ^40 y i n  The o t h e r f e a t u r e s  p r e s e n t on  were:  1.  Dimples o f cusps  2.  S l i p steps  3.  E l o n g a t e d dimples  The  t e n s i l e fracture surface  o f Al-22Zn f r a c t u r e d  d r i e d a i r a t 20 in/min. and 0.01 in/min. i s shown i n f i g u r e s The rim  edge i s shown i n f i g u r e s  slip  i s intergranular  l i n e s p r e s e n t on the g r a i n  e, g,h). that  (53a-f).  l i n e s i s p r e s e n t on the edge o f t h e n o t c h .  mode o f f r a c t u r e  i n dessicant (53a-h).  A c i r c u l a r band o f Near the c e n t r e , t h e  with dimples o f elongated dimples and  facets.  This  i s shown i n f i g u r e s  ( b-  The presence o f t h e r i n l i n e s can be e x p l a i n e d by t h e f a c t ,  t h e crack f r o n t  outwards.  a  i s n u c l e a t e d near t h e c e n t r e and then t r a v e l s  The c r a c k i s c o n f i n e d t o t h e notch and the f i n a l  i s t r a n s g r a n u l a r even though the i n i t i a l mode o f f r a c t u r e In notched and aged Al-Mg-Zn, the f r a c t u r e near t h e c e n t r e and t r a n s g r a n u l a r through the edge. f r a c t u r e i s p e r p e n d i c u l a r t o the t e n s i l e a x i s .  fracture  i s intergranular.  i s intergranular The t r a n s g r a n u l a r  The r i m l i n e s a r e shown  - 170 -  53(c)  5000 x  53(d)  2000 x  - 171 -  53(g)  53(h)  - 172 -  Figure (53).  Tensile fracture surface of aged, notched Al-22Zn a l l o y fractured i n dessicant dried a i r at crosshead speeds of 20 in/min (a-e) and 0.01 in/min (f-h).  (53a). Edge, e n c i r c l i n g a band of rim l i n e s . (53b). Near centre, dimples. (53c,d)Shear dimples with s l i p  lines.  (53e). Dimples of a smaller size than seen i n f i g . l a . (53f). Edge, intergranular facets and transgranular rim-  lines.  (53g,h)Near centre, dimples at two d i f f e r e n t magnifications.  - 173 -  54(c)  2000 x  54(d)  3000 x  - 174  54(e)  2000 x  gure (54).  -  54(f) 3000 x  Notched aged Al-2.6Mg-6.3Zn a l l o y pulled to fracture i n a i r at 0.01 in/min.  (54a).  Edge, rim lines on a transgranular fracture surface.  (54b).  Near centre, deformation on grain facets.  (54c,d).Magnified view of rim lines extending towards the centre (54e,f).Dimples  near the centre with s l i p lines on t h e i r surface  - 175 -  55 fa)  Figure  240 x  55(b)  4400 x  (55).  Aged Al-9Mg pulled to fracture i n a i r at 0.01 in/min  (55a).  Intergranular  (55b).  Dimples on the facets.  fracture with bright facets.  - 176 -  in figures (54a,c,d).  They seem to l i e p a r a l l e l to the d i r e c t i o n of  the crack propagation. Figure and transgranular  (54d)  fracture surfaces.  shows the junction of the i n t e r The intergranular fracture  surface with accompanying dimples and s l i p steps i s i l l u s t r a t e d by figures  (54b,e,f). The aged Al-9Mg a l l o y exhibits a bright faceted  without any evidence of rim l i n e s .  The absence of  surface  transgranular  fracture can probably be attributed to the wide PFZ and large intergranular |3 precipitates i n this a l l o y .  According to Unwin and Smith  intergranular fracture i s facored by wide PFZ and precipitates i n Al-Mg-Zn a l l o y s .  (106)  large grain boundary  Both features a s s i s t in the  nucleation  of voids ahead of an advancing grain boundary crack. The grain facets showing dimples can be seen i n figure  (55b).  No dimples can be seen under the microprobe as seen i n figure (55a). The fractographs  of smooth t e n s i l e specimens, showed the  same features as the notched specimens but without the presence of the transgranular  band of rim l i n e s .  In the absence of a notch the  crack presumably follows the usual intergranular mode of fracture.  - 177 Appendix II The straight l i n e f i t s were obtained by linear  regression  analysis using a suitable program i n the Hewlitt Packard 9100A calculator.  The program yielded the values of intercept b, slope m  and correlation c o e f f i c i e n t r . The significance of 'r' was determined by using the curves plotted by Thornhill  (110) .  In these curves the  value of r for d i f f e r e n t number of t r i a l s , i s plotted against the probability of exceeding r when true value i s zero.  I f the p r o b a b i l i t y  that the limited range of experiments might permit a value of r equal or greater than that actually obtained i s less than 0.05, i t i s generally  agreed that there i s genuine c o r r e l a t i o n . The c o r r e l a t i o n c o e f f i c i e n t s obtained during Arrhenius plots  of i n i t i a t i o n and propagation and t h e i r significance i s indicated i n , table  (19). The error i n apparent activation energy  AQ i s also  listed.  This error was calculated f o r a 95% confidence l i m i t as follows: The equation of the straight line f i t i s y = mx + b y = i n i t i a t i o n or propagation rate x = reciprocal absolute  temperature  Then error i n slope m i s given by m ± tR (111), where tR = confidence l i m i t  where x and y are the means and n i s the number of t r i a l s . The t values were taken from a table for appropriate number  - 178 Table (19) THE CORRELATION COEFFICIENT AND ERROR IN APPARENT ACTIVATION ENERGY FROM THE ARRHENIUS PLOTS Crack I n i t i a t i o n Alloy  Environment  Al-21.5Zn  Correlation  0 95  11  3.9  99% C H 0H  8  0 97  12  2.3  95% C H 0H  8  0 98  11  1.3  95% C H 0H  4  0 90  11  1.6  99% C H 0H  11  0 90  14.7  22  0 97  26  1.8  D.D. water  6  0 89  26  6.3  99% C H 0H  12  0 87  9.6  2.8  99% C H 0H  13  0 62  3.8  2.2  95% C H 0H  7  0 23  1.2  1.2 . Not s i g n i f i c a n t  99% C H OH  8  0 34  3.4  3.3  22  0 86  14.2  D.D. water  5  0 80  25  99% C H 0H  9  0 25  3.1  5  2  5  2  5  2  5  NaCl/K Cr0 2  Crack  Q ± AQ  ' 7  5  2  Al-8.6Mg  r  99% C H 0H 2  Al-2.6Mg-6.3Zn  Number of tests  2  4  5  Significant  4.6  Propagat Lon  Al-21.5Zn  2  5  2  Al-2.6Mg-6.3Zn  5  2  5  NaCl/K Cr0 2  Al-8.6Mg  2  4  5  2.9 15.3 3.2  Significant  ii  II  Significant Not s i g n i f i c a n t it  II  of experimental points and a p r o b a b i l i t y value. The log-log or semi-log plots were transposed to a linear scale before calculating the best f i t . In the.hydrogen  charging experiments  be shown that the differences  on Al-Zn alloys i t can  i n time to fracture with and without p r i o r  - 179 -  hydrogen charging are s i g n i f i c a n t . The appropriate t-function i s given by  Nl  N (N N -2) 2  1+  2  t = A + (N N ) 1+  2  In the present case N N A  = number of tests with p r i o r hydrogen charging 2  = number of tests without hydrogen charging = difference between the means of two sets of experiments,  x, and x„ are differences i n mean value and individual values 1 2 for the two sets of experiments. Substituting the appropriate values we obtain  t = 5.2 From Thornhill's curves (110) i t i s seen that a t value of 5.2 f o r 6 experiments can be regarded as highly s i g n i f i c a n t . Similarly the difference i n f a i l u r e times between p r i o r hydrogen charged specimens and those held f o r 8 days after hydrogen charging i s also highly s i g n i f i c a n t as i t gives a t value of 21.4 for 5 experiments.  - 180 Bibliography  1.  Hoar, T.P., Proc. Sec. Int. Conf. M e t a l l i c Corrosion, 1966, 14.  2.  Logan, H.L., The Stress Corrosion of Metals, John Wiley § Sons., 1966, 1.  3.  Staehle, R.W., Proc. Conf. Fund. Asp. St. Corr. Cracking. Ohio State Univ. 1967, 3.  4.  Logan, H.L.  5.  Logan, H.L. Ref. 2, 230.  6.  Staehle, R.W., Ref. 3, 6.  7.  Mears, R.B., Brown, R.H., Dix, J r . E.H., Symposium on Stress Corrosion Cracking of Metals, 1944, 329.  8.  Robertson, W.D., Tetelman, A.S., Strengthening Mechanisms i n Solids, A.S.M., 1962, 217.  9.  Stroh, A.N., Proc. Roy. S o c , 1954, 223, 404.  Ref. 2, 215.  10.  Hoar, T.P., Stress Corrosion Cracking and Embrittlement, Wiley and Sons, 1956, 107.  John  11.  Logan, H.L., J . Res. Nat. Bur. Std., 1952, 48, 99.  12.  Sprowls, D.O., Brown, R.H., Proc. Conf. Fund. Asp. St. Corr. Cracking, Ohio State University, 1967, 501.  13.  Pugh, E.N., Jones, W.R.D., Metallurgia, 1961, 63, 3.  14.  Farmery, H.K., Evans, U.R., J . Inst, of Metals, 1955-56, 8£, 413.  15.  Gruhl, W. M e t a l l , 1965, 1£, 206.  16.  McEvily, J r . , A.J., Clark, J.B., 60, 661.  17.  Thomas, G., J . Inst. Metals, 1960-61, 8_9, 287.  18.  H o l l , H.A., Corrosion, 1967, 25, 173.  19.  Gruhl, W., Co'rdier, H., Z. Metallkde., 1958, 49, 291.  20.  Jacobs, A.J., Proc. Conf. Fund. Asp. St. Corr. Cracking, Ohio State University, 1967, 530.  Bond, :A.P., Trans A.S.M., 1967,  - 181 21.  H e l f r i c h , W.J.,  Corrosion, 1968,  24,  423.  22.  Speidel, M.O., Proc. Conf. Fund. Asp. St. Corr. Cracking, Ohio State University, 1967, 561.  23.  Uhlig, H.H.,  24.  S t o l o f f , N.S.,  25.  Haynie, F.H., Boyd, W.K., Proc. Conf. Fund. Asp. St. Corr. Cracking, Ohio State University, 1967, 580.  26.  Sedricks, A.J., Slattery, P.W., 62, 238.  27.  Beck, A.F., Sperry, P.R., Proc. Conf. Fund. Asp. St. Corr. Cracking, Ohio State University, 1967, 529.  28.  G i l b e r t , P.T.,  29.  Perryman, E.C.W., Hadden, S.E., J . Inst. Metals, 1950,  30.  Paxton, H.W., Proctor, R.P.M., Proc. Conf. Fund. Asp. St. Corr. Cracking, Ohio State University, 1967, 509.  31.  Fink, W.L.,  32.  Thomas, G., Nutting, J . , Journal Inst. Metals, 1957-58, S6_, 10.  33.  Kelly A., Nicholson, R.B., 10, 151.  34.  Schmalzried H. , Gerold, V., Z. Metallkurde, 1958,  35.  Kelly, A., Nicholson, R.B.,  36.  Taylor, J.L. J . Inst. Metals, 1958-59, 87_, 24.  37.  Lorimer, G.W.,  38.  Embury, J.D.,  39.  Mondolfo, L.F., J . Inst. Metals, 1969,  40.  Dahl, 0., Detort, K.,Z.  41.  Garwood, R.D., 60, 88, 375.  42.  Fox, F.A.,  Treatise on Fracture, Academic Press, Johnston, T.L., Acta Met.  Hadden, S.T.,  Smith, D.W.,  1967.  1963,(n) 251.  Pugh, E.N.,  Trans. A.S.M., 1969,  J . Inst. Metals, 1950,  Trans. AIME 1937,  124,  77,  237. 77_, 207.  166.  Progress i n Materials S c i . , 1963, 49_, 64.  Progress Materials S c i . , 1963,  Nicholson, R.B. Nicholson, R.B.  Acta. Met., Acta. Met.,  Metallkunde,  1966, 1965,  14,  10_, 151.  1009.  13^, 403.  97, 95.  1955,  66_, 94.  Davies, A.L., Richards, G.L. , J . Inst. Metals  Lardner, E., J . Inst. Metals, 1943,  69,  373.  1959-  - 182 -  43.  Fisher, A., J . Inst. Metals, 1941, 67_, 289.  44.  Clark, J.B., Acta. Met.,  45.  Garwood, R.D.,  46.  Geisler, A.H., Barrett, C.S., Mehl, R.F., Trans. Amer. Inst. Min. Met. Eng., 1943, 152, 201.  47.  Kelly, A., Nicholson, R.B.,  48.  Dew-Hughes, D.  49.  Dew-Hughes, D., Robertson, W.D.,  50.  Ashby, M.F.,  51.  Thomas, G., Nutting, J . , J . Inst. Metals, 1959-60, 88_, 81.  52.  Mondolfo, L.F., Gjustein, N.A., Levinson, D.W. Inst. Min. Met. Eng., 1956, 206, 1378.  53.  Thackery, P.A.  54.  Embury, J.D., Nicholson, R.V.,  55.  Garwood, R.D.., Davies, A.L., J . Inst, of Metals, 1959-60, 88, 311.  56.  Carpenter, G.J.C., Garwood, R.D.,  57.  Niklewski, T., Spiegelberg, P., Sunbullik., Metal S c i . Jour., 1969, 3, 23.  58.  Talbot, A.M., 301.  59.  Fox, F.A.,  60.  Murakani, Y., Kawano, 0., Tanura, H., Mem. 1962, 24, Part 4, 411.  61.  Clark, J.B., J . Inst. Metals, 1968, 1_6, 141.  62.  Hansen, M. Constitution of Binary a l l o y s , McGraw-Hill, 1958,  63.  Kelly, A., Electron microscopy on strength of c r y s t a l s , Interscience.  64.  Beck, A.F., Sperry, P.R., Proc. Conf. Fund. Asp. St. Corr. Cracking, Ohio State Univ., 1967, 513.  1968, 1_6, 141.  Davies, A.L. , J . Inst. Metals, 1959-60, 88_, 311.  Ph.D.  Ph.D.  Progress Materials S c i . , 1963,  Thesis, Yale University,  1_0, 244.  1959.  Acta. Met. 1960, 8_, 147.  Thesis, Cambridge University,  J . Inst. Metals, 1968, 96, Acta. Met.,  1961.  Trans. Amer.  234. 1965, 13_, 409.  Metal S c i . Jour., 1967, 1_, 202.  Norton, J.T., Trans. Am.  Inst. Min. Engrs., 1936,  Lardner, E., J . Inst. Metals, 1943, 69,  122,  373.  Fac. Engng Kyoto Univ.  108.  - 183 -  65.  Greenwood, G.W.,  Interfaces, Int. Conf. Melbourne, 1969,  66.  Speidel, M.O., Proc. Conf. Fund. Asp. St. Corr. Cracking, Ohio State University, 1967, 567.  67;  Perryman, E.C.W., Brook, G.B.,  68.  Kocks, U.F.,  P h i l . Mag.  69.  Logan, H.L.,  Ref. 2,  70.  Edeleanu, C ,  71.  Sager, G.F., Brown, R.H., Mears, R.B., Cracking of Metals, 1944, 255.  72.  H e l f r i c h , W., 1967, 27.  73.  Gruhl, W.,  74.  P r i e s t , D.K., "Stress-Corrosion Cracking and Embrittlement, Ed. W.D. Robertson, Wiley, New York, 1956, 81.  75.  Logan, H.L.,  76.  Timanova, M.A., INtercrystalline Corrosion and Corrosion of Metals under Stress, 263.  77.  K r a f f t , J.M.,  78.  Hoar, T.P., West, J.M., 1962, 268, 304.  79.  Swann, P.R., Embury, J.D., Second Berkeley Int. Materials Conf. on High Strength Materials, S. Wiley, New York, 1964, 327.  80.  Holt, D.L.,  81.  Thomas, G., High Strength Materials, Ed. Zackey, V.F., John Wiley S Sons, 1965, 361.  82.  Wei, R.P., Proc. Conf. Fund. Asp. St. Corr. Cracking, Ohio State University, 1967, page 7.  83.  H e l f r i c h , W.J.,  84.  Nichols, H., Rostoker, W.,  85.  West, J.M., 1965, 15.  1966,  225.  J . Inst. Metals, 1951, 79_, 19. L3, 541.  206.  J . Inst. Metals, 1951, 80,  187.  Symp. Stress Corr.  STO 425, Symposium on Stress Corrosion Testing, ASTM  Z. Metallak, 1962, 53_, 671.  Ref. 2, 228.  Mulherin, J.H., Trans. A.S.M., 1969, 62_, 64. Roy. Soc. (London) Proc. (Ser. A),  Trans. A.S.M., 1967, 60,  152.  Corrosion, 1968, 24, no. 12,  ,  426.  Trans A.S.M., 1963, 56_, 498.  Electrodepostion and Corrosion Processes, Van Norstand,  - 184 -  86.  Gruhl, W., Cordier, H . , M e t a l l .  1963, 17, 197.  87.  Sedricks, A.J., Slattery, P.W., 1969, 62, 816.  Pugh, E.N., Trans. A.S.M.,  88.  Pourbaix, M., Deltombe, E., Corrosion, 1958, 16_, 496.  89.  Garofalo, F., Fundamentals of Creep and Creep Rupture i n Metals, MacMillan, 1965, 74-79.  90.  Troiano, A.R., Trans A.S.M., 1960, 5_2, 75.  91.  Smialowski, M., Proc. Conf. Fund. Asp. St. Corr. Cracking, Ohio State University, 1967, 462.  92.  Zapffe, C , Sims, C , Trans, AIME, 1941, 145, 221.  93.  Tetelman, A.S., Proc. Conf. Fund. Asp. St. Corr. Cracking, Ohio State University, 1967, 446.  94.  E l l s , C.E., Evans, W., A.E.C.L.-1286. Chalk River Report CRGM1008, 1961, 8.  95.  Haynie, F.H., Boyd, W.K., Proc. Conf. Fund. Asp. St. Corr. Cracking, Ohio State University, 1967, 580.  96.  Marshall, T., Schaffer, G.J., J . Appl. Chem., 1959, £, 38.  97.  Sherby, O.D., Lytton, J.L., Dorn, J.E., Acta Met.,  98.  Panseri, C , Fredrighi, T., P h i l . Mag., 1958, _3, 1223.  99.  Detert, K. , Stander, J . , Z. Metallkde. 1961, 52_, 677.  1957, 5_, 219.  100.  H i l l i g , W.B.:, Charles, R.J., High Strength Materials, Ed. Zackay V.F., Wiley, 1965, 685.  101.  Burkin, A.R. The Chemistry Spon, 1966, 41.  102.  Orowan, E., Symposium on Internal Stresses, Inst, of Metals, 1947, 451.  103.  Tiesenhausen, E. von., Lund, J.A., ILZRA Project No. LM-8, Report No. 4., f i g . 14.  104.  Greenler, R.G., J n l . Chem. Physics, 1962, 37_, 2094.  105.  Boreskov, G.K., Shchekochikhin, Y.M., Markarov, A.D., Filiminov, V.N., Doklady Akademii Nauk SSSR, 1964, 156, 901.  of hydrometallurgical processes.  London  - 185 -  106.  Unwin, P.N.T., S m i t h , G.C.,  107.  Kubaschewski, 0., H o p k i n s , B u t t e r w o r t h s , 1962, 33.  J n l . I n s t . M e t a l s , 1969, 97_, 299. B.E., O x i d a t i o n o f M e t a l s and A l l o y s ,  108.. C o n g l e t o n , J . , P a r k i n s , R.N., J n l . I n s t . M e t a l s , 1969, 97^, 134. 109.  A l d e n , T.H., A c t a Met., 1969, 17_, 1435.  110.  Evans, U.R., 943-947.  111.  Duncan, A . J . , Q u a l i t y C o n t r o l and I n d u s t r i a l S t a t i s t i c s , 1965, 694,  112.  Hansen, M., C o n s t i t u t i o n o f B i n a r y A l l o y s , M c G r a w - H i l l , 106, 149.  C o r r o s i o n and o x i d a t i o n o f m e t a l s , A r n o l d , 1960,  1958,  

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