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Stress corrosion cracking of rotor steels in carbonate/bicarbonate and sodium hydroxide solutions Rechberger, Johann 1986

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^ 7 STRESS CORROSION CRACKING OF ROTOR STEELS IN CARBONATE / BICARBONATE AND SODIUM HYDROXIDE SOLUTIONS By JOHANN RECHBERGER D i p l . Ing. ETH, Swiss Federal I n s t i t u t e of Technology, 1982 A THESIS SUBMITTED IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF MASTER OF APPLIED SCIENCE in THE FACULTY OF GRADUATE STUDIES Department of Me t a l l u r g i c a l Engineering We accept t h i s thesis as conforming to the required standard THE UNIVERSITY OF BRITISH COLUMBIA February 1986 © Johann Rechberger, 1986 In presenting this thesis in partial fulfilment of the requirements for an advanced degree at the University of British Columbia, I agree that the Library shall make it freely available for reference and study. I further agree that permission for extensive copying of this thesis for scholarly purposes may be granted by the head of my department or by his or her representatives. It is understood that copying or publication of this thesis for financial gain shall not be allowed without my written permission. Department of /r/?&. The University of British Columbia 1956 Main Mall Vancouver, Canada V6T 1Y3 7 DE-6(3/81) ABSTRACT In t h i s study, the stress corrosion behavior of 3.5%NiCrMoV and 1%CrMoV steels was investigated. Tests were conducted at 95°C in carbonate/bicarbonate and sodium hydroxide solutions. Results from slow s t r a i n rate t e n s i l e tests and fracture mechanics experiments were compared. A new, e a s i l y machinable, specimen geometry for fracture mechanics experiments was tested. The influence of pH, po t e n t i a l , microstructure and inclusions were studied. Stress corrosion cracking (SCC) was found with a l l three s t e e l s in the 3.5M NaOH solution at active, passive and transpassive p o t e n t i a l s . In the 1M carbonate/ bicarbonate solutions, the 1%CrMoV steels showed very severe cracking only during slow s t r a i n rate experiments. No major difference in cracking behavior was found between a regular 3.5%NiCrMoV steel and a super clean 3.5%NiCrMoV s t e e l with low Mn and Si content. In a c i d i c CO2/H2O environments, crack t i p blunting was observed due to strong d i s s o l u t i o n processes. TABLE OF CONTENTS page Abstract i i Table of Contents i i i L i s t of Tables v i L i s t of Figures v i i L i s t of Symbols and Abbreviations x i i i Acknowledgements xv I. INTRODUCTION 1 1.1 Rotor-Steel 2 1.2 Microstructure / Heat treatment 4 1.3 Stress corrosion cracking of rotor a l l o y s 5 1.3.1 General mechanisms of SCC 5 1.3.2 C o n t r o l l i n g factors for SCC 7 1.3.3 Environment 7 1.3.4 Electrochemical potential 9 1.3.5 Material composition 10 1.4 Present objectives 12 I I . EXPERIMENTAL 13 2.1 Material 13 2.1.1 Chemical composition of the materials 13 2.1.2 Heat treatment 15 2.1.3 Sampling of rotor material 16 2.1.4 Mechanical properties 18 i i i page 2.1.5 Production of CaS and MnS r i c h steels 20 2.2 Microstructure / carbides / inclusions 21 2.2.1 Optical microscope investigations .. 21 2.2.2 Steel- nA" (3.5%NiCrMoV super clean) 22 2.2.3 Steel-"R" (3.5%NiCrMoV) 23 2.2.4 Steel-"0" (1%CrMoV) 25 2.2.5 TEM investigations 26 2.2.6 Steels "A" and "R" 27 2.2.7 Steel "O" 28 2.2.8 Q u a l i t a t i v e carbide analysis by EDX 33 2.2.9 Electron d i f f r a c t i o n of carbides ... 35 2.2.10 Inclusions 37 2.2.11 Summary of microstructures 44 2.3 Electrochemical p o l a r i z a t i o n studies 45 2.3.1 Sample preparation, test set up .... 47 2.4 Slow s t r a i n rate t e n s i l e test 50 2.4.1 Sample preparation, test set up .... 50 2.5 Fracture mechanics test 53 2.5.1 Specimen design 53 2.5.2 Test set up 59 I I I . RESULTS 63 3.1 Anodic P o l a r i z a t i o n curves 63 3.1.1 Carbonate solutions 63 3.1.2 Sodium hydroxide solutions 67 3.1.3 Steel "R" doped with sulphides 69 i v page 3.2 Slow s t r a i n rate t e n s i l e tests 71 3.3 Fracture mechanics tests 80 3.3.1 Carbonate solutions 81 3.3.2 Sodium hydroxide solutions 82 3.3.3 Fractography 88 3.3.4 Crack path 96 3.3.5 Dynamic SCC-test 97 IV. DISCUSSION 99 4.1 Interpretation of p o l a r i z a t i o n curves 99 4.1.1 Carbonate solution 99 4.1.2 Sodium hydroxide solution 104 4.2 Slow s t r a i n rate vs. fracture mechanics test 108 4.3 Comparison with l i t e r a t u r e data 111 V. SUMMARY 114 BIBLIOGRAPHY 117 v LIST OF TABLES page Table I Mechanical properties 18 Table II Results of slow s t r a i n rate tests 72 Table III Results of fracture mechanics (SCC) tests ... 84 vi LIST OF FIGURES page F i g . 1 Chemical composition (wt.%) of the rotor steels "O", "A" and "R" 14 Fi g . 2 Heat treatment of steels "A" and "R" 15 F i g . 3 Sampling map for steels "0","A" and "R" 16 F i g . 4 Microstructure of st e e l "A", centre region ... 22 F i g . 5 Microstructure of st e e l "A", surface region .. 22 Fi g . 6 Microstructure of st e e l "R" , centre region ... 24 F i g . 7 Microstructure of steel "R" , surface region .. 24 F i g . 8 Microstructure of st e e l "0", r a d i a l d i r e c t i o n with r o l l i n g texture 25 Fi g . 9 Microstructure of st e e l "0" 26 F i g . 10 TEM-picture, microstructure of st e e l "A" with carbides and subgrains 29 F i g . 11 TEM-picture, microstructure of s t e e l "A" with carbides and subgrains 29 F i g . 12 TEM-picture, microstructure of steel "R" with elongated subgrains 29 F i g . 13 TEM-picture, microstructure of s t e e l "R" with carbides and subgrains 29 F i g . 14 TEM-picture, s t e e l "R", elongated carbides inside subgrains 30 F i g . 15 TEM-picture, s t e e l "R", carbides inside subgrains and at grain boundaries 30 Fi g . 16 TEM-picture, s t e e l "R", carbides at sub-grain boundaries 31 Fi g . 17 TEM-picture, s t e e l "A", elongated carbides at subgrain boundaries 31 v i i page F i g . 18 TEM-picture, s t e e l "0", microstructure with f e r r i t e and bainite 32 F i g . 19 TEM-picture, s t e e l "0", big carbides at subgrain boundaries 32 F i g . 20 TEM-picture, s t e e l "0", carbides inside grains and at grain boundaries 32 F i g . 21 TEM-picture, steel "R", locations for EDX - analysis 33 F i g . 22 EDX - analysis, steel "R", matrix and carbides 34 F i g . 23 Electron d i f f r a c t i o n pattern of a carbide, steel "R" 35 F i g . 24 Electron d i f f r a c t i o n pattern of several carbides, steel "R" 36 F i g . 25 D i s t r i b u t i o n of inclusions in s t e e l "A" 38 F i g . 26 Oxide inclusions in steel "A" 39 F i g . 27 Oxide and sulphide inclusions in s t e e l "A" ... 39 F i g . 28 CaS inclusions in steel "R" 40 F i g . 29 Oxide inclusions in steel "R" 40 F i g . 30 CaS inclusions with oxide core, s t e e l "R" .... 41 F i g . 31 Elongated MnS inclusions in steel "O" 42 F i g . 32 Small round inclusions in steel "O" 42 F i g . 33 Microstructure and inclusions in CaS doped steel "R" 43 F i g . 34 Microstructure and inclusions in MnS doped steel "R" 43 F i g . 35 I l l u s t r a t i o n of a t h e o r e t i c a l and experimental p o l a r i z a t i o n curve 46 F i g . 36 Test set up for p o l a r i z a t i o n studies 49 F i g . 37 Set up for slow s t r a i n rate t e n s i l e tests .... 52 v i i i page F i g . 38 Specimen geometry for fracture mechanics tests 54 Fi g . 39 Different specimen geometries for fracture mechanics tests 54 F i g . 40 Stress in t e n s i t y vs. crack length for di f f e r e n t specimen geometries 58 F i g . 41 Test set up for stress corrosion experiments with fracture mechanics specimen 60 F i g . 42 I l l u s t r a t i o n of crack propagation measurements 62 F i g . 43 Pol a r i z a t i o n curves for steels "A", "R" and "0" in C0 2/H 20 64 F i g . 44 Pol a r i z a t i o n curve for steel "0" in NaHC03 (3.5*10"3M) 64 F i g . 45 Pol a r i z a t i o n curves for steel "0" in NaHC03 (3.5 ,10~ 2M), fast and slow scan 66 F i g . 46 Pol a r i z a t i o n curve for steel "0" in NaHC03 (3.5'10"1M) 66 F i g . 47 Pol a r i z a t i o n curves for steels "A", "R" and "0" in IM NaHC03 + 1M Na 2C0 3 66 F i g . 48 Pol a r i z a t i o n curves for steels "A", "R" and "0" in 3.5M NaOH 68 F i g . 49 Pol a r i z a t i o n curves for steel "R" in 3.5M NaOH, f i r s t and second scan 68 F i g . 50 Po l a r i z a t i o n curves for steels "R" and st e e l "R" doped with CaS in 3.5M NaOH 69 F i g . 51 Pol a r i z a t i o n curves for steels "R" and st e e l "R" doped with MnS in 3.5M NaOH 70 F i g . 52 Influence of po t e n t i a l on the reduction in area during SSRT-tests, stee l s "O" and "R" 73 ix page F i g . 53 Reduction in area and surface cracks of ste e l "0" at di f f e r e n t potentials in the C0 2/H 20 system, SSRT-test 74 Fi g . 54 Reduction in area and surface cracks of ste e l "0" at di f f e r e n t potentials in the 1M NaHC03 + 1M Na 2C0 3 system, SSRT-test 75 F i g . 55 Surface cracks in SSRT samples afte r tests in 3.5M NaOH (steel "R"), C0 2/H 20 (ste e l "0") and 1M NaHC03 + 1M Na 2C0 3 (steel "A") 77 Fi g . 56 SEM-picture, surface cracks of steel "A" afte r testing in 3.5M NaOH at -1000 mV S C E, SSRT-test 78 Fi g . 57 SEM-picture, surface cracks of ste e l "R" aft e r testing in 3.5M NaOH at -950 mV S C E, SSRT-test 78 Fi g . 58 SEM-picture, surface cracks of ste e l "O" afte r testing in 3.5M NaOH at -1000 mV S C E, SSRT-test 78 Fi g . 59 SEM-picture, fracture surface inside a crack, steel "A" tested in 1M NaHC03 + IM Na 2C0 3 at -750 mV SCE, SSRT-test 79 Fi g . 60 SEM-picture, fracture surface inside a crack, steel "0" tested in 1M NaHC03 + 1M Na 2C0 3 at -750 mVSCE, SSRT-test 79 Fi g . 61 SEM-picture, crack t i p of st e e l "A" aft e r t e s t i n g in C0 2/H 20 at the free corr. pot., SCC-test 82 Fi g . 62 Sress in t e n s i t y vs. crack v e l o c i t y for steels "A", "R" and "0" in 3.5M NaOH, SCC-test 85 Fi g . 63 Potential vs. crack ve l o c i t y for steels "A" and "0" in 3.5M NaOH, SCC-test 87 Fi g . 64 SEM-picture, crack t i p with intergranular fracture surface, steel "A" tested in 3.5M NaOH at -1000 mVSCE, SCC-test 91 x p a g e F i g . 6 5 S E M - p i c t u r e , s t r o n g l y e t c h e d 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 , s t e e l " A " t e s t e d i n 3 . 5 M N a O H a t - 1 5 0 m V S C E , S C C - t e s t 91 F i g . 66 S E M - p i c t u r e , p i t t e d 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 a t t h e c r a c k m o u t h , s t e e l " A " t e s t e d i n 1M N a H C 0 3 + 1M N a 2 C 0 3 a t - 7 5 0 m V S C E , S C C - t e s t 92 F i g . 67 S E M - p i c t u r e , 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 a t t h e c r a c k t i p , s t e e l " A " t e s t e d i n 1M N a H C 0 3 + 1M N a 2 C 0 3 a t - 7 5 0 m V S C E , S C C - t e s t 92 F i g . 68 S E M - p i c t u r e , s t r o n g c o r r o s i v e a t t a c k o f t h e 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 , s t e e l " 0 " t e s t e d i n 3 . 5 M N a O H a t - 1 0 0 0 m V S C E , S C C - t e s t 93 F i g . 6 9 S E M - p i c t u r e , t r a n s g r a n u l a r f r a c t u r e a p p e a r a n c e , s t e e l " O " t e s t e d i n 3 . 5 M N a O H a t - 1 0 0 0 m V S C E , S C C - t e s t 93 F i g . 70 S e c o n d a r y c r a c k s , s t e e l " O " t e s t e d i n 3 . 5 M N a O H a t - 1 0 0 0 m V S C E , S C C - t e s t 9 3 F i g . 71 S E M - p i c t u r e , s t r o n g l y a t t a c k e d f r a c t u r e s u r f a c e , s t e e l " O " t e s t e d i n 3 . 5 M N a O H a t - 1 5 0 m V S C E , S C C - t e s t 94 F i g . 72 S E M - p i c t u r e , s t r o n g l y a t t a c k e d f r a c t u r e s u r f a c e , s t e e l " O " t e s t e d i n 3 . 5 M N a O H a t - 1 5 0 m V S C E , S C C - t e s t 94 F i g . 73 S E M - p i c t u r e , f r a c t u r e s u r f a c e , s t e e l " O " t e s t e d i n 3 . 5 M N a O H a t - 1 0 0 0 m V S C E , S C C - t e s t 94 F i g . 74 S E M - p i c t u r e , r o u n d e d c r a c k t i p w i t h s t r o n g l y a t t a c k e d f r a c t u r e s u r f a c e , s t e e l " O " t e s t e d i n 1M N a H C 0 3 + 1M N a 2 C 0 3 a t - 7 5 0 m V S C E , S C C - t e s t 95 F i g . 7 5 S E M - p i c t u r e , c r a c k t i p w i t h MnS i n c l u s i o n , s t e e l "O" t e s t e d i n 3 . 5 M N a O H a t - 1 0 0 0 m V S C E , S C C - t e s t 95 F i g . 7 6 S E M - p i c t u r e , c r a c k t i p w i t h MnS i n c l u s i o n , s t e e l " O " t e s t e d i n 1M N a H C 0 3 + 1M N a 2 C 0 3 a t - 7 5 0 m V S C E , S C C t e s t 95 x i page F i g . 77 Crack branching, s t e e l "A" tested in 3.5M NaOH at -1000 mV S C E, SCOtest 96 Fi g . 78 SEM-picture, crack t i p aft e r dynamic SCC-test , steel "0" in 1M NaHC03 + IM Na 2C0 3 at -750 mV S C E 98 Fi g . 79 SEM-picture, crack t i p aft e r dynamic SCC-test, steel "0" in 1M NaHC03 + 1M Na 2C0 3 at - 750 mV S C E 98 Fi g . 80 E-pH diagram for the iron / water system 99 Fi g . 81 Discussion of a p o l a r i z a t i o n curve in terms of chemical reactions in the Bicarbonate/Carbonate solution 102 Fi g . 82 Discussion of a p o l a r i z a t i o n curve in terms of chemical reactions in the sodium hydroxide solution 105 x i i LIST OF SYMBOLS AND ABBREVIATIONS Symbol a crack length Aa crack propagation B specimen width CC, CL, LC description of fracture plane F Faraday constant I moment of i n e r t i a <- i anodic current density Kj stress i n t e n s i t y for mode I opening Kj(a) stress i n t e n s i t y with crack length a K I C fracture toughness K I S C C threshold stress i n t e n s i t y L lever arm M bending moment P load p density Cr bending stress v crack v e l o c i t y Vg£E vo l t s with resp. to standard calomel electrode V g H E v o l t s with resp. to stand, hydrogen electrode W specimen thickness W equivalent weight x distance from neutral f i b e r Y specimen geometry correction factor xi i i Abbreviation DCB double cantilever beam EDX energy dispersive X-ray analysis HSLA high strength low a l l o y steel PTFE polytetrafluorethylene ("Teflon") R c Rockwell-C hardness SCC stress corrosion cracking SEM scanning electron microscope Steel-"A" 3.5%NiCrMoV steel with low Mn and Si content Steel - " 0 " UCrMoV st e e l Steel-"R n 3.5%NiCrMoV steel SSRT slow s t r a i n rate t e n s i l e test TN-DCB T-notch double cantilever beam TEM transmission electron microscope v-K plot crack v e l o c i t y vs. stress intensity xiv ACKNOWLEDGEMENTS I would l i k e to express my sincere gratitude to my supervisors Dr. Desmond Tromans and Dr. Alec M i t c h e l l for the i r f r i e n d l y guidance and encouragement during t h i s work. I also want to thank Dr. Bruce Hawbolt and Dr. Robert Jaffee for t h e i r h e l p f u l discussions. Special thanks go to Clinton Fong and Robert Kelly for th e i r help during the experimental work. The always generous help of the staff and graduate students i s appreciated. My wife V i r g i n i a has given me a great amount of support and exhibited great patience throughout t h i s project. F i n a n c i a l support has been provided by an U.B.C. scholarship which I received for two years. Additional f i n a n c i a l assistance was given by the E l e c t r i c Power Research I n s t i t u t e (EPRI) and the Department of M e t a l l u r g i c a l Engineering of the University of B r i t i s h Columbia. The author has been very g r a t e f u l for these contributions. xv 1 I. INTRODUCTION Most modern low pressure (LP) steam turbines (shafts and discs) are made out of steels that have Ni, Cr, Mo and V as major a l l o y i n g elements. Stress corrosion cracking (SCC) of these steels can cause catastrophic f a i l u r e of the turbine with obvious safety hazard to personnel and enormous economic impact to the power generating industry. A survey in the U.S. in 1980 showed that more than one t h i r d of the inspected power plants had LP disc cracking p r o b l e m s 1 T h e cracks were often found in keyways and the blade attachment regions, both areas that are highly stressed and conducive to the formation of a crevice. From the crack length and the operating time, t y p i c a l crack growth rates of 10~ 1^ m/s to 10~ 1^ m/s were calculated-*. Even i f the service environment i s , in most cases, highly p u r i f i e d steam^, various corrosion products are found in the stress corrosion cracks, for example, iron oxides, chlorides, sulphates and carbonates. In some cases, strongly a l k a l i n e hydroxide deposits are found in the cracked regions-*. It can be assumed that various steam impurities get concentrated at locations of steam condensation and therefore produce l o c a l l y very aggressive envi ronments. To determine the s u s c e p t i b i l i t y of rotor s t e e l to SCC, various test methods are being applied. Examples are: 2 constant load, slow s t r a i n rate t e n s i l e tests or fracture mechanics tests. The test environments chosen are often quite d i f f e r e n t from the o r i g i n a l steam composition because of the need to simulate the worst conditions which could e x i s t , for example, in l o c a l condensates at a keyway. This also includes the corrosion p o t e n t i a l which can be contr o l l e d by imposing a s p e c i f i c potential onto the test piece by means of an external DC e l e c t r i c a l source. With these stress corrosion experiments, data can be obtained in a r e l a t i v e l y short time. But even so, test periods of several hundred hours are quite common. One has to bear in mind that turbines are in operation for many years or decades and, therefore, even very small cracking rates cannot be tolerated. 1 . 1 . Rotor - Steel The size of rotor forgings has steadily increased over the past years accompanied by a corresponding increase in i n l e t steam pressures and temperatures. Requirements arose for improved material properties for the high pressure (HP) and the intermediate pressure (IP) rotors at elevated temperatures and for the low pressure (LP) rotors at low temperatures. These requirements have been- met by improvements in s t e e l making processes, refinements of a l l o y compositions and heat treatment, as well as forging p r a c t i c e ^ . 3 Steel making process: The most commonly adopted method i s basic e l e c t r i c arc melting followed by a vacuum treatment of the molten s t e e l , either by ladle degassing or by stream degassing of the ingot. Careful ladle metallurgy methods allows control of very low levels of P and S, whereas the contents of As, Sb, Sn and Cu can only be kept low by proper selection of raw materials^. F i n a l l y , u p h i l l pouring under protective atmosphere prevents hydrogen and nitrogen pickup during the casting process 7. A l l o y composition^: The high temperature HP and IP rotors are mainly made out of st e e l with approximately 0.25% C, 1% Cr, 1% Mo and 0.3% V. The high steam admission temperatures of up to 540°C require a material with a good balance of creep strength, d u c t i l i t y and toughness. For the low temperature LP rotors, a 3.5%NiCrMoV - s t e e l i s normally used with a balance maintained between strength l e v e l and toughness. Creep properties are not important at the gas i n l e t temperatures below 300°C. The ni c k e l content may vary between 2 and 3.75%, depending upon the diameter, strength, and toughness requirements for the rotor. 4 1.2. Microstructure / Heat treatment LP as well as HP rotor steels are heat treated to produce an upper b a i n i t i c structure. This microstructure proved to give the best combination of strength, toughness and stress corrosion cracking resistance in the LP rotor s t e e l s ^ ' 1 0 / 1 1 and gave e s p e c i a l l y good creep rupture properties in the HP rotor s t e e l s 1 2 . The b a i n i t i c microstructure i s normally obtained by an austenizing treatment at 960-975°C for 1%CrMoV - steels and 830-850°C for 3.5%NiCrMoV - s t e e l s , followed by o i l or water quenching respectively. Tempering treatments are in the temperature range of 690-710°C for 1%CrMoV-steels and 600-620°C for 3.5%NiCrMoV-st e e l s . To minimize residual stresses, slow furnace cooling from the tempering temperature i s necessary, e s p e c i a l l y for large forgings. A major disadvantage of the heat treated NiCrMoV-ste e l s i s th e i r s e n s i t i v i t y to temper embrittlement in the temperature range of 350-450°C. This detrimental e f f e c t i s caused by the segregation of impurity elements at the p r i o r austenite grain boundaries (P, Sn, As, Sb, together with Ni, Cr and Mo) 1^' 1^. It has been found that a lowering of the Si and Mn content together with additions of Mo can reduce the severity of embrittlement^' 1^. 5 1.3. S t r e s s c o r r o s i o n c r a c k i n g o f r o t o r a l l o y s 1 . 3 . 1 . G e n e r a l m e c h a n i s m s o f S C C T h e v a r i o u s m o d e l s o f S C C c a n b e l a r g e l y d i v i d e d i n t o t w o b a s i c c l a s s e s : t h o s e w h i c h c o n s i d e r t h a t c r a c k p r o p a g a t i o n p r o c e e d s b y a n o d i c d i s s o l u t i o n a t t h e c r a c k t i p a n d t h o s e w h i c h c o n s i d e r t h a t c r a c k p r o p a g a t i o n i s e s s e n t i a l l y m e c h a n i c a l . A t p r e s e n t , t h e f i l m r u p t u r e o r s l i p s t e p d i s s o l u t i o n m o d e l s a r e w i d e l y q u o t e d . V e r m i l y e a 1 ^ p r o p o s e s a f i l m r u p t u r e m o d e l w h e r e b y a p r o t e c t i v e f i l m i s r u p t u r e d b y p l a s t i c s t r a i n a t t h e c r a c k t i p a n d a l l o w s t h e c r a c k t o a d v a n c e b y d i s s o l u t i o n u n t i l r e p a s s i v a t i o n o c c u r s ( r e p a s s i v a t i o n t i m e may r a n g e f r o m 1 0 " ^ t o 1 s e c o n d ) . A f t e r r e p a s s i v a t i o n , p l a s t i c f l o w n e a r t h e c r a c k t i p a c c u m u l a t e s s t r a i n u n t i l a t a c r i t i c a l v a l u e t h e f i l m r u p t u r e s a n d d i s s o l u t i o n c a n t a k e p l a c e a g a i n . S t a e h l e 1 7 f o r m u l a t e s a m e c h a n i s m b y w h i c h a p r o t e c t i v e s u r f a c e f i l m i s r u p t u r e d b y s l i p p r o c e s s e s . T h e n e w l y e x p o s e d b a r e m e t a l a t s l i p s i t e s ( s o c a l l e d s l i p s t e p s ) a l l o w s m e t a l d i s s o l u t i o n . D e p e n d i n g o n t h e p o t e n t i a l a n d t h e c h e m i s t r y o f t h e e n v i r o n m e n t a t t h e c r a c k t i p , r e p a s s i v a t i o n t a k e s p l a c e a n d s u b s e q u e n t f i l m r u p t u r e a n d d i s s o l u t i o n c y c l e s l e a d t o f u r t h e r c r a c k i n g o f t h e m a t e r i a l . S t a e h l e p o i n t s o u t t h a t c r a c k a d v a n c e c a n b e n o t i c e d o n l y i n s p e c i f i c e l e c t r o c h e m i c a l c o n d i t i o n s a t t h e c r a c k t i p . T h r e e 6 cases are d i s t i n g u i s h e d : very f a s t r e p a s s i v a t i o n of the s l i p s t e p which g i v e s v i r t u a l l y no incremental advance; very f a s t d i s s o l u t i o n which causes p i t t i n g or crack b l u n t i n g ; a c r i t i c a l r a t i o of d i s s o l u t i o n - a t t a c k to r e p a s s i v a t i o n , at which c o n t i n u e d p e n e t r a t i o n takes p l a c e . U h l i g 1 ^ d e s c r i b e s an a d s o r p t i o n of s u r f a c e a c t i v e s p e c i e s , the consequence of which i s a r e d u c t i o n i n the s u r f a c e energy r e q u i r e d to form a cr a c k , and t h e r e f o r e , a reduced f r a c t u r e s t r e s s . V a r i o u s embrittlement phenomena have been e x p l a i n e d by t h i s model. There i s o f t e n d i s p u t e as to whether the f a i l u r e of a m a t e r i a l i s a c t u a l l y a t t r i b u t e d to SCC or to hydrogen embrittlement. From a thermodynamic p o i n t of view, one c o u l d say that a l l o y s c r a c k i n g p r e f e r e n t i a l l y at anodic p o t e n t i a l s are s e n s i t i v e to SCC, whereas c r a c k i n g i n the c a t h o d i c r e g i o n i s evidence of hydrogen induced c r a c k i n g . But t h i s p o t e n t i a l dependence i s an u n c e r t a i n c r i t e r i o n s i n c e i t i s w e l l known 1^' 2^' 2 1 that crack t i p c o n d i t i o n s ( p o t e n t i a l and pH) may vary c o n s i d e r a b l y from those measured at the bulk m e t a l / e l e c t r o l y t e i n t e r f a c e . P r e s e n t l y , there e x i s t s no s i n g l e coherent theory t h a t e x p l a i n s a l l of the observed s t r e s s c o r r o s i o n phenomena. In many c a s e s , no c l e a r d i s t i n c t i o n between SCC and hydrogen induced c r a c k i n g can be drawn. 7 1.3.2. Co n t r o l l i n g factors for SCC The following factors play important roles in the SCC behavior of rotor s t e e l s : -Chemical composition, concentration, and temperature of the "corrosive" environment. -Electrochemical potential -Material composition The stress l e v e l s and rotor design are also parameters that influence the SCC behavior of a turbine, but these w i l l not be discussed in further d e t a i l . 1.3.3. Environment Laboratory tests have shown that s t e e l s , heat treated to high strength values (>950 MNm~2), are susceptible to SCC in high purity water or steam as well as in aqueous solutions containing various anions , for example, C l ~ , HS~, S0 4 2", HC03~' C0 3 2~ and OH" 10,21-30^ However, i t must be reemphasised that even i f tests are conducted in d i l u t e aqueous solutions, the chemistry inside the crack may be very d i f f e r e n t . This can be due to concentration of damaging species by physical means ( d i f f u s i o n , evaporation of solvent) or by e l e c t r o chemical reactions inside cracks. Also, chemical l e a c h i n g 3 1 of the a l l o y or reactions with impurities in the m e t a l 3 2 can cause l o c a l changes in the envi ronment. 8 For steam turbine materials concentrated solutions of caustics, chlorides and hydrogen sulphide are found to cause cracking3 1/33,34,35_ More recent investigations showed that at very high stress l e v e l s cracking also occurs in steam and pure water 2' 4'36. McMinn et al-* studied the ef f e c t of temperature and caustic concentration on SCC s u s c e p t i b i l i t y and found that increasing temperature lowers the caustic concentration necessary for cracking. At temperatures below 50°C, even in very concentrated solutions, no cracking was found. Shalvoy et al-*? investigated the e f f e c t of various steam impurities and found that such species (eg. Nitrates) can cause an anodic s h i f t in the free corrosion p o t e n t i a l and enhance caustic SCC. Also gaseous impurities l i k e C0 2 can r e s u l t in accelerated c r a c k i n g 2 > -*8 ' 39. 1.3.4. Electrochemical p o t e n t i a l Agrawal et a l 4 ^ 1 did extensive p o l a r i z a t i o n studies on Fe-Ni-Cr a l l o y s in sodium hydroxide solutions. They found that a l l Fe-Ni-Cr a l l o y s exhibit active/passive behavior. They examined the e f f e c t of temperature and caustic concentration on the free corrosion p o t e n t i a l as well as the influence of a l l o y i n g additions on the corrosion resistance. 9 Unlike mild or pl a i n carbon s t e e l s , which are susceptible only over a narrow range of p o t e n t i a l s 1 0 ' 4 1 , rotor steels are susceptible to caustic SCC over a broad range of p o t e n t i a l s 3 4 , 3 7 . Shalvoy et a l 3 7 found the c r i t i c a l p o t e n t i a l ranges in 40%NaOH to be between -850 to ~600mV S H E and above +l50mVg H E. Bandyopadhyay 4 2 found cracking at -350mV S H E and at -750mV S H E in 9M NaOH. The former p o t e n t i a l i s the transpassive region of the po l a r i z a t i o n curve and the l a t e r is near the active to passive t r a n s i t i o n . At the pot e n t i a l of -350mV S H E where SCC of mild steel does not o c c u r 4 3 , the segregation of phosphorus to the prior austenite grain boundaries (temper embrittlement) i s suggested to cause the decrease in SCC resistance in rotor s t e e l s . McMinn et a l 2 did SCC tests at the free corrosion p o t e n t i a l and found cracking in concentrated and di l u t e d sodium hydroxide solutions, pure water and carbonated water. A l l tests were conducted in aerated and deaerated solutions. 1.3.5. Material composition Since the microstructure of most rotor steels i s b a i n i t i c , the emphasis for SCC investigation l i e s on the alloy-composition. A key role for the resistance to SCC seems to l i e in the a l l o y impurities. 10 Bandyopadhyay et a l J 1 ' 4 ^ studied the effect of Mo and P when segregated to the prior austenite grain boundary. They observed severe caustic cracking in the transpassive region with the presence of P and Mo grain boundary segregation, which was explained by the incorporation of Mo and P in the growing oxide f i l m . Both elements enhance the growth k i n e t i c s of the oxide f i l m and decrease i t s d u c t i l i t y . The presence of Mo in P-free steels,however, shows no detrimental e f f e c t . The e f f e c t of s i l i c o n and manganese i s not completely clear yet. Parkins et a l 4 4 suggest that elements that form carbides l i k e T i , Cr and Mn are b e n e f i c i a l to SCC resistance, whereas es p e c i a l l y Si (usually found in the f e r r i t e in preference to the carbides), seems to be detrimental. Also, Bandyopadhyay 4^ relates carbide composition and carbide morphology to the SCC resistance. McMinn et a l ^ found that the lowering of the n i c k e l content in rotor steels gives an improvement in SCC resistance. However, t h i s e f f e c t was only observed in s p e c i f i c environments and cannot be used as a general conclusion. They suggest that the y i e l d strength of the material has a greater e f f e c t on SCC behavior than the a l l o y composition. Greenfield et al^6 tested rotor steels with varying Ni, Mo and V contents and found no s i g n i f i c a n t difference in SCC behavior. Besides the influence of microsegregation at grain boundaries, Parker^2»46 points out the e f f e c t of sulphide inclusions on the SCC behavior of s t e e l s . Sulphides act as major cracking i n i t i a t o r s p a r t i c u l a r l y in a c i d i c solutions. The composition of the inclusions i s very important because i t determines th e i r s o l u b i l i t y in a c e r t a i n environment. Parker found that inclusions have a pronounced e f f e c t in very d i l u t e environments whereas in aggressive environments they may be of secondary importance. However, Roberts-*** tested rotor steels in pure steam and water and found no difference in cracking behavior of samples taken from heavily segregated areas (centre portion of forging) and samples taken from zones with no segregation. At present, there i s only very l i t t l e agreement about the influence of d i f f e r e n t a l l o y i n g or impurity elements on the SCC of rotor s t e e l s . The p r i n c i p a l objective was to study and c o r r e l a t e the SCC behavior of three d i f f e r e n t rotor steels with their electrochemical p o l a r i z a t i o n behavior in environments that simulate the steam condensates present in steam turbine operation i . e . carbonate and sodium hydroxide solutions. A second objective was to co r r e l a t e inclusions, p r i n c i p a l l y CaS and MnS in remelted and doped rotor s t e e l , with corrosion behavior of the s t e e l . In t h i s manner, new information may be generated to a s s i s t in the mechanistic understanding of reported f a i l u r e s of cracked steam turbine rotors and SCC results obtained in other laboratories. I I . EXPERIMENTAL 2.1. Material For our experiments three d i f f e r e n t rotor steels were av a i l a b l e , their designation i s "A", "R" and "0" - s t e e l . 2.1.1. Chemical composition of the materials Steel "A" i s a 3.5%NiCrMoV - st e e l with very low concentration of impurity elements. Also, the manganese and s i l i c o n content are kept considerably below the regular concentrat ion. The reference s t e e l "R" has a composition of a regular 3.5%NiCrMoV - st e e l used in commercial LP steam turbines. According to the ASTM Norm 4 7 i t represents an A470-Class5 rotor s t e e l . The 1%CrMoV steel "0" i s c l a s s i f i e d after the ASTM Norm as A470-Class8 rotor s t e e l . ( T y p i c a l for HP turbines) Figure 1 on page 14 gives the composition of a l l three s t e e l s . A l l concentrations are given in wt% and are provided by the producers of the m a t e r i a l s 1 ^ . 14 F i g u r e 1: Chemical composition (wt%) of the r o t o r s t e e l s Steel-A 3.44 0.3 1.68 0.41 10.09 0.02 0.04 0.05 C Ni Cr Mo V y h 0005 0005 ,001^ 0003 0002 0001 0.002 Mn Si Cu Al Sb W Sn As S P Steel-R 3.45 0.3 1.42 Q33 0.1 Q29 0.07 0.06 i | | |Q0Q5 0.QC3 r^no.006 Q.QQ6 0.002 0004 C Ni Cr Mo V Mn Si Cu Al Sb W Sn As S P 002 Steel-0 021 0.35 1.15 Q94 1 C Ni Cr Mo Q67 n 025 H r Q25 0.05 IQ006 Mn Si Cu Al 0.005 0.007 0007 Sn 15 2.1.2. Heat treatment The heat treatment of the steels "A" and "R" given below was conducted by the manufacturer and was not further changed for our investigations. Steel "0" was received in the h o t - r o l l e d and air-cooled condition and was not further heat treated for our experiments. Figure 2: Heat treatment of steels "A" and "R" Steel-A ac: air cool ing sq: s p ray quenching (water) 16 2.1.3. Sampling of rotor material Steel "A" and "R" were received as discs which were sectioned from big rotor forgings. Steel "0" was av a i l a b l e as a long round shaft. Figure 3 below shows the locations where specimens were sectioned from the forgings. The descr i p t i o n of the notched samples i s as follow: F i r s t l e t t e r : d i r e c t i o n normal to the fracture plane Second l e t t e r : d i r e c t i o n of crack propagation "L": length, d i r e c t i o n of metal working "C": chord of c y l i n d r i c a l cross section Figure 3: Sampling map Specimen: Steel 0 T? Fracture mechanics i==c Charpy-V Slow strain rate 0 "Surface "Centre" \ .Microstructure, Hardness 0-115mm 18 2.1.4. Mechanical properties Tensile-, hardness-, and charpy-V-notch tests were performed on a l l three steels at room temperature. To check the influence of temperature on the p l a s t i c deformation behavior, (slow) t e n s i l e tests were also conducted at 100°C and 160°C. (A hot o i l bath was chosen as the inert test medium) Table I below summarizes the results of the mechanical t e s t s . Table I: Mechanical properties Material/Temp. Y i e l d S t r e n g t h (MPa) •1) U l t . T e n s i l e S t rength (MPa) *1) Elongat i o n to F a i l u r e (50 *1) Hardne ss ( c e n t r e ) <v *2) Hardness ( s u r f a c e ) (R ) c *2) Charpy-V (J) *3) F r a c t u r e Toughness (MNm" 3 / 2) I | CC iCL|LC Steel "A" 3.5%NiCrMoV , i 1 1 1 1 RT 100°C 160°C 725 645 637 853 771 779 17.1 15.0 15.8 22.5 26.0 1 I 1 3311 33i --1 i 250 Steel "R" 3.5%NiCrMoV 1 , i ; i • ; i RT 1 00°C 160°C 727 712 670 838 810 797 16.7 14.0 15.7 25.0 26.0 1 37*1 37 1 __ __ — _ _ _ _ | 250 Steel "0" 1%CrMoV RT 100°C 160°C 931 924 839 1 100 1072 933 13.2 11.0 11.0 34.7 35.6 i 1 1 1 1 ! 6! 6| 20 ~ ~ i ~~r~ i " I » cc CL'LC 43 ,'43 | 55 i : • i 19 *1) For each temperature only one test was performed. The samples had a gage length of 25.4mm and a diameter of 4.0mm. Tensile tests at room temperature were run at £ =3.3 * 1 0 _ 4 s ~ 1 . At 100°C and 160°C e was 3.3'10" 5S" 1. *2) Hardness measurements were taken at locations indicated on the sampling map (page 16,17) *3) With steel "0" the toughness was measured in the CL, CC and LC d i r e c t i o n s . Because of the thin disc shape only CL and CC directions were investigated on steel "A" and "R". *4) Fracture toughness values of steel "A" and "R" are l i t e r a t u r e d a t a b U . Steel "0" values were obtained by breaking SCC samples in a i r . Steels "A" and "R" showed very similar properties. The hardness of "A" in the centre region was lower than in the R- s t e e l . This was due to a d i f f e r e n t cooling rate of each forging. The 1%CrMoV s t e e l "0" showed higher y i e l d and ultimate t e n s i l e strength but a l i t t l e less d u c t i l i t y . However, the toughness of "0" was very much lower than in the two NiCrMoV-materials. This correlated with the considerably higher hardness of the 1%CrMoV s t e e l . The toughness of "0" was strongly dependent on the d i r e c t i o n of the fracture plane. This observation can be explained with the presence of a r o l l i n g texture and the elongated manganese sulphides in the r o l l i n g d i r e c t i o n . For the steels "A" and "R" no d i r e c t i o n dependence of the toughness was assumed since neither r o l l i n g texture nor elongated inclusions were present. With increasing test temperature, both y i e l d strength and ultimate t e n s i l e strength decreased. However, t h i s could 20 also be due to the faster s t r a i n rate at room temperature. The e f f e c t of temperature on the percentage elongation to f a i l u r e was rather minimal. ( More tests would be necessary to confirm these r e s u l t s . ) On the stress s t r a i n curve, no s t r a i n aging phenomena was observed during p l a s t i c deformation at the elevated temperatures investigated. 2.1.5. Production of CaS and MnS r i c h steels .1.5 kg charges of the ordinary 3.5%NiCrMoV - st e e l "R" were remelted in a vacuum induction furnace. Sulphur was added to the charge in the form of FeS powder. The melting was conducted in an argon atmosphere (1.1atm). Calcium granulates and Manganese oxide powder were immersed into the superheated melt («M600°C) with a plunger. After the Calcium or Manganese additions, the steel was cast into a c h i l l e d s t e e l mold. Additions per charge: For CaS-rich s t e e l : 20gms Ca, 1gm FeS For MnS-rich s t e e l : 20gms MnO, 1gm FeS The casting was hot r o l l e d at 950°C. In 8 passes a 50% reduction was achieved. Afterwards, the heat treatment previously described for st e e l "R" (page 15) was conducted. 21 2.2. Microstructure / carbides / inclusions For viewing the very fine microstructure of the b a i n i t i c s t e e l s , both a l i g h t microscope and a transmission electron microscope (TEM) were used. The locations for the microstructural investigations are indicated on the sampling map (page 16). 2.2.1. Optical microscope investigations The samples were ground with SiC paper down to 800 g r i t . For the f i n a l p o l i s h i n g f i n i s h , diamond paste down to 1/4 urn grade was used. When inclusions were examined, alcohol was used as polish i n g lubricant and coolant instead of water. The microstructure was revealed with a p i c r i c acid etchant: 200ml ethanol, 20gm p i c r i c acid and a few drops of cone. HCL. Examination and photography were conducted with a "Zeiss-Ultraphot" o p t i c a l microscope. 22 2.2.2. Steel "A" (3.5%NiCrMoV super clean) The microstructure of the rotor forging was investigated at two locations. One was a few centimeters below the surface and the other one 23 centimeters further towards the centre. Typical microstructures are shown in Figures 4 and 5 below. Figure 4: Microstructure of s t e e l "A", centre region Bainite, carbide p a r t i c l e s in f e r r i t e matrix Mag.: 600x, etched Figure 5: Microstructure of s t e e l "A", surface region Mag.: 600x, etched 23 From the Figures 4 and 5 (pg. 22), i t can be seen that the b a i n i t i c microstructure was e s s e n t i a l l y the same at both locations. Also, no difference was observed between the r a d i a l and the longitudinal d i r e c t i o n at the same locations. However, i t was observed that in the centre region the carbides were more numerous and coarser. This could be explained with the faster cooling rate at the surface region. Neither P i c r i c acid nor N i t a l etching revealed p r i o r austenite grain boundaries. No further metallographic experiments were undertaken to determine the p r i o r austenite grain s i z e . However, from fractography studies, a size below 30 jum could be estimated. From the production report 4** of t h i s s t e e l , i t i s further known that throughout the whole rotor cross-section a b a i n i t i c microstructure was present. 2.2.3. Steel "R" (3.5%NiCrMoV) S i m i l a r l y to s t e e l "A", no difference in microstructure was observed between regions near the surface and 23 cm. below the surface towards the centre. According to Figures 6 and 7 (page 24), i t also seems that the size and d i s t r i b u t i o n of carbides were very much the same in both locations. This observation i s supported by the hardness measurements which also showed no difference in the two regions. Even with d i f f e r e n t a l l o y compositions and heat 24 treatment, both steels "A" and "R" appeared to have a very similar microstructure. Figure 7: Microstructure of s t e e l "R", surface region Mag.: 600x, etched 2.2.4. Steel "0" (l%CrMoV) According to the continuous cooling transformation diagram 4 9 of a 1%CrMoV-steel (similar to the one investigated), a i r cooling of the rotor s t e e l "0" should have resulted in a b a i n i t i c microstructure with possibly some f e r r i t e present. Photomicrographs taken at d i f f e r e n t locations in the material again did not show any major difference in structure between surface and centre regions. However, at low magnification a s l i g h t r o l l i n g texture could be observed in the longitudinal d i r e c t i o n . (See Fig.8) At higher magnification a coarser and less homogeneous carbide d i s t r i b u t i o n i s v i s i b l e compared to the one found in the "A" and "R" s t e e l s . (See Fig.9) Figure 8: Microstructure of s t e e l "0", radi a l d i r e c t i o n with r o l l i n g texture Bainite, bright areas are f e r r i t e , darker areas carbides + f e r r i t e , r o l l i n g texture v i s i b l e as horizontal bands ( s l i g h t l y darker) Mag.: 120x, etched 26 Figure 9: Microstructure of st e e l "0" Coarse carbides and f e r r i t e Mag.: 600x, etched 2.2.5. Transmission electron microscope (TEM) investigations Discs for TEM studies were sectioned from unstrained slow s t r a i n rate test pieces using a spark machining technique. The discs were 3mm in diameter and 2mm thick and were c a r e f u l l y ground on SiC paper to a thickness of about 0.1mm. (It was t r i e d to keep the mass of the samples as small as possible because of their disturbing e f f e c t on the magnetic f i e l d in the microscope lens system). 27 The f i n a l electrochemical thinning was conducted with a "Struers polipower" jet polishing device. A mixture of 20 percent perchloric acid in methanol was used as e l e c t r o l y t e . The potential was held at about 20-22 Volts with a medium flow rate s e t t i n g . The e l e c t r o l y t e temperature was maintained at 25°C. Examination and photography were conducted with a Hitachi H800 scanning transmission electron microscope. For a q u a l i t a t i v e chemical analysis of matrix and carbides, an Ortec - energy dispersive X-ray (EDX) analyzer was used. 2.2.6. Steels "A" and "R" In the l i g h t microscope steel "A" and "R" showed the same microstructure. However, the TEM studies revealed some differences. The Figures 10-13 on page 29 ( a l l electron micrographs taken at about the same magnification) show that the f e r r i t e grains in steel "R" are very narrow and needle-l i k e , whereas in "A" a rather irregular shape i s present. Also, the carbides in "R" seem to vary in size more than in st e e l "A". The Figures 14 and 15 on page 30 show that inside the f e r r i t e subgrains of st e e l "R", carbides are p r e c i p i t a t e d at an angle to the longitudinal d i r e c t i o n . In both s t e e l s , bigger carbides were often located along subgrain boundaries 28 and possibly along p r i o r austenite boundaries.(Figures 16 and 17) According to the 1 i t e r a t u r e ^ 1 ' ^ 2 ' ^ 3 t steel "R" represents the microstructure of a "lower b a i n i t e " whereas "A" has features t y p i c a l of "upper ba i n i t e " . 2.2.7. Steel "0" This steel had a much less uniform microstructure compared with "R" and "A". On Figure 18 (pg. 32), bainite and f e r r i t e are v i s i b l e . The very small p a r t i c l e s in the f e r r i t e may be carbides or could be some sort of contamination from the polishing solution. The ele c t r o p o l i s h i n g of t h i s s t e e l was not as s a t i s f a c t o r y as for the other two s t e e l s . The structure of the ba i n i t e would suggest "upper b a i n i t e " . On the boundaries between f e r r i t e and b a i n i t e rather large, elongated carbides were present. (See Figures 19,20 on page 32) 29 F i g u r e 10: TEM-picture, m i c r o s t r u c t u r e of s t e e l "A" with c a r b i d e s and subgr a i n s B a i n i t e , b r i g h t areas are f e r r i t e , dark p a r t i c l e s c a r b i d e s Mag.: 4,000x F i g u r e 12: TEM-picture, m i c r o s t r u c t u r e of s t e e l "R" with e l o n g a t e d subgra i n s F i n e , needle l i k e and round c a r b i d e s , b r i g h t area i n the c e n t r e i s a hole i n the sample Mag.: 4,000x Figu r e 11: TEM-picture, m i c r o s t r u c t u r e of s t e e l "A" with c a r b i d e s and subgrains Mag.: 3,500x F i g u r e 13: TEM-picture, m i c r o s t r u c t u r e of s t e e l "R" with c a r b i d e s and subgrains Mag.: 4,000x 30 Figure 14: TEM-picture, steel "R", elongated carbides inside subgrains Small carbides are not aligned with the long subgrain axis Mag.: 12,000x Figure 15: TEM-picture, steel "R", carbides inside subgrains and at grain boundaries Mag.: 9,000x 31 Figure 17: TEM-picture, steel "A", elongated carbides at subgrain boundar ies Carbides aligned with the long subgrain axis, fine dark l i n e s are d i s l o c a t i o n s Mag.:40,000x 32 F i g u r e 18: T E M - p i c t u r e , s t e e l " 0 " , m i c r o s t r u c t u r e w i t h f e r r i t e a n d b a i n i t e C a r b i d e s a r e a l i g n e d i n o n e d i r e c t i o n i n s i d e t h e b a i n i t e g r a i n . S m a l l p a r t i c l e s i n s i d e t h e f e r r i t e g r a i n may n o t b e c a r b i d e s , b u t c o u l d b e a c o n -t a m i n a t i o n f r o m t h e e l e c t r o -p o l i s h i n g M a g . : 8 , 0 0 0 x F i g u r e 1 9 : T E M - p i c t u r e , s t e e l " 0 " , b i g c a r b i d e s a t s u b g r a i n b o u n d a r i e s F e r r i t e a n d b a i n i t e g r a i n s M a g . : 5 , 0 0 0 x F i g u r e 2 0 : T E M - p i c t u r e , s t e e l " O " , c a r b i d e s i n s i d e g r a i n s a n d a t g r a i n b o u n d a r i e s B i g e l o n g a t e d c a r b i d e s a t t h e g r a i n b o u n d a r y t o t h e l e f t M a g . : 6 , 0 0 0 x 33 2.2.8. Qu a l i t a t i v e carbide analysis by EDX A l l carbides were iron carbides with d i f f e r e n t contents of a l l o y i n g elements. In steel "A" and "R" no major difference in composition could be detected between big carbides located at the grain boundaries and small carbides in the i n t e r i o r . However, a l l carbides in "A" and "R" seemed to be enriched in the al l o y i n g elements Mo, V and Cr. In steel "0", most carbides were enriched in Cr only, but to a lesser degree than in steel "A" and "R". However, the long, big grain boundary carbides of "0"-steel were enriched in Cr and Mo. Caution must be exercised, however, because these X-ray analyses were only q u a l i t a t i v e and d i d not allow a more detailed i n t e r p r e t a t i o n . Figures 21 and 22 below show an example of an EDX analysis on steel "R". 34 F i g u r e 22 : ED X - a n a l y s i s , s t e e l "R", matrix and c a r b i d e s . Q u a l i t a t i v e a n a l y s i s , peak h i g h t not d i r e c t l y p r o p o r t i o n a l to element c o n t e n t s . Enrichment of Mo, V, and Cr i n the c a r b i d e s P R E S E T : L T = 1 0 0 $ D E A D T I M E : 9'/. C O U N T S / S E C O N D : 4<5 F U L L S C A L E : IK L I N E A R 0 - 1 8 K E M 10 E L V C H M A T R I X 1 F e K a 1 1 F e K p fl NiK 0 . 0 0 2 . 5 « 5 . 1 2 7.<8 1 0 . 2 3 P R E S E T : L T * 1 8 0 S D E A D T I M E : 1 3 K C O U N T S / S E C O N D : 7 8 0 F U L L S C A L E : IK L I N E A R t 8 - 1 8 K E U 10 E U / C H C A R B I D E J C r K« F e K 4 M o L <t U K 8 . 8 0 2.Si 5 . 1 2 7.it 1 0 . 2 3 35 2.2.9. Electron d i f f r a c t i o n of carbides Electron d i f f r a c t i o n studies were attempted in order to characterize the carbides in more d e t a i l . However, because of the very strong e f f e c t of the sample matrix on the magnetic f i e l d of the microscope lenses i t was very d i f f i c u l t to analyze single carbides with the t i l t specimen holder. To avoid these problems, the extraction r e p l i c a method would have been preferable but was not employed in t h i s project. Attempts were made to obtain ring patterns from areas containing a multitude of carbides with l i t t l e success, due to the variety of carbides paterns which were present. Figures 23 and 24 i l l u s t r a t e patterns from single carbides and multiple carbides. In Figure 23 the pattern cannot be indexed as Fe3C, but may represent an orthorhombic structure analogous to Cementite. Figure 23: Electron d i f f r a c t i o n pattern of a carbide, s t e e l "R" Carbide o(Fe Fe-oxides 36 Bandyopadhyay et a l 4 5 conducted studies on the composition of carbides in rotor s t e e l s . In a tempered b a i n i t i c microstructure they found M 3 C , M 7 C 3 and M2C carbides. The M 3 C type were bigger carbides p r e f e r e n t i a l l y located along grain boundaries. The M 7C 3 and M2C p a r t i c l e s were found to be very small with diameters below 0.1 jjm. The M 3 C were found to be iron carbides enriched in Cr up to 25%, Mo up to 4% and V up to about 2%. The M 7C 3 contained mainly iron and chromium, whereas M2C represented Mo,V-carbides. This enrichment of a l l o y i n g elements in the carbides supports our q u a l i t a t i v e findings with the EDX analys i s . 37 2.2.10. Inclusions Since s p e c i f i c a t i o n s for rotor steels are very t i g h t on the content of impurity elements such as sulphur, the inclusions are very small and not very numerous. In t h i s study no attempt was made to f u l l y characterize the d i s t r i b u t i o n and size of a l l inclusions present in the test materials. However, attempts were made to characterize t y p i c a l inclusion types present in each of the three s t e e l s . Steel A; The big inclusions were found to have a diameter of about lO^im. Most inclusions, however, were smaller than 3jjm. At a few locations, very long and narrow p a r t i c l e s were detected with a length up to 50,um. The shape of most inclusions was quite i r r e g u l a r with sharp corners. Round p a r t i c l e s were rather seldom. Qu a l i t a t i v e investigations with an EDX analyzer attached to a scanning electron microscope (SEM) revealed the following data: -Small round inclusions were calcium sulphides. Sometimes aluminum or s i l i c o n was detected in the centre of these p a r t i c l e s , which most probably represents an oxide core (often observed in CaS - i n c l u s i o n s ^ 4 ) . See Figure 25. -Inclusions with sharp corners (often trian g u l a r shape) were composed mainly of S i , Al and Fe. Therefore they could be c l a s s i f i e d as oxides. See Figure 26. -Long flattened p a r t i c l e s could also represent oxides becouse they were also r i c h in s i l i c o n . See Figure 27 Figure 25: Di s t r i b u t i o n of inclusions in steel "A" Average inclu s i o n appearance in the centre region F i g u r e 2 6 : O x i d e i n c l u s i o n s i n s t e e l "A" M a g . : 1 , 0 0 0 x F i g u r e 2 7 : O x i d e a n d s u l p h i d e i n c l u s i o n s i n s t e e l "A" R o u n d p a r t i c l e i s C a S M a g . : 1,OOOx 40 S t e e l R: C o m p a r e d t o t h e "A" s t e e l , "R" s e e m e d t o h a v e f e w e r b u t b i g g e r i n c l u s i o n s . T h e i r s h a p e was m o s t l y r o u n d w i t h d i a m e t e r s b e t w e e n a b o u t 10 a n d 20 urn. A l s o , a f e w s m a l l e r i r r e g u l a r s h a p e d p a r t i c l e s c o u l d b e f o u n d . W i t h t h e E D X a n a l y s i s , t h e b i g r o u n d i n c l u s i o n s c o u l d b e c l a s s i f i e d a s C a S w i t h a n a l u m i n a / s i l i c a c o r e . T h e i r r e g u l a r s h a p e d p a r t i c l e s s e e m e d t o b e e x c l u s i v e l y s i l i c a i n c l u s i o n s . ( S e e F i g u r e s 2 8 - 3 0 ) F i g u r e 2 8 : F i g u r e 2 9 : C a S i n c l u s i o n s i n s t e e l "R" O x i d e i n c l u s i o n s i n s t e e l "R" D a r k a r e a i n t h e c e n t r e o f t h e M a g . : 1 , 0 0 0 x i n c l u s i o n i s a n o x i d e c o r e M a g . : 1,OOOx 4 1 Steel 0: A completely di f f e r e n t inclusion d i s t r i b u t i o n was found in t h i s s t e e l . In the r o l l i n g d i r e c t i o n long stringers of flattened inclusions were observed. Their length was generally 50um but stringers composed of several aligned inclusions were a few hundred microns long. Their thickness was mostly below lOum. In addition, only a few round inclusions could be found and the i r size was smaller than a few microns. Irregular shape oxide inclusions could only very rarely be detected. The composition of the elongated inclusions c l a s s i f i e d them as MnS. In a few MnS str i n g e r s , small p a r t i c l e s (probably oxides) of Al and Si were found. (See Figures 31 and 32, pg. 42) Figure 31: Elongated MnS steel "0" inclusions in Mag.: 1,000X Figure 32: Small round inclusions in steeel "0" Not i d e n t i f i e d as sulphides or oxides Mag.: 1,000x Remelted s t e e l "R" doped with inclusions, MnS and CaS. Figure 33: \ ; v i , .1.. Microstructure and inclusions in CaS doped steel "R" Black p a r t i c l e s are etched away inclusions, small p a r t i c l e s are carbides, most probably a b a i n i t i c microstructure i s present Mag.: 600x Figure 34: Microstructure and inclusions in MnS doped steel "R" Big p a r t i c l e s are etched away inclusions, small p a r t i c l e s are carbides, microstructure is not c l e a r l y i d e n t i f i e d , (hardness i s much lower than in the CaS doped steel) Mag.: 600x 44 2.2.11. Summary of microstructures In "A" and "R" st e e l s , a f u l l y b a i n i t i c microstructure was observed. The 1%CrMoV-steel "0" showed a f e r r i t i c / b a i n i t i c structure. Steel "A" showed a very uniform d i s t r i b u t i o n of fine carbides whereas bigger carbides were found along the grain boundaries in steel "0". The iron carbides in the 3.5%NiCrMoV-steels "A" and "R" were enriched in the al l o y i n g elements V,Cr and Mo. Steel "0" carbides showed only an enrichment in Cr, but to a lesser degree than in "A" and "R". Inclusions in steel "A" were mainly small oxides with a few CaS p a r t i c l e s . Steel "R" contained quite big CaS inclusions, but fewer oxides. In the hot r o l l e d steel "0", stringers of MnS inclusions were detected. 45 2.3. Electrochemical P o l a r i z a t i o n studies Under free corrosion conditions (the metal specimen i s immersed in a corrosive l i q u i d and no external voltage i s applied), both reduction and oxidation reactions occur on the metal surface. Therefore, both anodic and cathodic currents are present and equal in magnitude. For studying corrosion processes i t i s often advantageous to separate anodic and cathodic reactions. This can be done by use of a voltage source to force the metal specimen to a potential other than the free corrosion p o t e n t i a l . As a result of t h i s " p o l a r i z a t i o n " , depending on the p o l a r i t y , either the anodic or cathodic current predominates on the metal surface. Experimentally one measures the corrosion current as a function of the applied p o t e n t i a l measured with respect to a standard reference electrode. A record of p o t e n t i a l versus log current i s then c a l l e d a p o l a r i z a t i o n curve. Figure 35, on page 46 i l l u s t r a t e s how an experimentally obtained p o l a r i z a t i o n curve can be explained as the synthesis of oxidation and reduction reactions. 46 Figure 35 : I l l u s t r a t i o n of a t h e o r e t i c a l and experimental p o l a r i z a t i o n curve Theoretical cuves Experimental curves (Polarization curves) log current-density log current-density The position of oxidation and reduction curves i s dependent on variables l i k e temperature, concentration of species in solution or aeration and a g i t a t i o n of the soluti o n . Therefore, the interse c t i o n point(s) of the oxidation curve with the reduction curve change(s) according to these parameters. This means that the free corrosion p o t e n t i a l of a certain metal i s strongly dependent on the environment conditions. Therefore, by imposing a c e r t a i n p o t e n t i a l onto the metal (in a given environment), one can simulate the free corrosion behavior of the same metal for a d i f f e r e n t environmental condition. 47 2.3.1. Sample preparation, test set up Rods of approximately 4mm diameter and a length of 12cm were machined from a l l three materials. The samples were ground with SiC paper down to 800 g r i t and then electropolished in chromic-acetic a c i d . (1330ml acetic acid, 250 gms chromium t r i o x i d e , 70 ml H 20). For the test the samples were wrapped in PTFE-TefIon tape so that only the bottom section with an approximate surface area of 2.5 cm2 was exposed to the environment. Before each t e s t , the samples were cleaned with acetone and alcohol. The test c e l l (see Fig.36, pg. 49) was a 600ml PTFE beaker f i t t e d with a PTFE l i d . A heating mantle coupled with a temperature c o n t r o l l e r allowed adjustment of the solution temperature to - 1°C. An approximate 3cm2 platinum sheet in the solution served as counter electrode. A PTFE Luggin c a p i l l a r y was connected v i a a saturated KC1 s a l t bridge to a standard calomel reference electrode. The reference electrode was kept at room temperature(23°C). A l l solutions were made with reagent grade chemicals and d i s t i l l e d water. The solutions were purged with USP nitrogen before and during the t e s t , except for tests with carbonate solutions, where C0 2 was bubbled through the sol u t i o n . A l l tests were conducted at 95°C. 48 Investigations were conducted in the following solut ions: C0 2+H 20 ,p(C02)=1 atm C02+H20+NaHC03 ,p(C02)=1 atm , c(NaHC03)=3.5•10"3M to IM H20+NaHC03+Na2C03 ,c(NaHC03)=1M , c(Na 2C0 3)=1M H20+NaOH ,c(NaOH)=0.35M, 3.5M The pH range for the carbonate solutions at 95°C was from pH =4.5 for the C0 2/H 20 solution up to about 9.5 for the 1 molar carbonate/bicarbonate solution. The sodium hydroxide solutions were in the pH range of 12-13. A l l p o l a r i z a t i o n curves were obtained with a Princeton Applied Research Model 350A Corrosion Measurement System. The p o t e n t i a l scan rate was either 1mV/sec or 0.3mV/sec. A l l tests represent anodic p o l a r i z a t i o n curves (scanning from negative to po s i t i v e p o t e n t i a l s ) . Each test was started at a potential at least l50mV more negative than the free corrosion p o t e n t i a l in order to remove any oxide f i l m on the sample surface. No attempt was made to correct for po t e n t i a l differences a r i s i n g from l i q u i d junction p o t e n t i a l s in the s a l t bridge. (The correction would be in the order of a few m i l l i v o l t s ^ ) . 49 Figure 36 : Test set up for p o l a r i z a t i o n studies G 0 IEEE X R ) 1 VJ u Solution (95°) / ' / .-• / / / / / / / / / / / / / H. / / / / / / / /// / /. G Th 0 C E R P H L Gas i n l e t Thermometer Gas outlet Pt-counter electrode Working electrode Calomel reference c e l l Temperature c o n t r o l l e r Heater Luggin c a p i l l a r y 50 2.4. Slow s t r a i n rate t e n s i l e test The slow s t r a i n rate t e n s i l e test (SSRT) i s often applied for determination of SCC s e n s i t i v i t y . The advantage of t h i s method i s the r a p i d i t y with which the test results can be obtained. The s u s c e p t i b i l i t y to SCC i s usually indicated by a reduction in mechanical properties, eg. lesser degree of necking before f a i l u r e and the presence of secondary cracks in the gage section. From the length of the secondary cracks and the test duration, average crack v e l o c i t i e s can be calculated. In the present study only the time spent during p l a s t i c deformation was considered. 2.4.1. Sample preparation, test set up A l l slow s t r a i n rate t e n s i l e tests were conducted with a v e r t i c a l l y mounted Hounsfield tensometer f i t t e d with a reduction gear and 12 rph synchronous motor. A crosshead speed of 4.4•10~^mm/sec corresponded to a s t r a i n rate of 1.7-10~6 sec" 1. Rods of 25.4cm (10") length and 9.5mm (3/8") diameter were machined according to the sampling map (page 16). The gage section in the centre of the rod had a diameter of 4mm and a length of 25.4mm. Both ends of the t e n s i l e samples were threaded for mounting in t e n s i l e grips. 51 The gage section was ground with SiC paper and subsequently electropolished in chromic acetic a c i d . During the t e s t , the sample was wrapped with PTFE tape and only the gage section was exposed to the solution. The test c e l l (see F i g . 37, pg. 52) was a 600ml PTFE beaker, f i t t e d with a rubber l i d . A heating tape wrapped around the beaker coupled with a temperature c o n t r o l l e r allowed adjustment of the temperature to ±1°C. A platinum sheet served as counter electrode. A PTFE Luggin c a p i l l a r y was connected via a s a l t bridge to a standard calomel reference electrode in the same manner as during the p o l a r i z a t i o n tests. Solutions were made up as described in 2.3.1. and the tests were a l l performed at 95°C. The testing potentials were chosen according to the r e s u l t s of the p o l a r i z a t i o n studies. Tests were conducted at the free corrosion p o t e n t i a l , near and at the active peaks and in passive regions. The potential was co n t r o l l e d with a Wenking Potentiostat, Model 70HP10 or an ECO Model 549 respectively. The usual test duration to f a i l u r e was around 24 hours, depending on the s t e e l and the environment. After the test the two samples halves were cleaned in i n h i b i t e d a c i d (3ml HC1, 4ml 2-Butyne-1,4 d i o l , 50ml H 20) and one sample piece was e l e c t r o l y t i c a l l y coated with nickel 52 and sectioned l o n g i t u d i n a l l y (coating protects surface edge during mechanical p o l i s h i n g ) . On the polished l o n g i t u d i n a l cross section, the corrosion attack and secondary crack length could be examined. The other half of the broken sample was used for scanning electron microscope (SEM) investigations on an ETEC-Autoscan. Figure 37 Set up for slow s t r a i n rate t e n s i l e tests =-(R) G Th 0 C E R P H L Gas i n l e t Thermometer Gas outlet Pt-counter electrode Working electrode Calomel reference c e l l Temperature c o n t r o l l e r Heater Luggin c a p i l l a r y 53 2.5. Fracture mechanics test Results from p o l a r i z a t i o n and slow s t r a i n rate studies are often used as guidelines for the set up of long term fracture mechanics t e s t s . With a fracture mechanics test a more r e a l i s t i c simulation of SCC can be done. Instead of a smooth sample, such as in the SSRT studies, a prenotched and fatigue pre-cracked specimen i s used. This eliminates the problems of crack i n i t i a t i o n and allows loads to be employed below the macroscopic y i e l d strength. The characterization of the stress state at the crack t i p i s done by defining the so c a l l e d stress i n t e n s i t y factor "K". This parameter combines nominal stress with the crack length and a correction factor accounts for the specimen geometry^. The stress i n t e n s i t y factor i s , therefore, a very useful parameter for the comparison of various samples with d i f f e r e n t crack length and loading conditions. 2.5.1. Specimen design The most common specimen geometries used for SCC studies in the corrosion laboratory at U.B.C. are the T-notch double c a n t i l e v e r beam (TN-DCB) 5 7 and the double c a n t i l e v e r beam (DCB)^8. Both samples can be s e l f stressed (wedge or bolt loaded) or can- be connected to a load c o n t r o l l i n g device. 54 For t h i s study, an easy-to-machine cantilever beam type of specimen was designed. By bolting two specimens together a test sample similar to the T-notch double cantilever beam can be made up and allows two cantilever bend specimen to be tested simultaneously. Bars of size 101.6 x 25.4 x 25.4mm (1"X1"X4") were cut out according to the sampling map (page 16) . The notch was produced by d r i l l i n g a hole (3/16" in diameter) through the sample and removing the remaining metal between the hole and top face of the specimen. This produced a 3/16" wide s l o t . A fine jewellers-saw cut was then made along the bottom of the s l o t about 1mm deep. F i n a l l y a razor blade scratch produced the i n i t i a t i o n s i t e for the fatigue pre-crack. (See Fig.38) Figure 38 : Specimen geometry for fracture-mechanics tests 55 The specimen were fatigue pre-cracked by c y c l i c bending on a Sonntag SF-1-U fatigue machine. The applied load was selected so that the stress i n t e n s i t y at the crack t i p was at least 5-10 MNm~3/2 below the stress intensity applied during the SCC test. The stress i n t e n s i t y for our sample geometry was calculated using a formula described by Brown^. (see F i g . 39d, pg. 57) The same rel a t i o n s h i p was used by Robinson and S c u l l y ^ to do SCC experiments with a notched cantilever beam. However, t h i s formula assumes pure bending of a cant i l e v e r beam which i s not f u l l y s a t i s f i e d in our set up. The maximum bending moment that can be applied i s con t r o l l e d by the maximum t e n s i l e f i b e r bending stress. The analysis i s v a l i d only below the y i e l d point of the material. The r e l a t i o n s h i p between t e n s i l e f i b e r bending stress and bending moment i s given by 6" = Mx/I ...(1) M = bending moment; x = distance from neutral f i b e r ; I = moment of Ine r t i a Applied to our cant i l e v e r beam : <3 = (6M)/B(W-a)2...(2) Calculations for our case have shown that at stress i n t e n s i t i e s above about 50MNm"3/2 for steel "A" and "R" t h i s r e s t r i c t i o n of the maximum applied bending moment i s viol a t e d . . Therefore, l i n e a r e l a s t i c fracture mechanics should not be applied any longer to calculate the exact absolute stress i n t e n s i t y at the crack t i p . However, a comparison between d i f f e r e n t steels of the same geometry can s t i l l be done. In Figures 39a - 39d d i f f e r e n t methods for the determination of the stress i n t e n s i t y applied to our type of specimen geometry are i l l u s t r a t e d . Figure 39 : Different specimen geometries for fracture mechanics tests 57 Figure 39 : cont'd c) T-DCB (Russel Tromans57) d) W B Cantilever Beom (Brown 59) 6 M r~ K = , f a Y BW2 .. (5) Y = M = P L BW^ PL 2 ...(6) In Figure 40 (pg. 58) a l l four methods are applied to our specimens geometry and stress i n t e n s i t i e s are calculated for d i f f e r e n t crack length. One reason why the curves d i f f e r from each other i s the r e s t r i c t e d use of each c a l c u l a t i o n method to s p e c i f i c sample dimensions. This i s due to the experimental c a l i b r a t i o n of the "geometric factor" of each specimen type. Since the aim of th i s project was to compare d i f f e r e n t steel q u a l i t i e s under the same testing conditions no attempt was made to c a l i b r a t e our sample geometry with the compliance method. ( i t also would have been d i f f i c u l t to account for the y i e l d i n g of the bolt during a compliance c a l i b r a t ion). Figure 40 : Stress intensity vs. crack length for d i f f e r e n t specimen geometries Stress-Intensity vs. Crock-lengh 22 -20-E x ample tor P -- 1000N L = 615mm W.B= 25.<,rrm 16 -U • 12 -59 2.5.2. Test set up Before the test, the samples were cleaned with acetone and alcohol. Subsequently, the specimens were wrapped with PTFE tape so that only the notch area was exposed to the environment. After t i g h t l y bolting the two pieces together (with SAE Grade 5 b o l t s ) , the bolt ends were also c a r e f u l l y wrapped with PTFE tape. Copper wires were screwed into each sample and connected to the same potentiostat used for the SSRT studies (see page 51). The test c e l l (see Fig.41, pg. 60) was a 600 ml Teflon beaker f i t t e d with a Teflon l i d . From the side, openings were cut into the l i d for the grips. Through the top, the following devices were inserted: Luggin c a p i l l a r y , temperature probe, platinum counter electrode, gas i n l e t and condenser outlet. A heating mantle was mounted at the bottom of the beaker and the temperature was adjusted with a temperature c o n t r o l l e r as described previously (sections 2.3.1 and 2.4.1). The Luggin c a p i l l a r y was positioned about 1mm from the crack t i p . To minimize contamination of the test s o l u t i o n with KC1 from the reference c e l l , a "bridge" of test solution was set up between the test c e l l and a beaker f i l l e d with cold (room temperature) test s o l u t i o n . A s a l t bridge with saturated KC1 was then connected from the beaker to the calomel reference c e l l , ( i . e . a double bridge assembly), as shown in Figure 41. 60 F i g u r e 41 : T e s t s e t u p f o r s t r e s s c o r r o s i o n e x p e r i m e n t s w i t h f r a c t u r e m e c h a n i c s s p e c i m e n G : G a s i n l e t T h : T h e r m o m e t e r 0 : G a s o u t l e t C : P t - c o u n t e r e l e c t r o d e E : W o r k i n g e l e c t r o d e ( S C C - s a m p l e s ) R : C a l o m e l r e f e r e n c e c e l l P : T e m p e r a t u r e c o n t r o l l e r G r : G r i p s H : H e a t e r L : L u g g i n c a p i l l a r y 61 The gas i n l e t was used for C O 2 in tests with carbonate solutions and for N 2 in a l l other t e s t s . However, N 2 purging was always stopped after the test was run for a few hours (otherwise gas bubbles tended to form at the t i p of the Luggin c a p i l l a r y ) , and the gas outlet was sealed. Testing Procedure: The sample was immersed into the hot test solution (95°C for a l l experiments). Then the desired load was applied with a ho r i z o n t a l l y mounted Hounsfield tensometer and, at the same time the poten t i a l was set to a fixed value with the potentiostat. The l i d and a l l connections were subsequently sealed with s i l i c o n rubber to avoid evaporation of the test solution. During the whole test the load was kept constant. This required very l i t t l e adjustment of the loading mechanism becouse of the slow cracking rates. The normal duration of one test was 500 hours (3 weeks). After the test, the samples were broken open in l i q u i d nitrogen. Only two experiments required interruption of the test because of a bolt f a i l u r e . Consequently the specimen configuration proved to be adequate for SCC t e s t i n g . The major problem was the development of gas bubbles in the Luggin c a p i l l a r y which could cause a loss of poten t i a l c o n t r o l . Insertion of a cotton thread into the c a p i l l a r y helped in some circumstances, except when the test solution dissolved i t during the tes t , eg. in the sodium hydroxide solutions. 62 For investigations of the fracture surface with the S E M , the samples were cleaned with i n h i b i t e d acid as described on page 51. The crack propagation increment was measured with a t r a v e l l i n g microscope at f i v e d i f f e r e n t locations on the fracture surface, (see Fig.42) Figure 42 : I l l u s t r a t i o n of crack propagation measurements Crack propagation measurements Notch Fatigue precrack Stress corrosion crack (Aa) Overload f a i l u r e ( l i q . N 2) 63 I I I . RESULTS 3.1. Anodic P o l a r i z a t i o n curves In the following Figures anodic p o l a r i z a t i o n curves of a l l three rotor steels are shown. A l l curves are drawn with two d i f f e r e n t scales for the p o t e n t i a l : the standard hydrogen electrode (Vg H E) and the standard calomel electrode ( V S £ E ) . The standard hydrogen electrode scale was obtained by adding 242mV to the standard calomel electrode scale. A l l experimental measurements were made with the calomel reference electrode. A d d i t i o n a l l y , results from the remelted and "doped" steel R are shown. 3.1.1. Carbonate solutions Tests in the d i l u t e carbonate solutions were mainly done with the "0" s t e e l . However, some tests were repeated with "A" and "R" ste e l s but no major difference was found. In the a c i d i c C02/H2O solutions no passivation phenomena was observed and general corrosion occured. The corrosion current density increased towards more noble p o t e n t i a l s . See Figure 43, pg. 64. A small addition of bicarbonate to the C0 2/H 20 solution raised the pH and the p o l a r i z a t i o n behavior changed dramatically. See F i g . 44, pg. 64. A wide active peak at 64 around -500mVgcE appeared, followed by a passive region above -200mVgc E The corrosion current density in t h i s passive region was about two orders of magnitude smaller than in the C O 2 / H 2 O system where no passivation was observed. In the very d i l u t e bicarbonate solutions there was a great tendency for crevice corrosion and with samples mounted in epoxy the passive behavior was overwhelmed by c r e v i c i n g e f f e c t s . 0.6-0.4-0.2 a ao-3 -02-I £ -0.4--0.6. -0.8--1.0 stee l " A " steel " R " steel "0" 10-3 XT2 X)'1 1 10 Current 1.0 as 06 04 LU 02 g ao c o 02 a " 04 06 08 103 «* Figure 44: P o l a r i z a t i o n curve for s t e e l " 0 " in N a H C 0 3 (3.5-10" 3M) fa Figure 45: P o l a r i z a t i s t e e l " 0 " in N a H C 0 3 (3.5-10 2M), fast and o n curves for no"slow scan 65 With increasing bicarbonate concentration, the tendency for crevice corrosion decreased and p o l a r i z a t i o n tests could be done with the rod shaped samples as well as with the f l a t mounted specimen. Additions of bicarbonate below one mole per l i t e r did not have a great influence on the passive region of the curves. However, i t may be stated that the corrosion current in the passive region increased with the increased bicarbonate concentration. (compare Fig.44-46, pg. 64,66) The active peak in the 3.5*10~3 molar solution was a single wide peak. At higher concentrations t h i s peak subdivided into two peaks. In the 3.5'10~2 molar solution t h i s was only observed at the slow scan rate, whereas in the 3.5*10~1 molar solution both peaks were found at fast and slow scan rates. Upon r a i s i n g the pH further in the 1M carbonate/bicarbonate solution, only one active peak was observed. But now the behavior in the passive region changed. At about 300mV S C E a current peak ("transpassive" peak) was observed with current densities nearly ten times higher than in the passive plateau. (See Fig.47, pg. 66) Also in the more d i l u t e solutions, sometimes small peaks were found in the passive region but not as pronounced as in the 1 molar carbonate/bicarbonate solution. 66 W ] 0.6 CU-02-tt 00-I-02J -<u-j -0.6--08--1.0-i fast, Imv/s slow, 0.3mV/s 10-3 »"2 10 « 2 03 I f f 1 1 W « 2 ioT Cufrtnt [*^,2] I-to as - 06 - 04 02 _S .2 00 I o 02 ^ QA - 06 - as 10* Figure 45: Po l a r i z a t i o n curves for steel "0" in NaHC03 (3.5-10 2M), fast and slow scan Figure 46: Po l a r i z a t i o n curve for steel "0" in NaHC03 (3.5-10"1M) Figure 47: Po l a r i z a t i o n curves for steels "A", "R" and "O" in 1M NaHC03 + 1M Na 2C0 3 67 3.1.2. Sodium Hydroxide solutions In the strong a l k a l i n e sodium hydroxide solution (3.5 molar, pH 13), active/passive behavior of a l l three steels was observed. (See F i g . 48, pg. 68) The active peak consisted a c t u a l l y of two peaks but the i r separation was not well defined in a l l t e s t s . The current density at the peak was more than one order of magnitude smaller than in the carbonate solutions. In the passive region at about -150mVgrjg, again another peak was detected with a corrosion current density comparable to the one in the carbonate so l u t i o n . This transpasssive peak was more pronounced in the s t e e l "A" and "R", whereas in "0" a higher current density in the passive "plateau" diminished the peak appearance. In t h i s solution the position of the oxidation curve with respect to the reduction curve caused more than one inte r s e c t i o n point, producing a zero applied current (compare Fig.35, page 46). Therefore, another "cathodic peak" i s seen in some p o l a r i z a t i o n curves at about -800mV S C E. However, t h i s peak was not of interest for the present corrosion study. With a NaOH concentration of 0.35 mole/1, the po l a r i z a t i o n curves were about i d e n t i c a l to the 3.5 molar solutions. However, only a single active peak appeared. 68 M 06-04 -(if goo- Iii 1 -ou?-2 -0.4 • -0.6--0.8 --1.0. Steel "A" Sleet "R" steel "0" W3 IO2 10 W2 Current te. 1.0 hOJ 06 Q4 02 h 00 02 04 06 08 10* tn o I o 0. F i g u r e 4 8 : P o l a r i z a t i o n c u r v e s f o r s t e e l s "A", "R" a n d " 0 " i n 3.5M NaOH I n F i g u r e 49 a p o l a r i z a t i o n t e s t o f t h e s a m e s a m p l e w a s r u n t w i c e . I t c a n b e s e e n t h a t t h e same r e s u l t i s o b t a i n e d e x c e p t f o r t h e p e a k i n t h e p a s s i v e r e g i o n w h e r e a d e c r e a s e i n c o r r o s i o n c u r r e n t c o u l d b e s e e n . -06-0.4-02-~\— first scan 1-02-£ -<u--06-fl i ! 1 JJ —second scan -0.8- --1.0--ur3 Ho-2 lb'1 1 10 w2 Current fyrf] 06 02 J g ao J o 02 °-Q4 t)3 104 F i g u r e 4 9 : P o l a r i z a t i o n c u r v e s f o r s t e e l "R" i n 3.5M NaOH, f i r s t a n d s e c o n d s c a n 1 3.1.3. Steel "R° doped with sulphides Steel "R" was vacuum remelted and then doped with either CaS or MnS. Pol a r i z a t i o n tests were done in the as-cast, hot r o l l e d and heat treated condition. Tests were only performed in the 3.5 molar NaOH solution at 95°C. The results are shown in (Fig.50 and 51). For comparison "R"-steel "as received" (not remelted) i s also included. From the microstructure investigation i t i s known that CaS inclusions were also present in "R" but to a lesser degree than in the CaS doped s t e e l . 70 Figure 51: Polariz a t i o n curves for steel "R" and st e e l "R" doped with MnS in 3 . 5 M NaOH W 3 KT2 KT1 Current % 2 J No major difference between the doped s t e e l and the regular R-steel could be seen. The shape of the anodic peak seemed to consist of d i f f e r e n t peaks in both doped s t e e l s . The maximum corrosion current density at the peak was about a factor of five higher than in the regular R-steel. Also the current density in the passive region seemed to be higher in the doped stee l s and, therefore, no cathodic peak appears in t h i s region. The e a r l i e r observed peak in the passive region was only observed in the heat treated condition and appears to be uninfluenced by the inc l u s i o n s . P r a c t i c a l l y no difference could be seen between the two doped s t e e l s in the entir e potential range investigated. To f u l l y characterize the influence of inclusions, many more experiments are necessary. Different solutions with lower pH values have to be tested. 71 3.2. Slow s t r a i n rate t e n s i l e tests In Table II (pg.72) results of a l l SSRT studies are summarized. It can be seen that nearly a l l tests with the carbonate solutions gave much deeper cracks than in the sodium hydroxide solution. This was observed with a l l three s t e e l s . The same behavior i s r e f l e c t e d in the reduction in area, where lower values were obtained in the carbonate solutions (a small reduction in area i s equivalent to l i t t l e necking). Steel "A" and "R" seemed r e l a t i v e l y i n s e n s i t i v e to cracking in the sodium hydroxide solutions. Figure 52 page 73 i l l u s t r a t e s the e f f e c t of p o t e n t i a l on the reduction in area. The effect of small p o t e n t i a l changes on the cracking behavior of SSRT samples can be seen in Figures 53 and 54 (pg. 74,75). However, to f i n d a clear dependence between the reduction in area and the applied p o t e n t i a l many more tests would be necessary. 72 Table II : Results of slow s t r a i n rate t e n s i l e t e s t s Steel Envi ronment Potential m VSCE m VSCE Red in area % Max. crack length >um Crack v e l o c i t y m/s "0" C0 2+H 20 n B B B 0 -500 -600 -685 -1000 32.7 29.1 34.4 23.7 22.0 68 1 1 3 45 29 9.4-10~^° 1 .6-10"9 6.2- 10~10 4.0•10" 1 0 "A" C0 2+H 20 n -600 free 71 .8 73.5 68 <5 9 . 4 * 10" 1 0 <6.9•10~11 "0" 3.5*10"3M NaHC03 n B B n -100 -200 -370 -500 -1000 15.3 26. 1 32.6 40.8 26.6 250 159 1 59 3.5*10"^ 2.2- 10" 9 2.2-10"9 "A" 3.5*l0" 3NaHCO 3 -500 66.7 1 14 1 .6*10"9 "0" 1MNa2C03+1MNaHC03 B n B n n n -600 -680 -750 -750 40.8 39.6 38.0 38. 1 45 159 136 6.2- 10~ 1 0 2.2-10"9 1 .9-10~9 "A" 1MNa2C03+1MNaHC03 -750 67.4 1 14 1 .6-10 - 9 "R" 1MNa2C03+1MNaHC03 -750 68.7 182 2.5-10"9 "0" 3.5M NaOH n n B -150 -400 -1000 -1000 48.8 44.9 39.0 41.6 34 23 10 23 4.7-10~10 3.2-10 1 ° 1.4 * 10" 1 ° 3.2 * 10~ 1 0 "A" 3.5M NaOH B -880 -980 72.4 73.5 <5 18 6.9-lO - 1^ 2.5•10~ 1 0 "R" 3.5M NaOH B B -600 -880 -950 68.4 72.3 72.7 <7 9.7-10"1 1 "0" 0.35M NaOH -850 43.9 — "0" "A" "R" O i l (95°C) O i l (95°C) O i l (95°C) 48.7 75.8 72.5 --73 M 10 H 1<T Steel -"0" 35 M NaOH F i g u r e 52: V. Red inar*a I n f l u e n c e of p o t e n t i a l on the r e d u c t i o n i n area d u r i n g SSRT-tests "SCE 10' 10 4 io-H Steel-"0" 33-10'3M NaHC03 T | 7 ' t I A. • ! 1 / \ 1 1 / \ 1 , / \ \ / " \ 1 7 1/ OX 02 00 -02 -0.4 -06 -08 U> V. Red. in area 50 40 30 20 Steel -"R" w 35 M NaOH 10* • 10 • * ? ? I I l 1 ' 1 - ^ A s ft IO"1-• i > — i — V. Red. in area 90 70 60 50 04 02 00 -02 -04 -Q6 08 -10 ySCE 74 F i g u r e 53 : R e d u c t i o n i n a r e a a n d s u r f a c e c r a c k s o f s t e e l - " 0 " a t d i f f e r e n t p o t e n t i a l s i n t h e C 0 2 / H 2 0 s y s t e m , S S R T - t e s t - 5 0 0 m V S C E Steel-"0" M C O ^ O 102> 10 • T T t I 1 1 T • 1 1 • \ 1 1 10"1-w ..'WtutC ' i* > v - v m 04 02 00 02 -OX -Q6 -08 -10 V SCE •600 mV S C E 7. Red in a r e a 50 40 30 20 10 - 6 8 5 mV S C E 75 Figure 54 : Reduction in area and surface cracks of steel-"0" at d i f f e r e n t potentials in the 1M NaHC03 + Na 2C0 3 system, SSRT-test -600 mV SCE M K>2 10 10' 1. Steel -"0" % Red IM NaHC03/ / N Q 2 C O 3 in area •50 t •» 1 1 / X •40 1 / \ 1 l/ \ •30 1/ 1 •20 1 / 10 02 00 — 1 1 -1 — 1 •02 -OA -Q6 -0.8 —i ' •TJO VSCE -680 mV S C E -750 mV SCE 76 The environment and electrochemical p o t e n t i a l selected for the fracture mechanics tests were based on micrographs of sectioned SSRT samples. From these, conclusions could be drawn as to whether the corrosive attack during the SSRT experiment was of a general corrosion and p i t t i n g nature, or i f sharp, penetrating cracks were formed. The l a t t e r were assumed to be a sign for s e n s i t i v i t y to SCC. Different types of SSRT-cracks are i l l u s t r a t e d in Figure 55 (pg. 77). From the longest secondary crack length, an average crack v e l o c i t y could be estimated by d i v i d i n g the crack length by the time of p l a s t i c deformation during the t e s t . In the 1 molar carbonate/bicarbonate solution, a crack growth rate of about 2*10~9m/s was obtained for a l l three s t e e l s . However, in the 3.5 molar NaOH solution, steel-"0" was cracking at about 3.5 * 10~ 1°m/s, steel-"A" gave a maximum value of about 2"l0 _ 1°m/s and steel-"R" seemed to show a cracking-rate below I0~ 1°m/s. In the Figures 56-58 on page 78, cracking on the surface of the three steels in 3.5M NaOH i s compared. It was d i f f i c u l t to draw conclusions about the cracking path. Steel-"A" appeared to show some indications of intergranular cracking. 77 Figure 55 : Surface cracks in SSRT samples after tests in 3.5M NaOH (steel-"R"), C0 2/H 20 (steel-"0") and 1M NaHC03 + IM Na 2C0 3 (steel-"A") Steel "R"f 3.5M NaOH, -950 mV S C E Shallow surface cracks Mag.: 400x, etched Steel "0", C0 2 / H 20, -600 mV S C E Strong surface attack Mag.: 400x, etched Steel "A", 1M NaHC03 + 1M Na 2C0 3, -750 mV S C E Deep, penetrating cracks Mag.: 400x, etched 78 F i g u r e 56: SEM-picture, s u r f a c e c r a c k s of s t e e l "A" a f t e r t e s t i n g i n 3.5M NaOH at -980 mV S C E, SSRT-test The s u r f a c e of the samples was cleaned with i n h i b i t e d a c i d Mag.: 1,000x F i g u r e 57: SEM-picture, s u r f a c e c r a c k s of s t e e l "R" a f t e r t e s t i n g i n 3.5M NaOH at -950 mV S C E, SSRT-test Mag.: 1,000x F i g u r e 58: SEM-picture, s u r f a c e c r a c k s of s t e e l "0" a f t e r t e s t i n g i n 3.5M NaOH at -1000 mV S C E, SSRT-test Mag.: 1000X 79 In F i g u r e s 59 and 60, the f r a c t u r e s u r f a c e of s t e e l "A" and "0" i n s i d e the cr a c k s i s compared. No major d i f f e r e n c e can be seen, which may e x p l a i n the s i m i l a r c r a c k v e l o c i t i e s observed f o r both s t e e l s i n the ca r b o n a t e / b i c a r b o n a t e s o l u t i o n . F i g u r e 59: SEM-picture, f r a c t u r e s u r f a c e i n s i d e a c r a c k , s t e e l "A" t e s t e d i n 1M NaHC0 3 + 1M N a 2 C 0 3 at -750 mV S C E, SSRT-test At the top: s u r f a c e of the SSRT-specimen Mag.: 800x F i g u r e 60: SEM-picture, f r a c t u r e s u r f a c e i n s i d e a c r a c k , s t e e l "0" t e s t e d i n 1M NaHC0 3 -750 mV S C E, SSRT-test 1M N a 2 C 0 3 at At the top: s u r f a c e of the SSRT-spec imen Mag.: 800x 80 3.3. Fracture mechanics tests In Table I I I , page 84 the results of a l l fracture mechanics SCC tests are summarized. The numbering of each specimen has no meaning in respects to the testing conditions, i t was done during specimen f a b r i c a t i o n . Most tests were run for 500 hours or more. However, some tests needed to be terminated e a r l i e r because of instrumentation problems, ( i . e . unstable potential control) The stress i n t e n s i t y i s reported at the onset of cracking Kj(a) and a f t e r a crack growth increment of Aa at the end of a test Kj(a+Aa). The increase in crack length, Aa, i s an average value obtained from f i v e measurements (compare Fig.42, page62). The crack growth rate does not include an unknown incubation time and, therefore, the actual crack propagation rate may be higher (especially for short testing times). For s t e e l "A" and "R" samples were only tested in the "CC" and "CL" d i r e c t i o n s , becouse the thickness dimension of the discs was not s u f f i c i e n t for "LC" samples. For SCC tests, these specimens were not labeled as "CC" or "CL" since both mechanical tests and metallographic investigations did not reveal any difference between these two d i r e c t i o n s . Steel-"0" was tested in "CC" , "CL" and "LC" d i r e c t i o n s . Because of the d i r e c t i o n a l property dependence 81 in t h i s material, a l l samples were l a b e l l e d according to t h e i r d i r e c t i o n s . (See Table I I I , pg. 84) 3.3.1. Carbonate solutions The 1.5%CrMoV steel-"0" showed no cracking in the ac i d i c C0 2 / H 20 solution and none in the near neutral C0 2 + NaHC03 (3.5-10 - 3 molar) solution. In the a l k a l i n e 1 molar carbonate/bicarbonate solution, only three samples showed cracking. However, the specimen No."0"/2* may not be compared with the other samples since the solution in t h i s case was strongly contaminated with chloride ions and the pot e n t i a l during the test was not stable due to problems in the e l e c t r i c a l c i r c u i t . The reason why the other two samples cracked i s not clear since a l l the other tests with lower and higher stress i n t e n s i t i e s showed no cracking. The 3.5%NiCrMoV-steel-"A" cracked in the 1 molar carbonate/bicarbonate solutions at a stress i n t e n s i t y of about 50MNm~3/2. No cracking could be observed below 30MNm"3/2 . Steel-"R" was not tested in carbonate solutions. In the C0 2/H 20 solution, both steels "A" and "R" showed strong corrosion at the crack t i p . However, t h i s may not be c a l l e d stress corrosion cracking. It can be assumed that t h i s form of attack blunts the crack t i p and therefore, 82 reduces the tendency for further crack propagation. An i l l u s t r a t i o n of the "rounded" crack t i p i s given in F i g . 61. Figure 61: SEM-picture, crack t i p of steel "A" after testing in C0 2 / H 20 at the free corr. pot., SCC-test At the top: overload fracture in l i q . Nitrogen At the bottom: fatigue precrack At the centre: rounded SCC crack Mag.: 900x 3.3.2. Sodium hydroxide solutions Figure 62, pg.85 shows a v-K plot of a l l three ste e l s in 3.5 molar NaOH. The potential was set corresponding to the active peak at -870 to -lOOOmVgcE* (Compare p o l a r i z a t i o n curves F i g . 48, pg. 68) For the 1%CrMoV steel-"0", the stress i n t e n s i t y was varied between about 25 and 45 MNm~3//2. The crack growth rate seemed to decrease s l i g h t l y with decreasing stress i n t e n s i t y . However, even at a stress i n t e n s i t y below 30 MNm-3'2 a crack growth rate around 4*10 - 1 0 m/s was measured. The K I S c c value (stress intensity below which no measurable cracking i s found) must l i e below about 83 25MNrrr 3 / /2. However, e x p e r i m e n t s a t such low K - v a l u e s were not p e r f o r m e d . The i n f l u e n c e of c r a c k p l a n e d i r e c t i o n seemed t o be r a t h e r m i n i m a l . The LC d i r e c t i o n may have g i v e n a s l i g h t l y s l o w e r c r a c k r a t e than the " C L " - d i r e c t i o n a t t h e same s t r e s s i n t e n s i t y , but t h e r e was no c l e a r d i f f e r e n c e i n c r a c k i n g v e l o c i t y between "CC" and "CL" d i r e c t i o n s . The 3.5%NiCrMoV s t e e l "A" was t e s t e d between a p p r o x i m a t e l y 40 and 55 MNm - 3/ 2. I t was n o t i c e d t h a t the c r a c k growth r a t e d e c r e a s e d r a p i d l y around 40 MNm~3/2. T h e r e f o r e , a K i s c c v a l u e above about 35 MNm~3/2 c o u l d be e s t i m a t e d . The c r a c k growth r a t e above 40 MNm~3//2 was around I0~ 1°m/s and seemed r e l a t i v e l y independent of the s t r e s s i n t e n s i t y , ( t y p i c a l s t a g e I I b e h a v i o r i n a v-K p l o t ) . Not many t e s t s were run w i t h the 3.5%NiCrMoV s t e e l "R". No c r a c k i n g was ob s e r v e d below 40 MNm~3//2. At about 50MNm - 3/ 2, however, a c r a c k growth r a t e s i m i l a r t o the " 0 " - s t e e l was measured which was h i g h e r than f o r s t e e l - " A " . 84 Table III : Results of fracture mechanics (SCC) tests S l e e l /No Fracture p] » n i T e E t -solut ion Pol enl lal • v s c r T u t -durat i on hrs K <•) K 1 ( » - t»Bl («B> CracV growth r»t r ml t "0"/2" /JO /)] /)3 /14 nb nb m /16 ct: CC CC CL CI CL CL CL CL IM Carb/Blcarb -750 It 1000 500 500 500 500 720 720 720 720 40.6 36.4 25.3 34 .9 32.3 41 .0 38.5 39.6 4] .4 4r..0 35.4 32.7 900 110 100 2.5*10- 1 0 6.1x10"'' 5.6x10 "A'749 /SO /65 Ibb IM Carb/Bicarb -7 50 290 290 500 500 26.1 30.1 50.1 48.1 52.6 50. 6 330 360 ;,«>:;; 2.0x10 "0"IU 121 127 CC LC IX C0 2 + H 20 free corr 150 4 30 430 29.3 37.0 36.0 _ -_ "A'763 /64 C0 2 + H 20 free corr 340 340 65.3 51 .4 (-) (-) (30) (30) (-) (-) "R'757 /60 C0 2 + H 20 free corr 360 360 55.4 50.7 (-) (30) (20) (-) (-) "0'719 /20 CL CL C02+KaHC03 -4 00 500 500 41 .4 32.7 ---"0'723 /24 /25 /26 /34 /36 /37 /38 LC LC CL LC CC CL CL CC 3.5M NaOH -1000 360 430 360 430 500 500 500 500 41 .9 38.6 41 .5 37.0 26.2 35.2 28.5 34.5 44 . 2 43.7 46.8 4] .3 29.0 44 .2 30. 9 38.4 360 980 990 920 84 7 1460 640 730 6.3x10. 7.6x10. ^  5.9x10.]° i - 7 x l ° - 0 3.6x10. ]° 4.1x10 "A'745 /46 /47 /48 3.5M NaOH -1000 460 4 60 500 500 54.8 44 .9 44.2 41 .1 56.0 45.7 44 .7 42.1 200 150 100 190 1 .2x10"'° 9.3x10" 5.6x10" 1 .0x1 0" "R'753 /54 /55 /56 3.5M NaOH -1000 -950 500 500 500 500 40.0 39.0 47 .4 48.8 52.2 54 .3 660 690 3.7*1 cr1 £ 3.8x10 "0*761 /62 CC CL 3.5>i NaOH -650 460 460 36.5 37.6 38.2 50 3.0xl0" 1 3 "0*729 /30 LC LC 3.5K NaOH -400 770 770 39.0 49.2 39.9 51 .7 170 360 6.0x10"" 1.3x10 "0"/31 /32 CC CC 3.5M NaOH -150 500 500 36.5 36.9 37 .5 37 .9 210 210 1.2X101]° 1 .2x10 "0"/33 /35 CC CC 3.5K NaOH * 0 I 500 500 31.7 39.9 40.1 40.7 60 130 3.3xlO~" 7.2x10" "A'743 /44 3.5M NaOH -870 500 500 38.3 40.6 38.7 40.8 50 20 2.9xl0 :" 8.9x10" "A"/51 752 3.5M NaOH -150 460 460 56.1 48.6 56.5 49.0 40 40 2.4x1 o"" 2.4x10 "0'739 Ml CC CC 0.35M NaOH -1000 500 500 33.8 35.4 _ - -"0*727 /28 LC LC 0.35M NaOH -870 II 770 770 36.9 40.9 -- -85 Figure 62 : Stress i n t e n s i t y vs. crack v e l o c i t y for steels "A", "R" and "0". SCC-tests in 3.5M NaOH at the active peak pot e n t i a l (around -1000 mVcrp) v["te] 10-9-o •5 10' > o 10 10-11-3.5M NaOH Steel -0" Steel - A" Steel - R" - r — 10 CL CL cc * LC 1t. 20 30 'iO Stress - Intensity — i — 50 K, [MNm"^ ] 60 86 For steel-"0" (and "A"), the influence of test p o t e n t i a l on the cracking rate is i l l u s t r a t e d in Fig.63, pg. 87. The p o t e n t i a l was varied between 0 and -1OOOmVg^g. The stress i n t e n s i t y was chosen around 35 MNm-3/2 f o r s t e e l "0" and around 46 MNm~3/2 for steel "A". Steel-"0" seemed to show maximum crack growth rates at the a c t i v e peak at -1u00mVgcE and at the peak in the passive region at -150mVgcE. The difference between the cracking rates at the two peaks was less than one order of magnitude. Lower cracking rates were found between the two peaks and above -l50mV S C E. For comparison most tests were done with "CC"-spec imens. For steel-"A", not many measurements were done. However, the cracking rate at -lOOOmV S C E seemed to be only about 4 times higher than at ~l50mVgrj E > Steel "0" was also tested in 0.35M NaOH but no cracking was found at either -lOOOmVgCE or -870mV S C E. Both tests were run at stress i n t e n s i t i e s where cracking occurred in the more concentrated s o l u t i o n . 87 Figure 63 : Potential vs. crack v e l o c i t y for steels "A" and "0" in 3.5M NaOH, SCC-test 200 -200 -400 -600' -800 -1000 Potential [™ VSCE Steel-"0" • -1000mVS C E • -650mV A -400nW • -150mV • OmV (Kj«32-40MNm'&) Steel- "A" O -1000mVSCE • -150mV (Kr44-48MNm"^2) 88 3.3.3. Fractoqraphy Steel-"A" (3.5%NiCrMoV, super clean) In both the 3.5M NaOH and the 1M carbonate/bicarbonate solutions, intergranular cracking was observed at a l l potentials investigated. In the sodium hydroxide solution, however, the appearance of the intergranular fracture surface varied with the testing p o t e n t i a l . At -lOOOmVgcg (active peak) the individual grains on the fracture surface were hardly attacked by the environment whereas at -l50mVg£ E (transpassive peak) strong etching (or p i t t i n g ) of the whole fracture surface was observed. (Compare Figures 64 and 65 on page 91). In the carbonate/bicarbonate solution, a strong corrosive attack of the grains occurred at the crack mouth whereas smoother grain surfaces were found at the crack t i p (compare Figures 66 and 67, pg. 92). Even when the crack path in a l l solutions was intergranular, some tests at higher stress i n t e n s i t i e s (above 55 MNm~3/2) showed mixed cracking (intergranular/transgranular). Steel-"R" (3.5%NiCrMoV) Steel-"R n was tested only in 3.5M NaOH at the active peak po t e n t i a l and gave a fracture appearance similar to steel-"A". 89 Steel-"0" (UCrMoV) The determination of the crack path in a l l tests was rather d i f f i c u l t . This could be due to the mixed f e r r i t i c / b a i n i t i c microstructure. Most probably both intergranular and transgranular cracking occurred. (Compare Figures 68 and 69, pg. 93). In the 3.5M NaOH solution, once again, a strong influence of the potential on the appearance of the fracture surface was observed. At the transpassive peak p o t e n t i a l of -l50mV S C E, a much stronger attack on the grain surface occurred than at the active peak p o t e n t i a l of -1u00mV S C E. (Compare Figures 71, 72 and 73 , pg. 94). At the test p o t e n t i a l s between the two peaks, a stronger corrosion attack was observed compared to the tests at -l000mVgc E > As mentioned e a r l i e r , Steel-"0" cracked only in 3.5M NaOH with three cracking exceptions in the carbonate/bicarbonate s o l u t i o n . Figure 74 on page 95 shows the fracture surface at the crack t i p in the carbonate/bicarbonate solution, which looks very s i m i l a r to the fracture surface in the NaOH solution at -l000mV S C E (Compare Fig.73, pg.94). From the rounded crack t i p , i t may be concluded that the crack propagation was mainly a d i s s o l u t i o n process. 90 E a r l i e r i t was stated that the orientation of the crack plane did not have a big influence on the propagation rate. Figures 75 and 76 on page 95 show two cases where MnS inclusions were laying in the crack plane ("CL"-direction). Both pictures show the t i p of the stress corrosion crack. In Figures 75 (3.5M NaOH, -l000mV S C E) the crack front seems not to be influenced by the presence of the i n c l u s i o n . From Figure 76 (1M carbonate/bicarbonate,-750mVg£ E), on the other hand, one could conclude that the crack propagated further in the v i c i n i t y of the i n c l u s i o n . However, th i s " a d d i t i o n a l " cracking was only observed very close to the inclusion and may not have markedly influenced the o v e r a l l cracking behavior. In both examples no change in fractography adjacent to the inclusions was observed. 91 Figure 64: SEM-picture, crack t i p with intergranular fracture surface, steel "A" tested in 3.5M NaOH at -1000 mV S C E, SCC-test At the top: overload fracture in l i q . Nitrogen Mag.: 1,200x Figure 65: SEM-picture, strongly etched intergranular fracture surface, steel "A" tested in 3.5M NaOH at -150 mV S C E, SCC-test Mag.: 1,200x 92 Figure 66: SEM-picture, p i t t e d intergranular fracture surface at the crack mouth, s t e e l "A" tested in 1M NaHC03 + IM Na 2C0 3 at -750 mVc.rR, SCC-test Mag.: 1,700x Figure 67: SEM-picture, intergranular fracture surface, at the crack t i p , s t e e l "A" tested in IM NaHC03 + 1M Na 2C0 3 at -750 mV S C E, SCC-test Mag.: 1,700x 93 F i g u r e 68: SEM-picture, s t r o n g c o r r o s i v e a t t a c k of the 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 , s t e e l "0" t e s t e d i n 3.5M NaOH a t -1000 mV S C E, SCC-test Mag.: 800x F i g u r e 69: SEM-picture, t r a n s g r a n u l a r f a c t u r e appearance, s t e e l "0" t e s t e d i n 3.5M NaOH at -1000 mVe r p, SCC-test Mag.: 800x F i g u r e 70: Secondary c r a c k s , s t e e l "0" t e s t e d i n 3.5M NaOH at -1000 mV S C E, SCC-test Mag.: 200x, e t c h e d 94 F i g u r e 71: SEM-picture, s t r o n g l y a t t a c k e d f r a c t u r e s u r f a c e , s t e e l "0" t e s t e d i n 3.5M NaOH a t -150 mV S C E, SCC-test Mag.: 900x F i g u r e 72: SEM-picture, s t r o n g l y a t t a c k e d f r a c t u r e s u r f a c e , s t e e l "0" te s t e d i n 3.5M NaOH at -150 mV S C E, SCC-test Mag.: 1,000x F i g u r e 73: SEM-picture, f r a c t u r e s u r f a c e , s t e e l "0" t e s t e d i n 3.5M NaOH at -1000 mV S C E, SCC-test Mag.: 900x Figure 74: SEM-picture, rounded crack t i p (at the top), with strongly attacked fracture surface, steel "O" tested in 1M NaHC03 H 1M Na 2C0 3 at -750 mV S C E, SCC-test Mag.: 1,000x Figure 75: SEM-picture, crack t i p with MnS inclusion (dissolved away), steel "O" tested in 3.5M NaOH at -1000 mV S C E, SCC-test Crack front i s not influenced by the inclu s i o n MnS inclusion Mag.: 500x Figure 76: SEM-picture, crack t i p with MnS inclusion (dissolved away), steel "0" tested in 1M NaHC03 + 1M Na 2C0 3 at -750 mV S C E, SCC-test Crack propagated further in the v i c i n i t y of the i n c l u s i o n MnS inclusions Mag.: 200x 3.3.4. Crack path Strong crack branching was observed in a l l SCC-tests. Normally, two main cracks propagated from the t i p of the fatigue crack at an angle of about 60° to the crack plane of the fatigue crack and then branched further into smaller cracks. This cracking e f f e c t i s described by Brown 5 9 and was explained with the shape of the p l a s t i c zone at the crack t i p . He observed that in bend-bar stress-corrosion-tests, the crack tends to follow the e l a s t i c / p l a s t i c interface formed at the crack t i p . Green and Mundy^1 did a t h e o r e t i c a l and experimental analysis of the p l a s t i c zone in the notched bend specimen and found that the e l a s t i c / p l a s t i c interface ahead of the crack t i p has an e l l i p t i c a l shape. The observed crack branching angle during our SCC tests were consistent with these observations and explanations. An example of crack branching observed in s t e e l "A" i s given in Figure 77 below. Figure 77: Crack branching, s t e e l "A" tested in 3.5M NaOH at -1000 mV S C E, SCC-test At the top: v e r t i c a l fatigue precrack SCC cracks At the bottom r i g h t : overload overload fracture 97 3.3.5. Dynamic SCC-test A slow loading rate fracture mechanics test was employed to explain the d i f f e r e n t findings of SCC experiments and SSRT experiments with steel-"0". In the IM carbonate / bicarbonate solution, very severe cracking was found with the SSRT-test, whereas the conventional s t a t i c a l l y loaded fracture mechanics experiments gave no cracking beside a few exceptions. The dynamic test set up was analogous to a regular SCC tes t , but instead of applying a constant load, the sample was pulled slowly to f a i l u r e . The crosshead speed was the same as during the SSRT-tests which corresponded to a stress i n t e n s i t y increase rate of about 1.5 MNrn-3/2 per hour. The test was started at zero load and was run for about 48 hours. Therefore the stress intensity at the end of the test would be much higher than the KJQ value found by tes t i n g t h i s steel in a i r . The reasons for thi s are crack blunting and crack branching during the test. In the SSRT experiments (1M carbonate / bicarbonate, -750mV S C E) with a duration of about 24 hours, a crack length of about 0.16mm was obtained. From the fracture surface of the dynamic SCC-test a maximum crack length of about 0.05mm could be observed and the fracture surface was not similar to the fracture surface obtained with the SSRT-test. (Compare F i g . 78, 79 below and SSRT Fig.60, pg. 79). 98 In the dynamic SCC test, a strong corrosive attack could be observed and i t i s not very clear i f the crack advanced by stress corrosion cracking or by some di s s o l u t i o n mechanism. The same experiment was run in a CO2/H2O environment at the free corrosion p o t e n t i a l . No cracking or di s s o l u t i o n at the crack t i p could be found at a l l . Figure 78: SEM-picture, crack t i p a f t e r dynamic SCC-test, s t e e l "0" in 1M NaHC03 + 1M Na 2C0 3 at -750 mV S C E At the top: overload fracture Mag.: 1,000x Figure 79: SEM-picture, crack t i p a f t e r dynamic SCC-test, s t e e l "0" in 1M NaHC03 + 1M Na 2C0 3 at -750 mV S C E At the top: overload fracture Mag.: 3,200x 99 IV. DISCUSSION 4.1. Interpretation of p o l a r i z a t i o n curves 4.1.1. Carbonate solution The E-pH diagram 6 2 of iron in water at 100°C (Fig. 80), indicates, that p a s s i v i t y i s d i f f i c u l t to achieve at pH <4.5, except at higher ox i d i z i n g potentials. With t h i s , passive behavior was not observed in the a c i d i c CO2/H2O solution Figure 80 : E-pH diagram of iron in water at 100°C 0 « a 12 pH 100 By adding NaHCC*3 to the solution the pH i s raised. Oxide films can now form at much lower potentials which means that passivation behavior is possible. The occurrence of d i f f e r e n t active peaks (compare Figures 44-47, pg.64,66) suggests that several reactions occured place and the composition of the passive fi l m changed as the po t e n t i a l was raised from negative (active)to more p o s i t i v e (noble) potentials. Several authors have studied the f i l m formation on mild steel in carbonate/bicarbonate s o l u t i o n s 6 3 - 6 7 . The following possible passivation processes were suggested and i t is assumed that si m i l a r reactions could have taken place on the rotor s t e e l . At low concentrations of carbonate ions (below 10 - 3 mol/1), the p r i n c i p a l oxide films on the metal surface are Fe30£ at lower potentials and Fe 203 at higher po t e n t i a l s . At carbonate concentrations above about 1 mol/1, FeC03 and Fe 203 are the main oxidation products, where the l a t t e r i s again formed at higher p o t e n t i a l s . At concentrations in between, a mixture of Fe-^O^ and FeC03 are found at lower potentials and Fe 203 at higher p o t e n t i a l s . 101 Thomas et a l b proposed sequence takes place in potential i s raised from the noble d i r e c t i o n : that the following reaction bicarbonate solutions as the free corrosion potential in the At potentials just below the f i r s t active peak the following reaction (7) dominates which i s independent of bicarbonate concentration: Fe + 2H 20 Fe(OH) 2 + 2H + + 2e" ...(7) By r a i s i n g the poten t i a l further, FeCC>3 forms either by attack of the Fe(OH) 2 (8)or by di r e c t reaction with Fe (9): Fe(OH) 2 + HC0 3" FeC0 3 + H 20 + OH" ...(8) Fe + HC03~ FeC0 3 + H + + 2e" ...(9) At the f i r s t active peak, passivation occurs due to a p a r t i a l coverage of the metal surface with FeC0 3 and p a r t i a l coverage with F e 3 0 4 following the reaction (10). 3Fe(OH) 2 — F e 3 0 4 + 2H20 + 2H + + 2e . . . O O ) (Where the r a t i o of F e 3 0 4 to FeC0 3 i s controlled by the concentration of bicarbonate) 102 The second peak i s caused by the oxidation of the FeC0 3 (11) and F e 3 0 4 (12)to an Fe(III) f i l m namely F e 2 0 3 : 2FeC0 3 + 3H20 F e 2 0 3 + 2C0 3 2" + 6H + + 2e" ..(11) 2Fe 30 4 + H 20 — — 3Fe 20 3 + 2H + + 2e~ ...(12) Figure 81: Discussion of the p o l a r i z a t i o n curve in terms of chemical reactions in the Bicarbonate+Carbonate solutions / Fe 20 3 > Fe3Q;,FeC03-Fe203 / FeC03 Fe(0H)2-Fe3Q4 Fe-Fe(0H)j Fe.Fe(OH)2-FeC03 Current-density The occurrence of only the f i r s t active peak at low concentration of bicarbonate was explained by Thomas0,3 as being due to the very slow oxidation of ferrous carbonate at the low pH of these solutions. This could also explain why in the 3.5*10~2 molar bicarbonate solution the second peak was observed only with the slower scan rate. 103 The appearance of only the f i r s t active peak in the 1 molar carbonate/bicarbonate solution i s explained^ 4 by the i n h i b i t i o n of f i l m breakdown at the t r a n s i t i o n from a protective FeCC>3/Fe3C>4 f i l m to the protective F e 2 0 3 f i l m . The transpassive peak was consistent with the d i s s o l u t i o n of Cr. The aqueous thermodynamics of the Cr/H 20 system*^ shows that the chromium oxides dissolve at higher potentials. In Figures 43 and 47, pg. 64.66, a l l three steels are compared in the CO2/H2O solution and in the carbonate/bicarbonate solution. The reason why s t e e l "0" showed the lowest current density at the transpassive peak may be explained by i t s lower chromium content and the lack of a heat treatment which can cause chromium enrichment of the metal carbides. The higher current density in the passive "plateau" of s t e e l - " 0 " could be due to a less protective f i l m because of smaller chromium and n i c k e l contents compared to the "A" and "R" s t e e l s . At the active peak in the carbonate/bicarbonate solution s t e e l , "A" and R" showed p r a c t i c a l l y no difference whereas in the passive region s t e e l "R" gave s l i g h t l y higher current d e n s i t i e s . In the C 0 2/H 2 0 solution steel "R" and "A" gave much higher corrosion currents than steel "0", and again s t e e l "R" gave higher current densities than "A". The reasons for th i s could be the difference in microstructure, heat treatment (segregation phenomena), a l l o y composition or the 104 inclusion composition. More experiments are, therefore, necessary to c l e a r l y explain these observations. 4.1.2. Sodium hydroxide solution Agrawal et al^° investigated the p o l a r i z a t i o n behavior of Fe-Ni-Cr al l o y s in sodium hydroxide solutions and Crowe et a l 6 9 studied SCC of stainless s t e e l in caustic solutions. Also the behavior of iron and mild s t e e l in sodium hydroxide solution was studied by several r e s e a r c h e r s 7 ^ - 7 3 . Even so, not exactly the same steels were investigated as were used in our experiments, s t i l l very similar p o l a r i z a t i o n curves were obtained. Agrawal and Crowe found that the s i g n i f i c a n t features of the curves for Fe-Ni-Cr a l l o y s depend on the c h a r a c t e r i s t i c features of the pure components iron, nickel and chromium. Therefore, I consider p o l a r i z a t i o n in the anodic d i r e c t i o n , the f i r s t active peak (sometimes observed as double peak) above the free corrosion potential can be a t t r i b u t e d to the dissolu t i o n of iron at the lower and n i c k e l at the higher p o t e n t i a l . The following passive region i s due to a protective (Fe,Ni,Cr) -oxide f i l m . The peak following t h i s passive region (transpassive peak) i s found to be due to the d i s s o l u t i o n of chromium. At potentials more posit i v e than the transpassive peak an (Fe,Ni) oxide f i l m is responsible for the passive behavior. (Compare Figure 82) Figure 82: Discussion of the p o l a r i z a t i o n curve in terms chemical reactions in the sodium hydroxide solutions r c CD o Q _ (Fe.Ni)-Oxide } Chromium dissolution (Fe,CrNi) Oxide Nickel dissolution Iron dissolution Current - density 106 In Fig.48, pg.68, i t can be seen that a l l three steels behaved very s i m i l a r l y in the sodium hydroxide s o l u t i o n . The current d e n s i t i e s at the transpassive peak were comparable to the current densities found in the carbonate/bicarbonate solution. However, the peak currents at the active peak were about 10 times lower in the sodium hydroxide solution and both the active and transpassive peaks were displaced 200mV towards more negative potentials. These e f f e c t s were mainly due to the d i f f e r e n t pH of the s o l u t i o n . At potentials above and below the transpassive peak, s t e e l "0" seems to show s l i g h t l y higher current d e n s i t i e s compared to "A" and "R", which was also observed in the carbonate/bicarbonate solution. The active peak (double peak) appeared s l i g h t l y d i f f e r e n t for steel-"0" as compared to "A" and "R". In steel-"0" the second nose, which is a t t r i b u t e d to nickel dissolution was smaller. This may, therefore, be explained by the considerably lower nickel content of steel-"0". In a l l test solutions the free corrosion p o t e n t i a l for steel-"0" was s l i g h t l y more negative than both ste e l s "A" and "R". In the a c i d i c C O 2 / H 2 O solution, i t was found that steels "A" and "R" gave very much higher current densities than steel "0". However in the carbonate/bicarbonate and sodium hydroxide solutions similar current d e n s i t i e s were observed at the active peak. Therefore, d i f f e r e n t 107 d i s s o l u t i o n processes may be occurring on each s t e e l in the a l k a l i n e and acidic solution respectively. With the remelted steel "R" (doped with i n c l u s i o n s ) , i t was found that the transpassive peak did not appear in the as cast ot hot r o l l e d condition but only in the heat treated condition. Since this peak i s attributed to chromium d i s s o l u t i o n , an enrichment of chromium at c e r t a i n locations must take place during the heat treatment. The analysis of the carbides (compare chapter 2.2.8. on page 33) showed an enrichment in a l l o y i n g elements. Therefore, at the transpassive peak a p r e f e r e n t i a l d i s s o l u t i o n of chromium r i c h carbides may take place. As mentioned e a r l i e r , no major difference in p o l a r i z a t i o n behavior between the CaS doped s t e e l and the MnS doped steel was v i s i b l e . The difference in current densities at the active peak between the "R"-steel and both doped steels may be caused by the in c l u s i o n s . However, since the microstructure of the doped st e e l i s not exactly the same as the "R"-steel, these conclusions are not completely v e r i f i e d In further studies, "R"-steel needs to be tested in the remelted and heat treated state but without any additions of inclusions. The decreased current density at the transpassive peak during the second p o l a r i z a t i o n test in Fig.49, pg.68 may be a t t r i b u t e d to the di s s o l u t i o n of most chromium carbides on the metal surface during the f i r s t test. 108 4.2. Slow s t r a i n rate vs. fracture mechanics tests As mentioned in the introduction, SSRT tests are used for fast determination of s e n s i t i v i t y to stress corrosion cracking. Questions often a r i s e as to whether findings from SSRT-tests can r e a l l y be applied to samples with precracks. In most cases suspicion only arises when no cracking i s found in the SSRT-test, because one worries about a probable lack of i n i t i a t i o n s i t e s . In our experiments mainly the reverse was found and many cracking signs in SSRT-tests could not be reproduced in the fracture mechanics t e s t s . A possible explanation for t h i s i s given by Herbsleb et a l 74,75 yfoo found that in some metal/environment systems SCC only occurs when a c r i t i c a l loading - or st r a i n rate i s present, which allow a ce r t a i n p l a s t i c deformation at the crack t i p . B r o s e ^ also related s t r a i n at the crack t i p to SCC and compared the SCC i n i t i a t i o n resistance of smooth and notched samples. Takemoto?? combined the slow stra i n rate idea with fracture mecanics tests (analogous to our "dynamic" SCC experiment) and found a method for a quick determination of KISCC v a ^ u e s « H e found a c r i t i c a l loading rate below which the s t r e s s - i n t e n s i t y for stress corrosion crack i n i t i a t i o n no longer was dependent on the loading rate. With these explanations for stress corrosion cracking one could explain why, for example, steel-"0" showed no cracking in the s t a t i c fracture mechanics test (1M NaHCC>3 / 109 1M Na2C03) and only l i t t l e cracking ( i f any) in the dynamic fracture mechanics t e s t s . According to Takemoto's findings, the strain rate during the dynamic test must have been too fast to show appreciable cracking. In the s t a t i c test the p l a s t i c flow conditions at the crack t i p may not have been s u f f i c i e n t for any f i l m rupture mechanism to take place and even higher stress i n t e n s i t i e s (approaching K j r) would be necessary for cracking. Another factor could be the conditions of the environment (potential, pH) at the crack t i p which may have been quite di f f e r e n t in the SSRT-test (short crack) as compared to the fracture mechanics test (long crack). This could be responsible for the r e l a t i v e l i t t l e cracking of "A" and "R" steels in the SSRT-test with NaOH solutions, whereas with the precracked samples severe cracking was found in both s t e e l s . Also, differences in environment conditions may have accounted for the observations in 1M carbonate/bicarbonate (compare F i g . 67,68, pg.92) where the corrosive attack at the crack t i p was quite d i f f e r e n t from the crack mouth. This may be p a r t i a l l y explained with the longer exposure time of the crack mouth to the environment than that of the crack t i p , but d i f f u s i o n c o n t r o l l e d differences in environmental conditions could have caused th i s difference because of the narrow crack geometry. From these r e s u l t s , i t can be seen how very important the combination of environment ( i n c l . potential) material 1 10 and stress state can be for stress corrosion cracking to occur. The blunting observed with steel "A" and " R " in the CO2/H2O solution during the fracture mechanics test can be explained by strong corrosive attack, with no passivation. Steel-"0" showed no cracking in t h i s solution at a l l . From the p o l a r i z a t i o n curves (Fig.43, pg.64) i t can be seen that s t e e l "A" and "R" give about 10 times higher corrosion currents than s t e e l - " 0 " and therefore perhaps the corrosive attack at the crack t i p was noticed only in the two 3.5%NiCrMoV s t e e l s . If the crack blunting was due to a pure d i s s o l u t i o n process, then the "crack-advance" (v) should be related to the anodic current density via Faraday's law of e l e c t r o l y s i s : v = i - W ' / F - p ...(13) where F = Faraday constant (9.65*10 4A*S equiv" 1) p = density of the a l l o y ( 8.0•103kgm"3) W = equivalent weight of a l l o y i = anodic current density v = crack v e l o c i t y At 95°C and the free corrosion potential in C 0 2/H 2 0 , v = 2.4*l0" 1 1m/s , W = 2.8*10~2 kg (based on oxidation state of F e 2 + ) . A c a l c u l a t i o n of the anodic current density (i ) with these values gives about 0.7 A/m2. This value i s in good agreement with an anodic current density (Tafel 111 extrapolation of po l a r i z a t i o n curve) of about 1-1.5 A/m2 found from polarization curves of st e e l "A" and "R" (Compare Fig.43, pg.64) 4.3. Comparison with Literature data Most stress corrosion cracking data r e l a t i n g to carbonate/bicarbonate solutions was obtained on mild steel or high strenth low a l l o y (HSLA) l i n e pipe s t e e l . Tests were mainly conducted with the constant load or SSRT-t e c h n i q u e 2 5 ' 2 7 ' 2 8 ' 3 0 ' 4 1 ' 7 8 . The potential where cracking was observed l i e s between about -600 to -800mVgCE, depending on temperature and pH. Cracking in this range is reported to be intergranuler for the mainly f e r r i t i c microstructures and maximum crack v e l o c i t i e s are r e p o r t e d 2 7 ' 3 ^ as 2'l0 _ 9m/s. This compares well with our SSRT-tests studies, but, a d d i t i o n a l l y , we found similar cracking v e l o c i t i e s at much more noble p o t e n t i a l s . No published data were found to compare with the cracking rates of fracture mechanics samples in the carbonate/bicarbonate solution. In the aci d i c CO2/H2O system, investigations have been conducted on low a l l o y steels 2** and on 3.5%NiCrMoV s t e e l s 3 8 ' 3 9 . In the low a l l o y s t e e l s , cracking was found only at high C0 2 pressures and elevated temperatures. No quantitative cracking results were given. 1 12 In SSRT-tests, Mueller J° observed a strong corrosive attack of the specimen with many secondary cracks. The fracture surface showed no intergranular cracking. These findings are comparable with our results since testing conditions were very s i m i l a r . However, S p e i d e l 3 9 , found intergranular SCC in C O 2 containing water at 160°C when using fracture mechanics te s t s . Maximum cracking rates of 7*l0~ , 0m/s were reported. In our experiments at 95°C no cracking was detected, but rather strong crack t i p diss o l u t i o n took place. From the mechanical testing of our st e e l , no major difference in the p l a s t i c deformation behavior at 95°C and 160°C could be found. Therefore, differences between Speidel's results and the present study could be due to environmental e f f e c t s . For example, a higher pH of the solution may have allowed passive f i l m formation and prevented strong general corrosive attack. Our intergranular cracking rates of 2'10 _ ,° m/s obtained in the 1M carbonate/bicarbonate solution would be comparable to the findings of Speidel in the C0 2/H 20 system. In caustic sodium hydroxide solutions, tests of rotor steels have been reported with SSRT, constant load, and fracture mechanics t e c h n i q u e s 2 ' 3 ' 3 4 ' 7 9 . From SSRT-tests, i t was found that the minimum NaOH concentration required to produce SCC decreases with increasing temperature. Cracking was observed at 65°C above 0.8 Mol/lt and at 160°C above 3•10~2 Mol/lt cracking was observed. At constant NaOH concentration the crack growth rates decreased with 1 1 3 decreasing temperature. At 160°C crack growth rates of 4 - l0" l 0m/s to 3-1(T8m/s were found with precracked samples in 0 .3 molar NaOH solutions. The crack path was i n t e r g r a n u l a r 3 . Somuah et a l 3 4 did slow s t r a i n rate tests at 120°C in 12M NaOH, with additions of NaCl, and found no sign of an intergranular crack path. The tests were done at the free corrosion p o t e n t i a l . At more noble po t e n t i a l s , at the active peak and in passive regions, intergranular fracture was found. With fracture mechanics tests, intergranular corrosion was found at the free corrosion p o t e n t i a l . Crack growth rates were calculated to be -1 0~8m/s. Since most of the so far mentioned experiments (not including the present study) were done either at higher temperatures or with stronger NaOH solutions than our study, higher cracking rates were obtained in comparison to our measurements. However, both the crack path and cracking in the passive region was confirmed by our experiments. Estimated crack growth rates from cracked low pressure rotor d i s c s 1 ' 2 ' 3 ^ compare well with our r e s u l t s . In general, a large scattering of data i s reported for service f a i l u r e s becouse crack v e l o c i t i e s are estimated from the time of service, which does not include any incubation time for cracking or the e f f e c t of a changing Kj on growth rates. Typical crack growth rates in service have been estimated to vary from I0~ 1 0m/s to I0~ 1 2m/s. 1 1 4 A comparison of B r i t i s h and U.S. f i e l d experience data is given in the report of McMinn et a l 3 . A d d i t i o n a l l y , laboratory studies in environments ranging from steam to concentrated sodium hydroxide solutions are also l i s t e d . V. SUMMARY The microstructure of three d i f f e r e n t rotor s t e e l s was investigated with the transmission electron microscope. Stress corrosion studies were conducted in carbonate/ bicarbonate and sodium hydroxide solutions by means of slow s t r a i n rate t e n s i l e and fracture mechanics techniques. The hot r o l l e d 1%CrMoV ste e l - n O " showed a f e r r i t i c b a i n i t i c microstructure with elongated MnS inclu s i o n s present. The heat treated 3.5%NiCrMoV steels "A" and "R" revealed an upper b a i n i t i c and lower b a i n i t i c microstructure r e s p e c t i v e l y . Inclusions in "R " were much bigger than in "A", but less numerous. They were i d e n t i f i e d in both cases as CaS. In steel "A", oxide inclusions were also found. With slow s t r a i n rate t e n s i l e tests, severe cracking of a l l three steels was found in the carbonate/bicarbonate s o l u t i o n s . The deepest cracks were observed at test p o t e n t i a l s near or at the active to passive t r a n s i t i o n . In the caustic sodium hydroxide solutions only very l i t t l e cracking could be found. 1 15 SCC tests with fracture mechanics specimen showed cracking of a l l three steels in the sodium hydroxide so l u t i o n , whereas in the carbonate/bicarbonate environment cracks were found mainly in the 3.5%NiCrMoV s t e e l . SCC test s with s t e e l "0" and "A" showed cracking in NaOH solutions at various potentials in the active and passive region. At the transpassive peak, where chromium d i s s o l u t i o n takes place, crack growth rates s i m i l a r to the active peak were found. At a given stress i n t e n s i t y , the crack growth rates in NaOH were the smallest in s t e e l "A" and about a factor of 5-6 times larger in s t e e l "0". Cracking rates for "R" were similar to st e e l "0". The crack path in the steels "A" and "R" was mainly intergranular in a l l t e s t s . The crack path in s t e e l "0" could not be c l e a r l y defined but could have followed the f e r r i t e / b a i n i t e boundaries. In the a c i d i c C O 2 / H 2 O solution, strong d i s s o l u t i o n processes took place at the crack t i p of s t e e l "A" and "R", but no " r e a l " cracking could be observed. The e f f e c t of d i f f e r e n t inclusions could not be completely analysed, but from p o l a r i z a t i o n studies i t could be concluded that the eff e c t was rather minimal. The slower cracking of st e e l "A", as compared to "R", could be due to inclusions but i s more probably the result of d i f f e r e n t s t e e l composition (Mn,Si) and microstructure. 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