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1995

A Study of the Stress Corrosion Crack InitiationStage in Alpha-Brass.Kun LianLouisiana State University and Agricultural & Mechanical College

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A STUDY OF THE STRESS CORROSION CRACK INITIATION STAGE IN ALPHA-BRASS

A Dissertation

Submitted to the Graduate Faculty of the Louisiana State University and

Agricultural and Mechanical College in partial fulfillment of the

requirements for the degree of Doctor of Philosophy

in

The Interdepartmental Program in

Engineering Science

by Kun Lian

B.S., South China Institute of Technology, GuangZhou, China, 1982 M.S., Louisiana State University, 1991

August 1995

UMI Number: 9609102

UMI Microform 9609102 Copyright 1996, by UMI Company. All rights reserved.

This microform edition is protected against unauthorized copying under Title 17, United States Code.

UMI300 North Zeeb Road Ann Arbor, MI 48103

ACKNOWLEDGMENTS

The author would like to extend his sincere gratitude to the following

individuals: Dr. Efstathios I. M eletis, his major Professor, for his advice, tenacious

drive, help, and valuable guidance and suggestions throughout all stages o f this

research. W ithout the help o f Dr. Meletis, this research would not have been possible.

He would also like to express his appreciation and thanks to Dr. Aravamudhan Raman

for his enthusiastic help when the author came to the USA, and during the time he

spent at Louisiana State University. Appreciation and acknowledgm ents also go to Dr.

R obert J. Gale, Dr. Raul E. Macchiavelli, Dr. Su-Seng Pang, Dr. En M a, Dr. R oger R.

M cNeil for their support, time and input.

To his wife Cuihong Tao, he would like to express his deepest felt thanks for

her love, support, understanding , encouragement, and self-sacrifice. His son Guan

Lian, he thanks for his love.

H e wishes to thank his parents and the rest o f his family for their love, and

support.

H e also acknowledges the support by Dr. N. Sarkar and the Louisiana State

University School o f Dentistry, Operative Departm ent, Biomaterials Division for their

support during his study. His thanks also go to W uhan Research Institute o f M aterials

Protection, M inistry o f M achinery Industry, People’s Republic o f China for the support

and understanding.

TABLE OF CONTENTS

A C K N O W L E D G M E N T S ........ ii

L IS T O F T A B L E S ............................................................... vi

L IS T O F F IG U R E S ......... xii

A B S T R A C T ................................................................................................................................ xi

C H A P T E R

I. IN T R O D U C T IO N .................................................................................................... 1

H . O B JE C T IV E S ..............................................................................................................4

H I. R E V IE W O F P R E S E N T U N D E R ST A N D IN G ............................................. 5

A. Historical Background ................................... 5

B. Environm ent-Induced Cracking Phenom ena ............................................ 6

1. Stress Corrosion Cracking ................................................................ 6

2. H ydrogen Embrittlement ................................................................. 10

3. Liquid M etal Embrittlement ........................................................... 13

4. Corrosion-Fatigue ............................................................................. 14

5. W ear-Corrosion Cracking .............................................................. 15

C. Stress Corrosion Cracking in a -B rass ..................................................... 16

D. Present Stress Corrosion Cracking M odels ............................................ 20

1. Dissolution M echanisms ................................................................. 21

2. Film-Induced Cleavage M odel ...................................................... 21

3. Surface Mobility M echanism ........................................................ 22

4. Corrosion-Assisted Cleavage M odel ........................................... 24

5. Corrosion Enhanced Plasticity M odel ........................................ 26

6 . Hydrogen-Induced Cleavage M odel ........................................... 27

E. Corrosion-Deform ation Interactions ........................................................ 28

F. Stress Corrosion Crack Initiation Stage .................................................. 33

1. Deform ation Patterning U nder Environm ental Influence ....... 34

2. Effects o f Electrolyte on Deform ation Patterning .................... 37

IV . E X P E R IM E N T A L P R O C E D U R E S ................................................................ 41

A. M aterial ...... 41

B. Stress Corrosion Cracking Testing ........................................................... 41

C. Corrosion Testing ............... 44

D. Characterization ............................................................................................. 45

1. Scanning Electron M icroscopy ........................ 45

2. Transmission Electron M icroscopy ....................................... 47

3. X-Ray Diffraction ............................................................................. 47

4. Analysis o f Experimental D ata .......................... 48

V. R E SU L T S ........................................................................ 49

A. Evolution o f Deform ation Patterning During SCC .............................. 49

B. Effect o f the SCC Electrolyte .................................................................... 58

C. Dislocation Configuration During TG SCC ....................................... 69

V I. D ISC U SS IO N ..................................................... 75

A. Deform ation Evolution During TGSCC Initiation ............................... 75

1. Environm ent-Induced Surface Plasticity ........... 75

2. Effect o f Anodic Dissolution ......................................................... 77

B. Effect o f Electrolyte Type on TGSCC .................................................... 78

C. Dislocation Arrangement During SCC ............ 80

D. Crack Initiation and Propagation ................ 81

E. Proposed Environm ent-Induced Deform ation

Localization M echanism .............................................................................. 83

1. Vacancy-Dislocation Interaction ............. 87

2. Vacancy Contribution to Crack Propagation ............................ 90

iv

V II. C O N C L U S IO N S .................................................................................................. 96

Vm . SUGGESTED FUTURE RESEARCH ................................................ 98

R E F E R E N C E S ........................................................................................................................ 99

A PPE N D IX E S

A N O M E N C L A T U R E ............................................................................... 110

B H Y D R O S T A T IC STR ESS A N D P L A S T IC Z O N E

C A L C U L A T IO N ........... 113

C P O T E N T IO D Y N A M IC T E S T R E S U L T S O F

a-B R A S S IN D IF F E R E N T E L E C T R O L Y T E S .......................... 115

V IT A .......................................................................................................................................... 119

v

LIST OF TABLES

Table 1. Tafel Slopes and Corrosion Rate o f a -B rass under Stressand Stress Free Conditions........................................................................... 59

Table 2. Strain-to-Initiation, Strain-to-Fracture and Slip Band Spacingo f a -B rass imthe Different Environm ents................................................ 62

Table 3. Calculated Values o f Near-Surface Region Strain.................................. 76

vi

LIST OF FIGURES

F ig u re 1. Schematic plot o f the three stages o f SCC as a function o f stressintensity factor........................................................................................................... 7

F ig u re 2. Schematic anodic polarization curve showing zones o fsusceptibility to stress corrosion cracking.......................................................... 9

F ig u re 3. The relationship betw een the concentration o f copper and(a) the rate o f dissolution and (b) the time-to-failure for Cu-30Zn in 15 M aqueous am m onia" .............................................................................. 19

F ig u re 4. Schematic illustration o f the elements o f the film-inducedcleavage mechanism o f crack propagation........................................................ 23

F ig u re 5. Schematic description o f atomic processes associated withthe surface mobility m odel . 111 ........................................................................... 25

F ig u re 6 . Schematic descriptions o f hydrogen-induced cleavage m odel . 117 ............. 29

F igure 7. Slow strain rate test results on Type 304 stainless steel in boilingM gC k solution . 120 ........................................................................ 31

F ig u re 8 . Effect o f anodic dissolution on creep in pure copper . 121 ............................ 32

F ig u re 9. Relationship betw een plastic strain and dislocation density at thenear-surface region and the interior o f the material in naval brass . 13 ..... 36

F ig u re 10. Schematic diagram o f (a) the experimental specimen and (b) experimental set-up for SCC testing. (All dimensions are in millimeters.)...................................................................................................... 43

F ig u re 11. Schematic illustration o f param eters relating to emergent slipbands on specimen surface................................................................................ 46

Figure 12. Slip band spacing as a function o f plastic strain in a-b rasstested in 5 M N H 4OH solution plus Cu(N03)2............................................. 50

Figure 13. M axim um slip step height as a function o f plastic strain ina-b rass tested in 5 M N H 4OH solution plus Cu(N03)2............................. 51

Figure 14. Scanning electron m icrographs showing a-b rass specim ensurface appearance after straining at lx lO ' 5 s ' 1 in 5 M N H 4 OH solution plus Cu(NC>3 ) 2 (a) zp= 0.002, (b) ep= 0.017,(c) ep = 0 . 1 0 , (d) slip band m orphology in a specim en thatwas strained to fracture in laboratory air, e 0. 75.................................... 53

Figure 15. Scanning electron m icrograph o f an area close to the final stress corrosion fracture surface showing crack initiation patterning (a-b rass tested in 5 M N H 4 O H ).................................................... 55

Figure 16. Scanning electron fractograph showing stress corrosion crack initiation (indicated by arrows) and propagation sites (a-b rass tested in 5 M N H 4 O H )....................... 56

Figure 17. Scanning electron m icrograph showing flat {111} facets producedduring stress corrosion testing (a-brass tested in 5 M N H 4 O H ).............. 57

Figure 18. Scanning electron m icrograph showing (a) surface appearance o f a specim en after 10 m inutes im m ersion in 5 M N H 4OH solution w ith Cu(N03)2, and (b) h igh m agnification o f an area shown in (a)................ 60

Figure 19. Scanning electron m icrograph showing surface appearance o f aspecim en after 10 m inutes im m ersion in 0.1 M CUSO4 solution 61

Figure 20. Fracture surfaces produced by testing (a) in 5 M N H 4OH with C u (N 0 3)2, (b) in 1 M N a N 0 2, (c) in 0.1 M C u S 0 4 and (d) in laboratory a ir ...................................................................................... 64

Figure 21. Slip band spacing developed during SCC testing in thethree electrolytes (solid signs represent crack initiation).......................... 65

F ig u re 22. Typical surface appearance o f slip bands developed in the threeelectrolytes after testing for a total strain o f 0.006 in (a) 5 M N H 4OH with Cu(N03)2, (b) 1 M N a N 0 2 and (c) 0.1 M C u S 0 4 solution 6 6

F ig u re 23. The early stages o f the stress-strain behavior o f a -b rass inthe three electrolytes and in laboratory air.................................................... 67

F ig u re 24. The full w idth ha lf m axim um (FW HM ) m easurem ents o f the {111} peak obtained by X-ray diffraction o f a -b rass specim ens tested in different environm ents....................................................................... 6 8

F ig u re 25. D islocation configuration in an a -b rass SCC specim en(a) a region near the stress corrosion fracture surface showing intensified pile-ups on { 1 1 1 } slip planes and (b) a pure mechanically deform ed region showing a m ore hom ogeneous dislocation distribution................................................. 70

F ig u re 26. A m icrocrack in a -b rass following the < 211> direction(intersection o f the { 1 1 1 } w ith the {0 1 1 } crack plane)............................... 71

F ig u re 27. D islocation configuration in a pure C u SCC specim en (a) a region near the stress corrosion fracture surface showing cellular dislocations w ith traces o f planar slip in the w alls o f the dislocation cells and (b) a pure m echanically deform ed region showing a more hom ogeneous dislocation substructure.......................................................... 73

F ig u re 28. Transm ission electron m icrograph showing dislocation pile-ups associated with dissolution m icrocracks in pure copper. The m icrocrack tip is shown on the upper part o f the m icrograph................. 74

F ig u re 29. Schematic representation o f proposed environm ent-induced dislocation em ission (a) perfect lattice; (b) vacancy generation due to anodic dissolution; (c) generation o f a dislocation and out o f surface displacem ent by one interatomic distance; (d) and (e) form ation o f planar dislocation arrangement; and (f) crack nucleation by opening o f a dislocation pile-up in a slip band ................. 85

F ig u re 30. Schematic representation o f proposed m ultiple dislocationnucleation and fracture surface characteristics produced duringcrack propagation. ............................................................................................. 8 6

F ig u re 31. (a) Stress distribution in front o f a blunted crack in a-brass;(b) Transient divacancy concentration profiles o f stress free(600 s) and stress-assisted (60 s and 600 s) diffusion (C s denotesthe divacancy concentration at the surface.)................................................ 94

F ig u re 32. Potentiodynam ic test results o f a -b rass in 5 M NH 4 0 Hplus Cu(N 03)2 solution...................................................................................... 114

F ig u re 33. Potentiodynam ic test results o f a -b rass in 1 M N a N 0 2 solution 115

F ig u re 34. Potentiodynam ic test results o f a -b rass in 0.1 M C u S 0 4 solution 116

x

ABSTRACT

The objective o f the present work was to provide an insight to the nucleation and

evolution o f deformation patterns occurring during transgranular stress corrosion cracking

(TGSCC) and produce new alternatives for addressing the nature o f the embrittlement

process. Flat, tensile a-brass (72Cu-28Zn) specimens were tested in 5 M NH4OH, 0.1 M

CUSO4 and 1 M N aN 0 2 solutions at a strain rate o f lx lO " 5 s_ 1 . Slip band spacing (SBS)

and slip band heights (SBH) were measured as a function o f strain by conducting

interrupted experiments in the SCC environments and were compared with those developed

during laboratory air experiments. The presence o f the TGSCC-causing environment

during straining was found to promote localized plastic deformation at the near-surface

region, induce strain hardening and more importantly to produce an entirely different

deformation pattern compared to that developed in laboratoiy air. The deformation evolved

in the presence o f the TGSCC electrolytes was highly localized, exhibiting a dense SBS but

coarse SBH. Also, a periodicity was exhibited by the crack initiation process. The amount

o f localized strain developed at the specimen near-surface region prior to nucleation o f

stress corrosion cracks was found to be equivalent to the strain required for ductile fracture

o f the material in air, suggesting the existence o f a fundamental fracture criterion. In view

o f the present observations, an environment-induced deformation localization mechanism is

introduced to explain TGSCC initiation and propagation. The main elements o f the

proposed mechanism are: (i) strain localization due to corrosion instability and periodicity;

(ii) vacancy-induced dislocation emission at the surface region and (iii) vacancy-dislocation

interaction localizing deformation and modifying dislocation arrangement.

CHAPTER I. INTRODUCTION

The problem o f stress corrosion cracking (SCC) has been attracting significant

interest for the last 100 years. A number o f critical engineering com ponents and

structures suffer daily from this and similar types o f environment-induced cracking

(EIC) phenomena with catastrophic consequences in terms o f human life and economic

losses. Fundam ental mechanistic aspects and advances in understanding this

environment-induced embrittlement have been addressed recently . 1 ,2 Even though

several theories have been invoked, there is no broadly accepted mechanism at the

moment to explain these types o f phenomena. The way the environment causes

embrittlement during SCC still remains unresolved and this is an issue o f significant

engineering concern and academic interest.

The m ost significant characteristic o f SCC is that normally ductile alloys

undergo an essentially brittle fracture at relatively low stresses. A specific chemical

environment/material combination and the presence o f a tensile stress are the necessary

requirem ents for SCC to occur. Crack propagation can occur in an intergranular (IG)

and/or transgranular (TG) fashion. Regarding IGSCC, a consensus exists that cracks

initiate at and propagate along grain boundaries according to the film rupture 3 ,4 or the

slip dissolution m odels . 5 In these models, the crack grow s by preferential anodic

dissolution o f the grain boundary region. In contrast, T G cracks generally initiate at

emergent slip bands and propagate on crystallographic planes6 ' 8 by an environmentally-

induced cleavage process, although the way in which the environment induces cleavage

in normally ductile alloys has not been determined yet.

O f particular interest is the case o f TGSCC o f the otherwise ductile face-

centered cubic (FCC) metals and alloys, since they have multiple slip systems and are

1

2

characterized by high dislocation velocities. As has been shown by M eletis and

Hochm an , 6 ' 8 an im portant element o f the underlying TG SCC mechanism is the

ciystallographic nature o f cracking. It has been shown in the above studies that in all

FCC metals and alloys, TG cracks initiate on {111} slip planes but propagate on {0 1 1 }

for non-ferrous metals and on {001} for austenitic stainless steels. All fracture surfaces

consist o f parallel and relatively smooth facets separated by crystallographic or fan­

shaped steps. Acoustic emission 9 ,1 0 and electrochemical transien t" studies suggest

that crack propagation is discontinuous, and mating fracture surfaces are matching and

interlocking . 12 Also, crack-arrest markings are usually observed on the fracture

surfaces, and they have been correlated to the discontinuous nature o f crack

propagation.

All the above physical evidence suggests that cracking occurs by a micro

cleavage process. Thus, it is only logical to expect that the deform ation m ode o f the

material would play a key role in the embrittlement process. Previous studies have

shown that the presence o f the SCC environment produces a much higher dislocation

density at the near-surface region , 13 and cracking occurs when a critical value o f

dislocation density is reached . 14 In addition, the fact that cracks initiate on {111} slip

planes, but then change direction and propagate on another plane, suggests that the

param eters involved are contributing to a different extent during crack initiation and

propagation. In the presence o f a growing crack, relatively high stress intensities

prevail and the re-initiation stage is anticipated to be o f short duration and thus difficult

to observe. D ue to their transient character, stress/environment interactions are

expected to have longer lifetime during the very first crack initiation event (occurring

on smooth material surfaces, where a crack has not been formed yet). This provides

the opportunity to study the evolution o f the deform ation process in the presence o f a

SCC environment and to identify possible differences with the deform ation pattern in

the absence o f the cracking-causing environment. Second, it allows the investigation o f

the role o f the plastic deform ation on the TG stress corrosion crack initiation stage, in

an effort to shed m ore light on the fundamental mechanism o f the embrittlement

process. The nature o f the initiation process and its impact on crack propagation may

be fundamental and despite its im portance the crack initiation has attracted very little

attention in the past. A significant role in the SCC initiation process is played also by

the type o f the environment. The fact that in the same metal or alloy certain

environments can cause faster crack initiation compared to other environments that are

m ore corrosive has not been understood yet. The present investigation was conducted

in an effort to provide a better fundamental understanding o f the crack initiation

process during TG SCC and explore its consequences on the crack propagation stage.

CHAPTER II. OBJECTIVES

The specific objectives o f the present research are:

(a) to investigate the nucleation and evolution o f deform ation patterns

occurring during the initiation stage o f TGSCC;

(b) to study the effects o f various TGSCC-causing electrolytes on the

deform ation patterning and correlate them to cracking susceptibility and

the degree o f embrittlement exhibited in the particular electrolyte; and

(c) in view o f the above findings, to produce new insights and new

alternatives for addressing and better understanding the nature o f the

embrittlement process. It is hoped that this investigative effort will

contribute new knowledge to m ore closely defining the operating SCC

mechanism in FCC metals and alloys.

4

CHAPTER III. REVIEW OF PRESENT UNDERSTANDING

A. Historical Background

It is believed that the first docum ented environment-induced cracking is the

“season cracking” o f brass in ammonia-bearing environments in the early tw entieth

century . 1 5 ,1 6 Extensive early w ork on season cracking o f brasses w as carried out by

M oore, Beckinsale, and M allinson . 17 As a result o f their research, an acid m ercurous

nitrate solution w as used to detect susceptibility to season cracking in brasses.

In 1909, another type o f E lC -stress corrosion cracking o f boiler steel was

reported . 18 In 1913, Andrew w as one o f the first to suggest that embrittlem ent caused

by the ingress o f hydrogen into ferric materials, which provided a possible explanation

for SCC in boiler steels . 19 A couple o f years later, Wolff2 0 proposed a fatigue

mechanism accounting for the caustic cracking o f boilers. In 1920, Griffith2 1 first gave

a quantitative relationship betw een the applied energy and crack surface area. By

studying the SCC o f aluminum alloys in 1940, Dix2 2 suggested that chemical

heterogeneity at the grain boundaries was the reason for cracking. In 1947, Keating 23

and Evans2 4 independently suggested that crack propagation occurred by alternating

mechanical and electrochemical stages, in which localized corrosion produced notches

that induced brittle fracture o f a mechanically weakened path. This mechanism requires

a region in the material that is susceptible to corrosion and mechanically weakened,

providing the site for notch o r bulk defect formation.

W ith the developm ent o f new fracture characterization techniques, it was

realized that m icrostructure features such as dislocations may play an im portant role in

the fracture mechanisms. A t the same time, those early hypotheses have been tested

5

and some o f these have gained acceptance. Today, the efforts o f explaining the EIC

are focused on tw o aspects. One is to explain the observed cracking in normally ductile

materials at relatively low stresses with little attendant deformation, and the second is

to determine the role o f specific environments that prom ote such failure.

B. Environment-Induced Cracking Phenomena

EIC can be defined as a brittle failure o r ductility reduction phenom enon in a

particular metal o r alloy when stressed in a given environment. Thus, cracking can

occur only in specific m aterial/environment systems, even though many environments

may be able to produce EIC in a certain material. The various param eters involved in

EIC may include (i) material param eters: composition, crystal structure, m icrostructure,

nature o f grain boundary, deform ation characteristics, and surface condition; (ii)

environmental param eters: nature o f ions, concentration, conductivity, and type o f

phase (i.e. liquid o r gaseous), and (iii) mechanical factors: stress intensity, strain rate,

plane stress/plane strain, loading mode, cyclic frequency, wave shape, etc. Thus, it is

evident that a num ber o f complex physical-mechanical-chemical aspects are involved in

these environm ent-induced fractures.

EIC phenom ena can be classified into the following categories:

I. Stress Corrosion Cracking

SCC can occur in a large number o f material/environment systems. The most

common ones are: brass in ammonia, stainless steel in chloride containing

environments, high strength steels in caustic and nitrate environments, aluminum alloys

in sea water, and Ti alloys in chlorides and methanol. It has been established, that a

certain relationship exists betw een SCC velocity, expressed as: V-da/dt, and the stress

intensity factor Kj, Figure 1. For engineering materials, the rate o f crack grow th in

SCC as a function o f the stress intensity factor at the crack tip shows three regions o f

crack development. In region I, cracks initiate when loading above a threshold value

log

& (

SCC

velo

city

, m

/s)

7

Stage m Rapid SCC Growth /

Stage n Crack Velocity Plateau

Catastrophicfailure

SubcriticalSCCpropagation

Stage I

Timedependent K

SCC

'SCC

1/?Stress Intensity K (MPa* m )

Figure 1. Schematic plot o f the three stages o f SCC as a function o f stress intensity factor.

8

Kiscc and crack velocity is a strong function o f Ki In region II, the crack grow th is

essentially independent o f Ki, indicating the existence o f another controlling param eter

(m ore than likely diffusion controlled). In region III, the crack velocity increases

rapidly with Ki, and unstable crack growth occurs as Ki approaches Kic, the fracture

toughness o f the material. In both stages I and II, the fracture process is influenced

significantly by the environment and generally follows intergranular and/or

transgranular paths. In region III, fracture occurs by a mixture o f environmental and

mechanical cracking (rupture); the proportion o f rupture increases until K i= K ic where

fast fracture occurs.

The electrochemical potential has a critical effect on SCC, especially in

m aterial/environment systems with an active/passive behavior. Figure 2 shows the

potential ranges where SCC occurs in relation to a potentiodynamic anodic polarization

curve for a typical active/passive metal/environment system. Three electrochemical

corrosion reaction zones basically exist in this type o f system: active, passive, and

transpassive or pitting zones. SCC usually occurs in the potential ranges that are at the

beginning or at the end o f the passive region, where the possibility for film breakdow n

exists.

Plastic strain and strain rate also play an important role in stress corrosion

cracking due to the synergistic effects o f the time-dependent corrosion reactions and

microplastic strain. SCC processes require both strain energy and corrosion reactions

in order to initiate and prom ote the crack advancement. Continuous straining exposes

fresh metal surface to the environment and functions as the catalyst to subsequent

environmental reactions. Because the corrosion reactions are chemical reactions and

have a time dependency, a fast strain rate may not allow enough time for corrosion

processes to fulfill their role, thus SCC may not be produced and a ductile fracture is

PO

TE

NT

IAL

9

Pitting

+

A

PassiveCracking zones

Active

lo g C U R R EN T DENSITY

Figure 2. Schematic anodic polarization curve showing zones o f susceptibility to stress corrosion cracking.

10

obtained. A very slow strain rate on the other hand prom otes excessive corrosion

causing crack blunting rather than SCC. A t intermediate strain rates, plastic strain and

corrosion reactions interact critically (synergistically) with each other, and SCC has the

opportunity to occur. It has been shown that slow strain rate testing provides an

excellent way to study these critical corrosion-deform ation interactions.

2. Hydrogen Embrittlement

Even when present in very small amounts, hydrogen can cause severe

embrittlement in many metals, ceramics and intermetallics. Body-centered cubic (BCC)

and hexagonal close-packed (HCP) metals are m ost susceptible to hydrogen

embrittlement (HE), with FCC metals less susceptible . 2 5 The fracture process in HE

can involve cleavage, intergranular, o r ductile fracture depending on the hydrogen

concentration and the stress level . 2 6

The main characteristics o f H E are its strain-rate sensitivity and tem perature

dependence. Unlike m ost EIC processes, HE is enhanced by slow strain rates since

m ore time is provided for hydrogen diffusion. At low and high tem peratures, H E can

be ignored but it is most severe at intermediate tem peratures. H ydrogen can be

introduced in the material in the atomic form (H) in tw o ways. These are,

electrolytically through a discharge reaction (IT+ + e- —> H ) or from a gaseous

atm osphere through molecular dissociation, followed in both cases by H absorption on

a metal surface. D ue to its small size, hydrogen occupies interstitial lattice positions

and diffuses with a high mobility. Because o f the larger interstices and close packing o f

host atom s in FCC and HCP metals than those in BCC metals, usually the mobility o f

hydrogen is less in FCC and H CP metals compared to BCC metals.

The hydrogen transport rates in association with dislocation motion can be

several orders o f magnitude greater than those occurring by diffusion alone . 2 7 In the

presence o f hydrogen, dislocations exhibit higher mobility and velocities and the

opposite effects are observed when hydrogen is removed. The effect o f hydrogen in

stressed metals is to reduce the critical stress required for dislocation motion.

Enhanced dislocation velocities resulting from hydrogen have been observed for screw,

edge, and mixed dislocations and for dislocations that were in tangles, in slip bands, and

far from other dislocations. Similar behaviors have been observed for dislocations in

FCC, BCC, and HCP crystal structures and in alloys as well as in pure m etals . 2 8 ' 3 6

The main proposed H E mechanisms are: (a) lattice decohesion, (b) hydrogen-

dislocation interaction, (c) hydride formation, and (d) hydrogen-enhanced localized

plasticity mechanism.

D ecohesion is one o f the oldest HE mechanisms. In general, decohesion

associates H E with a decrease in the interatomic bond strength. Thus, cleavage

fracture occurs when the applied stress exceeds the lattice “cohesive stress.” The

fracture resulting from HE can be transgranular or intergranular. In IG fracture, the

relevant param eters are the cohesive energy and cohesive force o f the grain boundary,

which are decreased by the presence o f hydrogen as well as the segregation o f many

other solutes. In TG fracture, the fracture surface is usually associated with significant

local plasticity.

For many years, it w as believed that hydrogen dissolved in metals did not

interact significantly with dislocations . 3 7 ,3 8 In the late 60s and early 70s, the opposite

suggestion was made, with strong supporting evidence. Experim ents have shown that

(a) the large values o f the partial molar volume o f hydrogen causes hydrogen-

dislocation interaction in terms o f elastic interactions, (b) internal friction experiments

have shown almost conclusively large hydrogen-dislocation interaction , 3 9 ' 41 and (c) high

resolution autoradiography o f tritium in metals has shown clearly and directly hydrogen

segregation at dislocations . 4 2 ' 4 5 It has been suggested that hydrogen plays an important

12

role in metal embrittlement by interacting with and pinning dislocations, thus permitting

brittle fracture . 4 6 "4 8

A num ber o f metals have dem onstrated HE resulting from stress-induced

form ation o f hydrides or other relatively brittle phases and the subsequent brittle

fracture o f these phases . 4 9 "5 2 The basic requirem ents are that these phases be stabilized

by the presence o f hydrogen and the crack-tip stress field , 5 3 ,5 4 and that the formed

phases be brittle . 55 There are also a num ber o f alloy systems that form “pseudo­

hydrides” under cathodic charging. These “pseudo-hydrides” are high-concentration

solid solutions formed in the presence o f a miscibility gap . 5 6 These systems often

exhibit embrittlement in the presence o f this high-concentration, solid solution (or by

the stress-induced formation o f this high-concentration solid solution), even though this

phase is not a true hydride; i.e., it lacks the ordering o f the hydrogen in the interstitial

positions. These stress-induced hydrides remain at the fracture surface and can be

detected in some cases , 5 7 but in other cases, these hydrides must be observed while

under stress . 5 8

Beachem first suggested that hydrogen embrittlement o f steels was in fact

associated with locally enhanced plasticity at the crack tip . 2 6 U nder an applied stress,

hydrogen distribution can be highly non uniform. The high local stress field at the tip

o f a crack reduces the chemical potential o f solute hydrogen , 5 9 and as the result o f

diffusion, the concentration is locally increased. The tip o f the crack is also the most

likely place for hydrogen entry because the plastic deform ation minimizes the surface

barriers. The resistance to dislocation motion, and thus the flow stress, is decreased by

the presence o f hydrogen. So, in the region o f high hydrogen concentration, the flow

stress is decreased and slip occurs at stresses below those required for plastic

deform ation in other parts o f the specimen; i.e., slip localization occurs in the vicinity o f

the crack tip. Thus, the flow stress can be locally reduced, resulting in localized

13

deform ation that leads to highly localized failure by ductile processes, while the total

m acroscopic plastic deform ation remains small.

3. Liquid Metal Embrittlement

Liquid metal embrittlement (LM E) was first reported in the literature in 1874.60

This type o f embrittlement is observed when many ductile metals are coated with a thin

layer o f certain liquid (or low melting point) metals and then are loaded in tension.

Time-to-failure is extremely short, with crack velocities as high as 500 cm/s being

reported for aluminum alloys and brass in the presence o f liquid m ercury . 6 1 ' 6 4 LM E

depends strongly on tem perature. Almost all solid metal/liquid metal couples can

becom e embrittled if a specific tem perature can be reached. LM E usually causes a

cleavage type o f fracture. For FCC metals fracture is mostly IG, while TG fracture

occurs in BCC and HCP metals. Actually, in most situations, a mixture o f IG and TG

fractures co-exists in both FCC and BCC metals.

The m ajor LM E mechanism is chemisorption. It is believed that liquid metal

atom s reduce the cohesive atomic bond strength by chemisorption in the stress

concentration region, allowing fracture to occur at a much low er stress . 6 5 ’6 6 Another

very interesting proposal is the surface charge model. W hen tw o electrically

conducting solids are brought into contact, an electrical double layer is established on

the metal interface, with a potential difference building up at the liquid/solid interface.

This electronic charge layer at the interface would extend only a few atomic diameters

into the solid and could have a strong interaction with dislocations . 6 7 Latanision has

shown that the hardness o f a metal (resistance to the dislocation motion) can be

influenced by the application o f an electrical potential across the surface . 6 8 He

explained this phenomenon by proposing a change in the surface energy by the double

layer potential.

14

4. Corrosion-Fatigue

The simultaneous action o f cyclic stress and environment leading to fracture is

know n as corrosion-fatigue (CF). Usually, materials undergoing cyclic loading in

vacuum at room tem perature show no fatigue limit, and they are not affected by the

cycling frequency (from 1,000 to 12,000 cycles/min.). W hen a corrosive environment

is present, materials show a definite fatigue limit, and they are also affected by the

loading frequency. As a m ajor component o f the EIC phenomena, CF differs from

SCC in several respects. First, it undergoes alternate compressive and tensile stresses.

Second, CF may occur under very low cycle frequency or in the ultrasonic frequency.

Third, the load may be sinusoidal or stochastic. Fourth, the corrosion system may be in

the active or passive state.

It has been shown by a number o f investigators that for pure metals and alloys,

fatigue crack initiation is associated with the production o f intrusions and extrusions by

conventional slip processes. W ithout corrosive environment, the intrusions and

extrusions occur after a saturation o f the cyclic stress occurs . 6 9 ,7 0 During CF processes,

the cyclic loading usually causes cyclic hardening, cyclic softening, or a combination o f

both, provided the stresses exceed the cyclic elastic limit. W hen a corrosive

environment is present, the formation o f microscopic fatigue cracks and crack

propagation can be observed under the stresses below the cyclic fatigue strength in an

inert environment. The morphology o f CF fracture is mainly TG 71 and in general

crystallographic in nature.

Some early w ork on the CF behavior o f mild steel in aqueous sodium chloride

environments suggested that (a) there is a critical corrosion rate below which the

environment does not affect the fatigue crack initiation process; (b) pitting o f the alloy

surface is not necessary to lower the fatigue resistance; (c) passivity enhances fatigue

15

resistance; (d) corrosion rates are progressively higher for higher cyclic stresses, and

(e) the surface deformation characteristics o f the steel are altered by active corrosion.

For stainless steels, corrosive attack w ithout superimposed stress often results

in pitting o f metal surfaces. The pits act as notches and produce a reduction in fatigue

strength. How ever, when a corrosive attack occurs simultaneously with fatigue

loading, a very pronounced reduction in fatigue properties is produced. Traditionally,

CF crack initiation has been associated with pitting7 2 ,7 3 or, with enhanced

electrochemical dissolution resulting from deform ation . 7 4 When films are involved,

rupture o f the protective films by fatigue generated deform ation has also been

implicated in crack initiation . 7 5 ' 7 6

5. Wear-Corrosion Cracking

W ear-corrosion cracking (W CC) is a synergistic phenomenon resulting from

conjoint interactions o f w ear and corrosion in materials. W hen mechanical w ear and

electrochemical processes occur simultaneously, materials deterioration can be

accelerated greatly. It has been suggested by Meletis and Lian7 7 that the wear

corrosion process involves: (a) the disruption and removal o f surface films leading to

exposure o f fresh metal to the environment, and (b) interactions betw een the

environment and elastically or plastically deformed sites at asperities in contact, which

can lead to cracking through wear-corrosion induced embrittlement.

During sliding-wear for example, tension/com pression stress fields are

generated at the vicinity o f the contact. Recent results show that depending on the

properties o f the substrate material and surface film, the combined w ear-corrosion

process can significantly accelerate material degradation due to synergistic effects . 7 8 ,7 9

U nder the activation o f the w ear process, the corrosion potential can be significantly

reduced, and the corrosion current can increase greatly (for example, m ore than tw o

orders o f m agnitude higher for dental amalgams ) . 7 8

16

Sliding-wear/corrosion interactions can also produce subsurface cracking. The

occurrence o f cracking has been observed in our laboratory at LSU in amalgams tested

in chloride solutions . 8 0 W CC also has been reported in ZrCte / w ater system, in which

intergranular WCC was observed under low speed sliding-wear. M eletis81 proposed a

hydrogen-assisted maximum plastic yielding mechanism to account for the observed

embrittlement and cracking in previous studies on w ear-corrosion o f amalgams.

C. Stress Corrosion Cracking in a-Brass

a-b rass is an FCC Cu-Zn alloy with less than 37% Zn. Like m ost o f the FCC

m etals and alloys, a -b rass generally exhibits a ductile behavior and it can undergo

extensive plastic deformation before fracture. How ever, SCC has proved to be one o f

the m ost frequent causes o f failure o f brass equipment in industry. For many years,

brass w as thought to fail only in ammonia environments, but later w ork has shown that

cracking also occurred in non-ammonia environments. These media are sulfate , 8 2 ' 86

acetate, formate, tartrate, hydroxide , 8 7 acidic chloride , 8 8 sodium chlorate , 8 9 m ercurous

nitrate , 9 0 ’9 1 sodium nitrite , 9 2 ’9 3 sodium fluoride , 9 4 and cupric sulfate9 5 solutions.

A lthough many kinds o f solutions can cause SCC in copper-zinc alloys, the a -

brass/aqueous ammonia system has been overwhelmingly studied. In literature, SCC o f

the a-brass/am m onia system has the longest study history, and a large amount o f

available data exists. One m ore advantage o f this system is that it can exhibit both IG

o r T G cracking modes.

Evans suggested that in the a-brass/am m onia system SCC is caused by

ammonium ions which can significantly attack only the grain boundaries or lattice

im perfections . 9 6 Dix has shown that the cracking o f brasses in ammonia solution is

electrochemical in character and that the grain boundary regions are strongly anodic . 9 7

The most common environment used in a -b rass SCC studies is aqueous

ammonia containing dissolved copper, which is the fundamental work done by

M attesson . 9 8 The general corrosion behavior o f a -b rass in such environment is

illustrated in Figure 3(a), which shows the well known relationship betw een corrosion

rate and dissolved copper content . 9 9 ' 101 In order to understand the corrosion behavior

o f a -b rass in the aqueous ammonia solution, the tw o essential redox reactions must be

considered:

Cu = Cu+ + e- ( 1 )

Cu+ = Cu++ + e- (2)

The standard potential for reaction (1) is 520 mV vs N H E, and for (2) is 159 mV vs

NHE. Reaction (2) is slower than (1), so the reaction (2) is controlling the process . 102

The electrochemical reaction is dominated by the form ation o f very stable and soluble

complexes. In moderately concentrated ammonia solution (N H 3 > 1M), cuprous ions

are overwhelmingly present as [Cu(NH 3 )2+] complexes, while cupric ions are present

either as [Cu(NH 3 )4 ++] or [Cu(NH 3 ) 5 + + ] . 103 As the ammonia concentration

increases, the second form complex becomes the main species. So, the reactions (1)

and (2 ) should be rew ritten as:

3Cu + 3N H 3 = Cu+ + Cu(N H 3 )+ + Cu(N H 3 )2+ + 3e" (3)

2Cu+ + 2C u(N H 3 )+ + 2Cu(N H 3 )2+ +9N H 3 = Cu++ + Cu(N H 3 )++ +

C u(N H 3 )2++ + Cu(N H 3 )3++ + Cu(N H 3 )4++ +Cu(NFI3 )5++ + 6 e ' (4)

There are no kinetic data available for the complex reactions. The equilibrium states o f

complex reactions are established very fast, so they do not seem to be the rate

controlling processes. Because the complexes are highly stable, the alkaline ammonia

solution can dissolve fairly large amount o f copper ions before the solubility product

for precipitation o f Cu20 is reached. The solubility product depends on pH and

18

ammonia concentration. By adding reactions (1) and (2), the total corrosion reaction is

obtained:

Cu++ + Cu = 2Cu+ (5)

This corrosion reaction will reach the equilibrium state when the ratio between the

activities o f the uncomplexed species, [Cu++]/[Cu+ ]2, equals the equilibrium constant

K , which is 1.2x10^ at 25°C .104 Because o f the formation o f complexes, the electrode

potential cannot be calculated from the N em st equation. The presence o f zinc in a -

brass does not substantially change the electrochemical behavior because zinc has

relatively fast electrode kinetics in ammonia . 1 0 5 ,1 0 6

Early experimental data showed the strong relationship betw een brass SCC

behavior and copper content in solution. Figure 3(b) shows the results under constant

load at open corrosion po ten tia l." The results indicate that (a) a minimum copper

concentration is required for cracking to initiate, (b) tim e-to-failure is greatly reduced

when dissolved copper concentration increases, and (c) the cracking path changes from

mainly T G to IG. Researchers also like to divide the electrolyte type into tarnishing

and non tarnishing solutions, which is directly related to the copper concentration in the

solutions. It has been established that in tarnishing solutions, IG cracking prevails; and

in non tarnishing solutions, the T G cracking prevails . 107

In non tarnishing solutions, the crack propagation path is dependent on the zinc

content. Brass usually shows IG cracking when the zinc content is less than 18%, and

T G cracking above this value. Zinc content also affects the slip character and the

dislocation structure. One o f the most important observations made in the 1960s and

1970s w as that the TG crack surfaces were cleavage-like in appearance with linear

features resembling fatigue striations. These features are parallel with each other,

TIM

E TO

FA

ILUR

E (s

ec)

RATE

OF

W

EIGH

T LO

SS

(mg/

min

)

19

TarnishFree

0.5

Tarnishing

IntergranularCracking

2 Trans- granular Cracking

COPPER CONTENT (g/I)

F ig u re 3. The relationship between the concentration o f copper and (a) the rate o f dissolution and (b) the tim e-to-failure for Cu-30Zn in 15 M aqueous am m onia."

20

greater than the grain size, perpendicular to the crack propagation direction, and

present atdifferent stress levels . 10 8 All these phenomena indicate that crack

propagation is a discontinuous process, and crack is blunted by plastic deformation

within its arrests. Finally, TG SCC in a -b rass is crystallographic with cracks initiating

on {111} slip planes and propagating on {110} planes. The parallel {110} SCC facets

are separated by crystallographic steps lying on alternating segments o f { 1 1 1 } planes.

This crystallography is com m on to all non-ferrous FCC metals and alloys , 8 and is a

m ajor element o f the TGSCC phenomenology.

D. Present Stress Corrosion Cracking Model

Trem endous efforts have been devoted in the recent years to develop an

understanding o f the SCC mechanism(s). M ost current mechanisms o f SCC are based

on tw o assumptions. One involves the embrittlement o f the metal as a consequence o f

corrosion interaction, and another attributes the SCC to the extremely localized

dissolution processes. All these proposed mechanisms are limited to certain systems or

to some special situations, and still are controversial. At present, there is no

mechanism that can be universally applied to all SCC; however a unified SCC theory

still remains the goal o f most researchers in the field.

A major fundamental aspect o f SCC is the IG or TG m ode o f the cracking path.

In IGSCC, grain boundaries comprise the initiation and propagation sites o f the crack

and a consensus exists that in m ost cases a preferential anodic dissolution mechanism

prevails . 12 On the contrary, as has been shown by M eletis and Hochm an , 8 TG cracking

is crystallographic in nature with the crack initiation and propagation planes not

coinciding. In the latter case, crack propagation is occurring by a microcleavage

process. However, the way in which the environm ent/stress combination induces

cleavage in normally ductile materials has not been established yet.

21

1. Dissolution Mechanisms

The m ost common dissolution mechanism has been used to explain mainly

IGSCC. This mechanism involves tw o elements: (a) preferential and rapid chemical

attack along the plane o f the boundary, which significantly reduces the fracture stress

o f the boundary plane, and (b) pile-up o f dislocations that result in a concentration o f

normal stress across the boundary plane. The unique feature o f this mechanism which

cannot be used on TGSCC is the segregation o f solutes or the precipitation o f discrete

phases that can occur at grain boundaries and that may result in electrochemical

heterogeneity. The driving force for dissolution is related to the potential difference

betw een the matrix and the segregate atoms forming a galvanic cell. Beside the simple

galvanic effects at grain boundaries, it is also possible that film characteristics at grain

boundaries are different with those within the grains. This also can result in preferential

dissolution.

The slip-step dissolution mechanism has been proposed to account for TGSCC.

Slip-steps resulting from plastic deformation emerge at the metal surface. The

protected surface film is disrupted by these slip-steps, exposing the reactive surface to

the environment. Then, the exposed area will be preferentially corroded and become

the crack initiation site. For such a mechanism to operate, the slip-step height should

exceed the thickness o f the surface film.

Another dissolution mechanism that has been introduced by Swann and

Pickering , 1 0 9 is the tunnel model. This model suggests that cracks are initiated at slip

steps by forming arrays o f fine corrosion tunnels which grow in all directions until the

remaining metal ligaments fail by ductile rupture.

2. Film-Induced Cleavage Model

The film-induced cleavage model proposed by Sieradzki and N ew m an 10 is based

on early w ork by Edeleanu and Forty, and emphasizes the role o f dealloyed surface

layers that can initiate microcracking in some materials (e.g. Cu-Zn). As has been

shown by atom ic modelling studies, a brittle crack can initiate in a thin surface film and

obtain a velocity fast enough to penetrate a small distance into the underlying ductile

metal matrix prior to being arrested (Figure 4). The surface film m ust reform at the

crack tip surface before a new burst o f brittle crack grow th is possible. This surface

film could be an oxide, dealloyed layer, or any other kind o f surface layer. The extent

o f the “film-induced cleavage” may be governed by the film-matrix misfit, the strength

o f bonding across the film-matrix interface, the film thickness, film ductility, and the

film structure. It would be reasonable to suppose that hydrogen absorption may also

play an im portant role in such cleavage process. In practice, m ost passive films are

very thin and hydrated, and the dealloyed layers are also unlikely to have sufficient

adhesion to the substrate or have inherent brittleness to sustain cracking. Another

criticism for this mechanism is that it has difficulty in explaining branching o f

propagating cracks.

3. Surface Mobility Mechanism

The surface mobility mechanism, proposed by Galvele , 11 0 suggests that the

advance o f stress corrosion cracks may be modeled by a process that involves surface

diffusion o f species from a stressed cracked tip. Surface diffusivities are significantly

enhanced due to the contam ination o f the metal surfaces from the environment. This

model can be used to explain HE, SCC, LME. The basic principle o f this model is

dem onstrated in Figure 5. An atom at the crack tip is transported by surface diffusion

from its highly stressed location at the crack tip to a new less-stressed site on the crack

walls.

The high stress concentration at the crack tip reduces the free energy o f

vacancy formation, and the equilibrium vacancy concentration at the crack tip is

23

Brittle crack: initiates in brittle film.

P ro p a g a tes in ductile crack tip metal.

Figure 4. Schematic illustration o f the elements o f the film-induced cleavage mechanism o f crack propagation.

24

increased. The surface mobility mechanism assumes that under the action o f the

environment, only the first atomic layers o f the metal are susceptible to measurable

movement. The stress concentration at the tip o f a crack generates a very localized

vacancy deficient region. Every time the stressed lattice at the crack tip captures a

vacancy, the crack propagates by one atomic distance, and a surface depletion o f

vacancies will be created. The diffusion o f those vacancies will be the rate controlling

process, leading to the crack propagation . 111 The main criticisms for this model relate

to relatively low vacancy diffusivities inside the material compared to the observed

SCC velocities and its difficulty in explaining crack branching.

4. C orrosion -A ssisted C leavage M odel

The corrosion-assisted cleavage (CAC) model has been proposed by Flanagan

and Lichter . 1 1 2 ,1 1 3 CAC is based on the interactions between localized anodic

dissolution, localized adsorption and dislocation behavior at the crack tip. It w as shown

that corrosion can enhance local plasticity, which leads to a macroscopically brittle

cracking. This model can be used in both transgranular and intergranular cracking.

In CAC, dislocations generated by mechanical deform ation will pile-up around

the crack tip due to the presence o f Lom er-Cottrell locks near that area. This pile-up

will form an active path for corrosion reactions. This path will be preferentially

corroded away, creating a sharp micro crack, increasing the local stress intensity and

finally initiating a ( 1 10) < 0 0 1 > crack. As the crack grows, the local stress intensity K

drops, and the crack will stop growing when K reaches a critical value. Then, the

corrosion dissolution processes will initiate again, the same process will occur

cyclically, until the material fails. One o f the main deficiencies o f this model is that it

requires the presence o f sessile Lom er-Cottrell locks which are observed only in certain

systems with low stacking fault energy (SFE).

25

S u rfa c e m o b il i ty

C rack g row th

I Metal ion | | Va c a n c yC o n t a m i n a n t

Figure 5. Schematic description o f atom ic processes associated with the surface mobility m odel.1"

26

5. Corrosion Enhanced Plasticity Model

The corrosion enhanced plasticity m odel1 1 4 ,1 1 5 proposed by Magnin, is based on

the principle o f interaction betw een dislocations and corrosion at the crack tip. D ue to

the corrosion interaction with dislocations, there is a large amount o f plastic

deform ation in the micro scale at the crack tip, and the local plasticity is largely

increased. For FCC metals and alloys, this model consists o f the following steps : 116 (a)

depassivation causes a localized anodic dissolution and adsorption on { 1 1 1 } slip planes

at the crack tip; (b) localized dissolution and adsorption result in a stress

concentration, which enhances localized plastic deformation because o f the interaction

among dislocations, adsorption, and localized stresses. The role o f corrosion is

essential but indirect. It enhances the local plasticity at the very crack tip; (c)

dislocations will interact with obstacles, and this will induce the form ation o f pile-ups

where the local stress will increase; (d) if the obstacles are strong enough, local Kic

will be reached. Then an embryo crack will form by a kind o f Stroh mechanism at the

obstacle; (e) the decohesion energy o f { 1 1 1 } microfacets may be decreased by the

adsorption (i.e. hydrogen). The {111} plane will become weaker com pared with other

lattice planes, and cracking will preferentially initiate on this plane. Dislocations are

emitted on asymmetrical planes versus a general crack plane, shielding the new crack

tip. Depending on the crystallographic orientation, this cracking can occur on {111} or

{ 1 0 0 } facets; (f) in the mixed I/II loading mode, this process is expected to lead to

regular changes o f crack planes. A zigzag m icrocracking can occur on {111} and/or

{100} facets. This model has some interesting features but it cannot explain the

formation o f dislocation pile-ups in systems with high SFE such as pure Cu, an element

essential in the model. An advantage o f the corrosion enhanced plasticity model is that

it can be also applied to CF, HE and SCC.

27

6. Hydrogen-Induced Cleavage Model

Hydrogen-assisted cracking is invoked sometimes for BCC materials, but in

conjunction with anodic dissolution has been suggested also for FCC structures. An

interesting proposal has been made recently by Jani et al.,ni based on transmission

electron microscopy (TEM ) studies to characterize the deformation substructure o f 304

stainless steel tested for TG SCC and on earlier w ork by M eletis and H ochm an . 7 The

results showed that the SFE o f the material immediately ahead o f the crack tip is

lowered, with the deform ation m ode at small distances (a few microns) in front o f the

growing crack front being entirely coplanar, while at larger distances hom ogeneous.

Based on these and previous observations o f TGSCC in austenitic stainless steels, a

"hydrogen-induced cleavage" model was proposed with the following sequence o f

events: (a) in the area ahead o f the crack tip, due to the triaxial stress state, the lattice is

expanded and a "hydrogen affected region" (HAR) is formed by the diffusion o f

hydrogen (Figure 6 (a)). (b) Hydrogen in the H A R serves to reduce the SFE within this

volume, therefore restricting cross slip o f dislocations, (c) Planar slip occurs on

intersecting {111} planes that form Lom er-Cottrell locks that are supersessile and lie

on {100} planes, as shown in Figure 6 (b). (d) Since the Lom er-Cottrell barriers are

supersessile, they act as obstacles against glide o f dislocations on the tw o original slip

planes. Therefore, dislocations will pile-up against Lom er-Cottrell barriers, resulting in

extremely high stresses. These stresses will be directed normal to the {100} planes,

and at a critical value, given by the Griffith criterion, the {001} plane ahead o f the

crack will cleave and join up with the pre-existing crack. This is shown schematically

in Figures 6 (c) and 6 (d). (e) After extension, the tip o f the crack will be blunted by

plastic deform ation in front o f it and be arrested. Then the same series o f events will be

repeated.

The new im portant element involved in the "hydrogen-induced cleavage" model

that is supported by the experimental observations is that the corrosion/deform ation

interaction causes a modification in the deform ation mode and the dislocation

configuration at the crack tip. This model requires the presence o f hydrogen for the

m odification o f the deform ation m ode in front o f the crack tip, and indeed, it has been

shown that hydrogen ion form ation and discharge occur readily during TG SCC o f

austenitic steels . 1 1 8 ,1 1 9 How ever, hydrogen involvement has been ruled out from the

TGSCC o f other non-ferrous FCC alloys, such as a-b rass and pure Cu, which basically

exhibit the same cracking phenomenology except that their fracture facets are on {0 1 1 }

planes. Thus, a similar study o f the deformation mode involving these systems would

be o f great interest in establishing the fundamental mechanism o f the embrittlement

process.

E. Corrosion-Deformation Interactions

It is becoming increasingly apparent, that local embrittlement and m icrofracture

can be interpreted in term s o f critical localized corrosion-deform ation interactions

occurring at the near-crack tip region. Both corrosion and deform ation processes can

be greatly enhanced when they occur simultaneously. In the absence o f corrosion,

under slow strain conditions, each element in the metal or alloy undergoes a sequence

o f elastic deform ation, plastic deformation, strain hardening, and finally fracture. In the

presence o f a corrosive environment, the interior elements are unaffected by the

corrosion media, but the exterior elements are attacked (influenced) by the

environment, and becom e "weaker" so the stress required to achieve the same degree

o f strain is significantly reduced. Mom et al.u0 studied the slow strain behavior o f

stainless steels in boiling M gC h and showed that as the strain rate decreases, there is

m ore time for corrosion reactions developing fully, the strain hardening before failure is

gradually diminished by the concurrent effects o f corrosion, but the yield strength is

29

CXACXMOUTH

SUP AND CRACK TIT BIXTNTWO

CSUCXMOUTH

Figure 6. Schematic descriptions o f hydrogen induced cleavage model.

30

unaffected (Figure 7). Because the yielding strength is not affected, the major effect on

the exterior elements by corrosion seems to be a decrease in strain hardening that can

be translated into higher dislocation motion.

Uhlig and co-w orkers 1 2 1 ' 1 2 3 have dem onstrated that anodic surface dissolution

increases the ambient tem perature primary creep rate for both copper and iron. One o f

their examples is shown in Figure 8 , where an applied anodic current clearly increases

the creep rate, which then decreased to the original rate when the applied current was

switched off. Smialowski and Kostanski1 2 4 and Petit and Desjardins125 m ade similar

observations for austenitic stainless steel tested in boiling M gCk solutions. The creep

rate enhancement has been invoked to vacancies produced by the anodic dissolution at

slip steps that migrate into the material and interact with dislocations influencing their

m otion at the near-surface region.

Similar effects have been observed during cyclic straining under a corrosive

environment, with the m orphology o f the C-F fracture surface being different from that

produced in an inert environment. Hahn7 6 ,1 2 6 for example, reported that fatigue surface

shows persistent slip band (PSB) formation early in the 10% cyclic life, while at 90% o f

life cycle the samples show secondary slip bands, which form after significant hardening

has occurred. In the presence o f an anodic corrosion current, the slip bands are

preferentially attacked and the slip distribution at the surface is considerably altered.

On both polycrystals and single crystals o f pure copper, the numbers o f PSB s are

significantly increased and both their heights and breadths are larger.

Besides mechanical effects, corrosion/deform ation interactions can also have

significant effects on corrosion potential and corrosion current, which are the

thermodynamic and kinetic param eters o f the corrosion system, respectively. M eletis et

a l 19 found that amalgam corrosion potential dropped greatly and the corrosion current

increased intensively when sliding-wear was activated on the material's surface

31

O)

2 00

0=glycerine 1=MgQ2, =6 .4 ‘ 1 0 - 5 / s 2=MgC12, =6.4- 10-6 /s 3=MgCl2, =6.4 *10-7/s 4=MgCI2, =6.4 • 10 - 8 / s

100

4 95 6 81 2 3 1070

* • Deformation (mm)

F ig u re 7. Slow strain rate test results on Type 304 stainless steel in boiling M gC h solution . 120

Elo

ngat

ion

32

4 .3 5

4 .3 0

4 .2 5

-P current off

4 .2 0

5

^current on

40 5 0 60 70 8 0 9 0 1 1 0 12 0 100 130 14 0 150

Time, minutes

Figure 8. Effect o f anodic dissolution on creep in pure copper . 121

33

producing tension/com pression stress fields. Jones and co-w orkers 12 7 performed rapid

straining tests and reported a shift in the active potential along with very high anodic

currents. They also noted that a significant increase in anodic current did occur only

when the elastic limit was passed and plastic deform ation had developed. The increase

in the anodic current was attributed to the film rupture by em ergent slip bands.

Parkins 1 2 8 ,1 2 9 studied the interactions in the reversed way. Carbon steel specimens were

stressed and polarized at non-cracking electrochemical potentials in an effort to

suppress the corrosion-deform ation interaction and indeed no cracking was observed.

A considerable volume o f experimental data exists in the literature that can be

interpreted in term s o f corrosion-deform ation interaction that may play a fundamental

role in the embrittlement process. This interaction can mainly be expressed in tw o

forms; as an influence o f the corrosion (environment) process on the dislocation

behavior and as an influence o f plastic deform ation on the electrochemical reactions. A

better understanding o f these interactions would have a lot to offer in shedding more

light on the embrittlement mechanism in an effort to resolve this long standing issue.

F. Stress Corrosion Crack Initiation Stage

The processes preceding crack initiation and propagation should be o f

fundamental importance in understanding the embrittlement mechanism during SCC.

D ue to their transient nature, environm ent-deform ation interactions are expected to

have longer lifetime during the initiation stage, perm itting their easier observation and

study. Thus, the motivation behind the present w ork was first, to study the evolution

o f the deform ation process in the presence o f a SCC environment and identify possible

differences in the absence o f the environment. Second, to assess the role o f the plastic

deform ation patterning on the TGSCC initiation stage, in an effort to shed m ore light

on the crack propagation process.

34

1. Deformation Patterning Under Environmental Influence

There is evidence suggesting that significant differences should be expected in

the deform ation patterns developed in a material that was strained in SCC-free and

SCC-causing environments. For example, K ram er130 reported that a direct correlation

exists betw een the stress at the surface layer and crack propagation rate for com pact

tensile specimens o f Ti-6A1-4V and AISI 4130 steel exposed to H C l-CFbO H and 3

w t% NaCl solution, respectively. It was also shown that different environment/material

combinations will cause different magnitudes o f surface stresses, which is inversely

proportional to the resistance to SCC. A similar observation was reported 131 for

OFHC copper strained in a cupric nitrate-ammonium hydroxide solution that was

known to cause SCC . 1 3 2 The surface layer stress was much higher in the above

solution than in the ammonium hydroxide-copper solution, that causes SCC in brass

but not in copper.

A study o f direct observations o f the dislocation density in SCC specimens was

conducted . 133 The results showed that for 304 stainless steel in a boiling 42 wt%

M gC k solution, the dislocation density in the surface layer increased continuously

under constant stress. On the contrary, samples under the same stress condition but

exposed to an inert atm osphere, showed a slight dislocation density increase at the

beginning, and then remained constant. Yanici134 m easured the dislocation density

change as a function o f distance from the surface o f 304 stainless steel stressed under

constant load in boiling 42 w t% M gC h solution. Initially, the dislocation density in the

surface layer was much higher than that in the interior o f the sample, and both

increased with exposure time. Finally, the tw o dislocation densities became equal and

fracture occurred. However, both dislocation densities were greatly enhanced when

stressing in the SCC environment. Also, both dislocation densities increased very fast

with strain, and cracks occurred when the dislocation density reached a critical value.

K ram er et al.95 studied the relationship betw een dislocation densities and strain

for TG SCC o f naval brass in 0.1 M CuSO t solution, Figure 9. There are several

highlights in their study: (a) the surface layer thickness with a high dislocation density is

about 200 pm and beyond it the dislocation density is constant; (b) cathodically

polarized samples have the same fracture character as the samples stressed in air,

showing a ductile dimple-type fracture; (c) dislocation densities increased linearly with

strain, and the density on the surface is always higher than that in the interior; (d) the

surface layer dislocation density curves do not pass through the origin except for the

interior dislocation density stressed in air. This indicates that once the material is

exposed to the environment, the environm ent-deform ation interaction results in a

sudden dislocation density increase; (e) for naval brass in 0.1 M C uS04 solution, SCC

11 2occurred w hen the surface layer dislocation density reached a value o f 4.5x10 cm' .

This is also the value reached in plastic fracture o f specimens strained in air and under

cathodic polarization; (f) the dislocation density on CF fracture surfaces also has the

same value as that obtained from specimens fractured in a SCC environm ent and in air.

An interesting finding in this report is that by removing periodically the surface layer o f

the SCC specimens, the time before cracking initiation w as extended and the total

strain could reach 30%. When cracking finally occurred, the surface and interior

dislocation densities reached the same values as in the SCC samples. This shows that

crack initiation has a close relationship with the surface layer strain and dislocation

density.

The crystallographic nature o f cracking is another im portant clue to

understanding the initiation and propagation mechanisms o f SCC, especially o f

TGSCC. This topic was reviewed by M eletis7 ,8 for FCC metals and alloys. The results

showed that for TGSCC in austenitic stainless steels, {0 0 1 } is the predom inant

cracking plane with some secondary cracking on {110} planes. The crack propagation

36

7

O l * I ' I I 1 1 I I0 4 8 12 16 20 24

€ %28 32

Sym bol Sym bol Environment Potentialfor data for data (m V (S C E ))obtained obtainedat the in thesurface interior

o , *, A irA , - — A , - - ~ 0.1 M C u S 0 4

0 , ------- « , -------- 0 .1 M C u S 0 4 + 5400 fi 0 .1 M C u S 0 4 + 3 0 0V T 0.1 M C u S 0 4 - 5 4 0

F ig u re 9. Relationship between plastic strain and dislocation density at thenear-surface region and the interior o f the material in naval brass.13

37

process has been shown to be discontinuous1 2 ,1 3 5 and fractographic features on mating

fracture surfaces are found to be matching and interlocking . 8 For non-ferrous metals or

alloys, such as Cu, Cu3 Au, a-brass, etc., the crack initiates on {111} planes, but

propagation occurs on {110} planes rather than lying along {111} slip planes. Lichter,

et al.,m found that the TG fracture surfaces generated during SCC o f ordered and

disordered CU3 -AU are essentially identical to those that form during TG cracking o f a -

brass in ammonia. The fractures appeared to occur by discontinuous cleavage with

{ 1 1 0 }-type cleavage steps, and crack-arrest marks perpendicular to the direction o f

crack grow th w ere present.

2. Effects of Electrolyte on Deformation Patterning

It is generally accepted that SCC occurs easily in certain specific environments.

The chemistry o f the electrolyte in which the material undergoes SCC should be a key

factor to understand the mechanism o f the SCC phenomena. Unfortunately, very little

attention has been given to this very important issue. An understanding o f the effects

o f the electrolyte type on the SCC process can contribute significantly to the

establishment o f the SCC mechanism, but also aid in the practical prevention or

inhibition o f SCC in service.

Althof4 3 7 made an important contribution to understanding the role o f the

ammonia environment in SCC o f a-b rass back in 1944. He dem onstrated that cracking

did not occur until the testing solution turned blue and that the stress corrosion life was

remarkably reduced when the test solutions were pre-saturated with copper ions. It is

known that in aqueous ammonia solutions, copper exists as cupric complex ions o f the

type Cu(NH 3 )n2+, the number o f ammonia ligands, n, varying from 1 to 5 depending on

the ammonia concentration o f the solution . 138 The presence o f these complex ions gives

rise to the blue color o f the solution. Another important variable o f am m onia/a-brass is

the concentration o f ammonia. M ost studies in the past were performed in 15 N

38

aqueous N H 3 , but Pugh et a / .139 showed that experiments conducted in 15 N and 1 N

solutions pre-concentrated with copper produced rapid cracking in both solutions. The

spectrophotom etric studies illustrated that the predominant complex species were

Cu(NH3)52+ and Cu(NH3)42', respectively. It seems that for SCC to occur in the

ammonia/ot-brass system, the concentration o f ammonia can be varied in certain range

as long as the cupric complex ions can be formed.

The effects o f varying the complex ion concentration on the SCC behavior o f

70-30 brass tested in 15 N aqueous ammonia w ere extensively studied.99 Figure 3

illustrates the relationship betw een time-to-failure and the rate o f weight loss for

specimens tested under constant load and copper content o f the solution, which was

directly proportional to the concentration o f Cu(NH3)52+ ions. The tim e-to-failure

decreased w ith increasing concentration o f copper. Also, a well defined inflection

occurred at a critical concentration. Specimens tested in solutions o f copper content

exceeding the critical value were coated with the characteristic black oxide coating

commonly term ed the tarnish, while the specimens tested in solutions o f lower

concentrations w ere apparently free from this coating. The relationship betw een rate o f

weight loss and copper content o f the solution also exhibited a maximum at this critical

concentration.

Thus, it is evident that the electrolyte chemistry is an im portant param eter in

determining the SCC behavior o f a system. Num erous experimental studies and

modeling efforts have been conducted in the past and the present w ork will not make

any attem pt in that direction. Having an overall look o f previous observations on

environm ent/deform ation interactions, it seems that the environment exercises a critical

effect on deform ation and dislocation behavior in particular. It is interesting to note

that this effect is present even in environments that may cause minimal o r no corrosion

at all in the material. Apparently, some type o f interaction may exist betw een the

39

surface layer o f the material as has been modified by the presence o f the electrolyte and

the dislocations present in the near-surface region. A possibility is that an important

effect on aqueous environment/material systems may be exercised by the double layer

structure betw een the electrolyte and the m aterial’s surface (usually metal surface).

Different electrolyte/material combinations will form different double layer structures

and different tendencies to corrosion reactions, which are reflected on the different

corrosion potentials for different material/electrolyte combinations. It has been

discussed before that the SCC phenomena have intrinsic relation with dislocations in

the material and that there is a much higher dislocation density on the crack tip and the

specim en’s surface when SCC occurred.

For clean metallic material surfaces (under vacuum), it was suggested by

Fleischer140 that dislocation nucleation at the surface will be different from that in the

bulk material because the lattice spacing in the surface layer w as different from the bulk

spacing, which is called “relaxation” . This difference will inhibit dislocation nucleation

on the surface layer. Low-energy-electron-diffraction (LEED ) studies141,142 have

shown that “relaxation” o f the lattice occurs in the atomic spacing for several atomic

layers ju st beneath clean surfaces, where the atomic spacing between the first and

second layers is generally contracted compared with the bulk. Occasionally, the surface

lattice even has a different structure from the bulk, which is term ed “reconstruction” .

B oth relaxation and reconstruction are expected to change the dislocation nucleation

and m otion in the near-surface layer o f the material.

Surface relaxation and reconstruction essentially happen because surface atoms

have few er neighbors than those in the bulk. W hen a metal comes in contact with an

electrolyte and the double layer structure is established on the m etal’s surface,

especially in the form o f complex o f the metal ions, the atomic forces acting on the

atom s in the near-surface layer will change significantly. Thus, it should not be

surprising that dislocation nucleation, configuration, and motion may be significantly

influenced by the double layer. This expectation can be relating to the effects o f the

double layer on modifying the image force. A reduction o f the image force by the

particular double layer (producing a low modulus surface layer) can significantly

enhance dislocation motion to the surface and cause localization o f strain. Another

possibility relates to the interaction o f the electrical charges from the double layer with

dislocations and vacancies in the near-surface region o f the material that can also

change dislocation nucleation, configuration, and movement, that is o f primary

im portance to the SCC mechanism.

CHAPTER IV. EXPERIMENTAL PROCEDURES

A. Material

The tes t m aterial selected in the present study w as a -b ra s s due to the large

am ount o f available da ta and extensive previous studies. Tensile specim ens w ere

p repared from a cold rolled a -b ra ss sheet w ith 10 mm gauge length, 5 mm width,

and 0.5 mm thickness, F igure 10 (a). The nom inal com position o f the m aterial w as

72 w t% C u and 28 w t% Zn. Specim ens w ere first stre tched to about 3% and then,

encapsulated in quartz tubes under vacuum and recrystallized by heating at 900° C

fo r 48 hrs to p roduce a final grain size o f about 1 mm. All specim ens w ere ground,

fine polished, and electropolished in 3:1, m ethanol/conc. H N 03 a t -3 0 ° C and 8V

for about 2 min. in o rder to ge t m irror-like, stress-free surfaces.

In addition, sm ooth, cylindrical, single crystal specim ens o f a -b ra s s and

pure copper (99 .999% purity ) w ere used fo r the dislocation configuration studies.

The rationale behind selecting pure Cu in addition to a -b ra ss is th a t Cu has a

relatively high SFE com pared to a -b ra ss and an additional insight can be gained.

Single crystal specim ens w ere used in o rder to avoid com plexities arising from

grain boundary effects easing TEM observations. The gauge lengths o f these

specim ens w ere 25 mm and specim ens w ere 5 mm in diam eter. The single crystals

w ere g row n using the B ridgm an technique. They w ere all e lectropo lished as the

flat a -b ra s s tensile specim ens.

B. Stress Corrosion Cracking Testing

SCC testing was conducted using the slow strain rate technique to study the

evolution o f the deform ation patterning in the a-brass/5 M NHUOH system. The effect

o f the electrolyte type on the deform ation patterning was investigated by conducting

41

42

slow strain ra te tes ts in th ree different electro lytes and determ ining the deform ation

characteristics on the surface o f strained specimens. In addition, electrolyte effects

w ere studied by conducting straining experiments and subsequently determining near­

surface dislocation densities as described later. The three electrolytes selected for the

present w ork were: 5 M N H 4 OH with 8 g/1 Cu added as Cu(N03)2, 0.1 M CuSO t, and

1 M NaNCte solutions. The rationale behind this selection is that, while a -b rass shows

different corrosion behavior in these electrolytes, they all have been found to cause

TG SCC in this material.

The SCC testing included the following slow strain rate experiments: (i) testing

in the particular electrolyte to obtain the strain-to-fracture under SCC conditions, ascc

and (ii) testing for various levels o f straining in the SCC solution to allow determination

o f the deform ation characteristics as a function o f strain. Similar control experiments

w ere conducted in laboratory air and served as standards for comparison. Three repeat

tests w ere conducted for each testing condition.

The slow strain rate tests were carried out using an electrically driven

M echanical Testing System (CORTEST Model 6100). A strain rate o f lx lO '5 s"1 was

applied. It has been established previously by Kramer,95 and by the author, that this

strain rate is suitable for the SCC study o f a-brass. The specimens w ere attached to a

pair o f fiberglass grips in order to allow electrochemical testing during straining. A

plastic cup was positioned and sealed around the gauge length o f the specimens and

filled with the electrolyte. Finally, a salt bridge that was connected to a reference

saturated calomel electrode (SCE) and a pair o f graphite counter electrodes were

positioned in the cell and connected to the electrochemical testing system (EG& G

Potentiostat M odel 273 with M 342C software), Figure 10 (b).

Slow strain rate experiments were also conducted on a-b rass and pure copper

single crystals to study the dislocation substructure during TGSCC. Cylindrical single

II

h

c E

vve

Figure lo o ,

Sel-«P A r sc c ° f(a; "><> ®per^ testing ( a i. , .p rimental cnQ •(■AH dimensions P omen and Ch\"S'0nsare i„ mi„im ™ W eXperimen(a;

44

crystal a -b rass and pure copper specimens were strained to fracture in 5 M NH4OH

and 1 M N aN ( > 2 aqueous solution, respectively. A strain rate o f lxlO "5 s '1 was also

used in these experiments. Part o f the gauge length o f these specimens was covered

with an insulating lacquer to avoid contact with the electrolyte and to restrict

environmental action in a certain section o f the gauge length. This allowed preparation

o f TEM foils from areas w ith and w ithout environmental interaction as described in the

next section.

C. Corrosion Testing

The effect o f the electrolyte type on SCC was studied by performing

electrochemical tests in strained and unstrained specimens. First, the corrosion

behavior o f a -b rass in the three different electrolytes was studied by carrying out

anodic polarization tests. Dissolution rates o f unstrained and strained a-b rass

specimens w ere determined by conducting polarization resistance experiments. The

Tafel slope values needed to calculate the dissolution rates were obtained from the

anodic polarization tests o f unstrained specimens in all three electrolytes. The

polarization resistance tests during straining were conducted to determine the effect o f

mechanical deform ation on the corrosion rate. This m ethod was used for estimating

the corrosion rate in order to minimize applied potential disturbances with corrosion

processes especially for stressed specimens. Second, pure immersion tests w ere also

conducted o f unstressed a -b rass in all three electrolytes to investigate surface

characteristics and possible patterning o f the corrosion process.

All corrosion experiments were conducted by using an EG & G Corrosion

M easurem ent System (M odel 273 with M 342C Corrosion Software). For all

polarization experiments, the specimens were exposed to the solution for 10 min. prior

to testing. The anodic polarization scans started at about 300 mV below the corrosion

potential, Ecorr, and progressed up to 300 mV above Ecorr at a rate o f 1 mV/s. The

polarization resistance tests w ere conducted by polarizing at ±20 mV from the open

45

circuit corrosion potential at a scan rate o f 0.3 mV/s. The tensile specimens used for

polarization resistance tests during straining were masked by using an insulating lacquer

leaving only a certain area o f the gauge length exposed to the electrolyte. Five

polarization tests with reproducible results were performed during straining in each

electrolyte, (Table 1) starting before the elastic limit o f the material and progressing up

to a total strain o f about a = 6%.

Pure immersion tests w ere conducted by placing annealed a -b rass specimens

approximately lx l cm^ in area and about 0.5 mm thickness in the three electrolytes.

The duration o f these tests w as 10 minutes. After testing, the specimens w ere rinsed in

distilled water, dried by com pressed air and immediately examined using scanning

electron m icroscopy (SEM ).

D. Characterization

1. Scanning Electron Microscopy

Scanning electron m icroscopy (ISI 60A and HITACHI S2700) was used in the

present study to characterize the evolution o f deform ation patterning under the

environmental influence. The development o f the deform ation patterning was

characterized by m easuring the slip band spacing (SBS), slip band height (SBH ) and

slip band length (SBL) at various strain levels. Figure 11 schematically illustrates these

param eters in relation to the specimen surface. Slip band height refers to the thickness

(or height) o f the slip band and is directly proportional to the number o f slip planes

involved in the slip band. Crack initiation spacing (CIS), which is the distance between

fine cracks was also used to describe the deformation pattern. D ue to the

polycrystalline nature o f the specimens, measurements were taken in at least three

different grains showing the largest amount o f plastic deformation. Also, in the

interrupted deform ation experiments, markings were lightly scribed on the specimen

surfaces to allow identification o f the same locations for measurements.

slip direction

slip b a n d length

S S S 'v$Cq /} O f

ra.ers

Qtj,

er.

!Pb,Of}

47

SEM was also used to conduct fractographic studies after SCC testing and

examine specimen surfaces after immersion testing in the three different electrolytes.

2. Transmission Electron Microscopy

An analytical TEM (Jeol 100c) was used to characterize the deform ation

substructure in the cylindrical single crystal SCC samples. Thin discs o f the tested

specimens were obtained from the SCC fracture surfaces and sites protected by the

insulating coating (no environmental influence). Electron transparent thin foils were

prepared by electropolishing in a solution containing 75% methanol and 25% H N 0 3 at

a tem perature o f -30° C and voltage o f 10 V. Electropolishing was carried out on a

Tenupol-2 unit (Struers) using both the single-jet and tw o-jet technique. TEM was

performed at 100 keV and electron diffraction was used to identify crystallographic

orientations in the foils.

3. X-Ray Diffraction

The dislocation densities in the near-surface layers o f the a -b rass specimens

produced by straining in the different electrolytes w ere studied by X-ray diffraction

(XRD) techniques. The a -b rass specimens used in the XRD study had an average

grain size o f about 70 pm. These specimens were produced from the originally

stretched specimens (3% ) by annealing at 550 °C for 1 hr. The purpose o f using

sm aller grain size specim ens in the X RD study w as to increase the accu racy o f

th e measurement and reduce the variation among different specimens. XRD in the

present study was conducted by using a Rigacu D-M ax 2B X-ray diffractometer, which

w as equipped with a diffraction beam graphite monochrometer. Copper Ka radiation

w as used with a wavelength o f 1.542 A. An acceleration voltage o f 35 keV and a

current o f 25 mA were used. The full-width half maximum (FW HM ) o f the {111}

diffraction peak was m easured for annealed specimens and specimens strained in air

and in the various electrolytes. The broadening o f FW HM was taken as a measure o f

48

plastic deformation (dislocation density) in the near-surface region o f the specimens to

allow comparison o f the effect o f the electrolyte type on dislocation density.

4. Analysis of Experimental Data

The effects o f electrolyte type on FW HM values were studied by using the

analysis o f variance (ANOVA) m ethod. Paired comparisons w ere conducted to study

the specific differences. The effects o f the electrolyte type on the corrosion rate in the

stress free condition were analyzed by conducting tw o tails heteroscedastic T test. The

effects o f stress on the corrosion rate were investigated by applying general linear

models procedure (GLM ) with the data transfer. The data transform ation was

performed by taking unstressed data as unity. All the analyses w ere conducted by using

statistical analysis system (SAS) program with a 0.95 confidence coefficient.

CHAPTER V. RESULTS

A. Evolution of Deformation Patterning During SCC

Figure 12 shows the variation o f SBS as a function o f plastic strain for

experiments conducted at the same strain rate in laboratory air and in the 5M N H 4 OH

solution. As would have been expected, the SBS for the specimens tested in laboratory

air is larger initially and gradually decreases with strain since the slip band density is

proportional to the magnitude o f the plastic strain. These specimens w ere quite ductile

and showed a strain-to-failure o f about s j = 0.75. The results also show that a t the

initial stages o f deformation, there is a significant difference betw een prominent

(coarse) SBS and average SBS, but eventually this difference is diminished. In

contrast, slip during SCC testing exhibits an entirely different pattern. The SBS for

these experiments is very fine at the initial stages o f the deform ation process and

increases with strain reaching a plateau value o f about 14 pm at fracture. The

approxim ate strain-to-failure for the stress corrosion specimens was about escc = 10%.

SEM examination showed that crack initiation had ju st occurred in the specimens that

w ere strained at a plastic strain o f 8p= 0.002. Also, a strain hardening effect was

observed in the stress/strain curves o f the SCC experiments compared to the

experiments in air.

Figure 13 shows the variation o f the maximum SBH as a function o f plastic

strain. The slip band heights for straining in laboratory air w ere too small (in the order

o f nm) to be measured precisely and w ere therefore theoretically calculated by using a

model developed by Eshelby et al.143 This model has been previously applied for ex-

brass by Yu et al.144 and showed very reasonable agreement with experimental values

obtained by using optical interferometry. Based on the above model, the SSH (h) is the

49

Slip

B

and

Sp

ac

ing

50

20 n a-brass L a b o ra to ry Air □ Average SBS ■ Coarse SBS

SCC E nv ironm en t O C oarse SBS

- 5

gra in

ED

0 -

6 -

4 -

250 5 10 15 20P l a s t i c S t r a i n , %

F ig u re 12. Slip band spacing as a function o f plastic strain in a -b rass tested in 5 M N H 4OH solution plus Cu(N 03)2.

Slip

St

ep

Hei

gh

t,

51

Ot—brass □ L a b o ra to ry Air O SCC E n v iro n m e n t

- 5

1 0 -

- 0.8

0.6

6 -

- 0.4

4 -

- 0.2

0.0

P l a s t i c S t r a i n , %

Figure 13. M aximum slip step height as a function o f plastic strain in a -b ra ss tested in 5 M N H 4OH solution plus Cu(N 03)2.

Slip

St

ep

Hei

gh

t,

52

result o f a pile-up o f n screw dislocations with Burgers vector b gliding out o f the free

surface:

G M ' (6)

where <p is an orientation factor (between 0.71 and 0.75), G is the shear modulus

(about 37 GN/m2 for a-b rass), M = 2.2 is a Taylor average orientation coefficient for

shear in polycrystalline material, k is the slope in the Hall-Petch relationship and I is the

average grain diameter (1 mm). Arm strong et al.145 showed that for 70Cu-30Zn k

changes very little (from 0.34 M N/m 3/2 at yield to 0.374 M N/m 3/2 at 0.2 strain)

producing a relatively insensitive variation o f the SSH with strain in laboratory air. The

SSHs developed during the SCC experiments were substantially larger (by a factor o f

about 25) compared to the step heights in air. Thus, the presence o f the SCC

electrolyte was found to result in environment-induced slip, prom oting deformation

localization in the form o f finely spaced coarse slip bands.

Figures 14 (a)-(c) are scanning electron micrographs from surfaces o f

specimens that were strained at £p = 0.002, 0.017 and 0.10, respectively, in the stress

corrosion electrolyte (5 M NFUOH). It is evident from these micrographs, that crack

initiation occurs along {111} traces at the very early stages o f straining. Such TGSCC

along emergent slip bands on the surface is typical for FCC materials. Also, a fine SBS

develops initially in these experiments with coarse SBH and slip band length. As strain

increases, the deformation becom es coarser in all respects, with increasing SBS, SSH

and SBL. Figure 14 (d) presents a typical slip band m orphology on a specimen surface

that was strained to fracture in air {ej « 0.75). The slip band spacing in specimens

Figure 14. Scanning electron m icrographs showing a-brass specimen surfaceappearance after straining at lx lO '5 s '1 in 5 M N H 4 OH solution plus Cu(NCb ) 2 (a) £p= 0.002, (b) &p- 0.017, (c) sp = 0.10, (d) slip band m orphology in a specimen that was strained to fracture in laboratory air, £/■« 0.75.

54

strained to failure in air was very similar to that m easured in specimens strained at 8p

0.002 in the N H 4 O H environment {d « 1.5 pm, Fig. 12). Nevertheless that morphology

w as quite different with very shallow slip band heights as predicted by Equation (6).

Figure 15 presents an area close to the final fracture surface in a SCC specimen.

Besides the presence o f coarse slip at the near-surface region, the m icrograph shows a

series o f regularly-spaced micro- and m acro-cracks that were initiated along the traces

o f the {111} primary slip planes. M ore importantly, it was observed in the present

w ork that a periodicity exists in the occurrence o f crack initiation (between 6-10 pm,

Fig. 14). A relation also exists betw een SBS and CIS. For example, specimens that

w ere strained at ep - 0.002 exhibited a CIS o f about 6 pm and a SBS o f about 2 pm,

showing that a crack initiation event was activated for every three slip bands. The

transition from initiation to propagation was evident on the SCC specimens, Figure 16,

with the crack initiation site extending for about 50-80 pm. Also, the crack initiation

facets w ere highly crystallographic and featureless as shown in Figure 17.

The potential measurements showed that the corrosion potential o f the

specimens stabilized at about -450 mV (SCE) prior to straining and decreased by about

20-50 mV during straining. The anodic and cathodic Tafel slopes determined from the

anodic polarization experiments were pa = 383±45 mV and fic = 387±32 mV,

respectively. The anodic dissolution rates determined from the polarization

experiments, w ere found to be 5.9 ±0.6 mA/cm2 and 6.7 ±0.05 mA/cm2 for unstressed

and stressed specimens, respectively. It is interesting to note that no statistical

difference was determined in the corrosion rates between elastically (just before the

elastic limit) and plastically stressed specimens. This behavior is in agreement with

earlier slow strain rate results, where strain induced dealloying o f Cu-Zn alloys was

evidenced even with final loads appreciably lower than the 0.1 pet offset yield stress.146

55

F ig u re 15. Scanning electron m icrograph o f an area close to the final stress corrosion fracture surface showing crack initiation patterning (a-b rass tested in 5 M N H 4OH).

56

■ . '■' ■ ■' v*'- i Vu a ■ -■--*,* ■# <**>&'*':■■?■**:•?£. *• ’-• -,

Figure 16. Scanning electron fractograph showing stress corrosion crack y initiation (indicated by arrows) and propagation sites (a-b rass tested in 5 M N H 4OH).

57

F ig u re 17. Scanning electron m icrograph showing flat {111} facets produced during stress corrosion testing (a-b rass tested in 5 M N H 4 OH).

58

B. Effect of the SCC Electrolyte

Table 1 presents the corrosion rates o f stressed and unstressed brass in the three

different electrolytes. The results showed that the corrosion rate o f a-b rass is

significantly different in the three electrolytes. The 5 M N H 4 OH solution was the most

aggressive environment exhibiting the highest corrosion rate, whereas 1 M NaNCh

showed the lowest corrosion rate. In all cases the application o f stress was found to

increase the corrosion rate significantly. The GLM analysis also showed that the

relative increase in the corrosion rate due to stress was similar for N H 4 OH and CuSCh.

It is interesting to note that the greatest increase (almost one order o f magnitude) was

observed for 1 M NaNCh which exhibited the lowest corrosion rate in the experiments

w ithout the application o f stress. Consistent with the above results, the immersion tests

showed that significant corrosion had taken place in the N H 4 OH and CuS04 solutions,

but no evidence o f any attack was found for the NaNCh solution. Observations o f

specimens that w ere exposed in the form er tw o solutions indicated that corrosion in

N H 4 O H exhibited preferential dissolution resulting in a m orphological pattern with

evidence o f periodicity, resembling that o f slip bands developed under stress, Figures

18 (a) and (b). A rather random, corrosion attack was observed on surfaces o f

specimens that were exposed in the CuSC>4 solution, Figure 19. The anodic

polarization curves (Appendix C) showed that a -b rass passivated only in the 1 M

NaNCh solution. Thus, the low corrosion current in the absence o f stress in the above

electrolyte was m ore than likely due to passive film formation. Upon application o f

stress, the film was ruptured exposing fresh metal surface to the electrolyte, thus

increasing the corrosion rate.

Table 2 presents the total strain to crack initiation, total strain-to-fracture and

the SBS at fracture for all three SCC electrolytes and for the experiments in laboratory

air. T otal strain rather than plastic strain w as used due to possible environm ental

59T ab le 1. Tafel Slopes and Corrosion Rate o f a-B rass under Stress and Stress Free

Conditions

SolutionsTafel slope

Pc(mV) Pa(mV)

2C orrosion C urren t(p .A /cm )

Stress free U nder stress

0.1 M CuS04 350±9 276±5 1535±142 1796±61

1 M NaN02 469±67 56±2 7±1 67±1

5 M NH^/Cu** 387132 383±45 5979±592 6777155

Values are based on five measurements.

F ig u re 18. Scanning electron m icrograph showing (a) surface appearance o f aspecimen after 10 m inutes immersion in 5 M N H 4OH plus Cu(N 03)2, and (b) high magnification o f an area shown in (a).

61

¥£&&5w m & ^ $w a S S & s t f M g *

H R i8 ® ^ S t |8 f e s

KiSSSMffiSWEi

1 P ® §

F ig u re 19. Scanning electron m icrograph showing surface appearance o f a specimen after 10 minutes immersion in 0.1 M C u S 0 4 solution.

T ab le 2. Strain-to-Initiation, Strain-to-Fracture and Slip Band Spacing o f a -B rassthe Different Environments.

Environment£ i S f

SBS(nm) at S f

Air — 0.932410.0065 1.310.3

0.1 M CuS04 0.030010.0002 0.803410.0054 1.210.5

1 M NaN02 0.0150+0.0005 0.380910.0065 1.510.4

5 M NH4+/Cu++ 0.005410.0001 0.099010.0032 1.510.4

All values are based on three measurements. £/ = strain for crack initiation £f = strain for fracture

effects on the stress-strain behavior o f the material. Figures 20 (a)-(d) are electron

m icrographs o f typical fracture surfaces o f specimens tested in the three different SCC

solutions and in laboratory air. The values o f strain-to-crack initiation and strain-to-

failure consistently show that the NH4OH environment caused the m ost severe

embrittlement, w hereas the least susceptibility was shown in 0.1 M CuSOt. Table 2

also shows that the SBS at fracture was similar for all three environments as well as for

the experiments in laboratory air. Furtherm ore, Figure 21 provides m ore details on the

SBS developm ent at the initial stages o f straining and Figure 22 gives typical surface

appearances o f slip bands in these environments. It was evident in all o f these tests that

dense SBS o f about 1.5-2 pm had to develop prior to crack initiation. Also, the

presence o f the SCC electrolyte prom otes the development o f dense SBS in

com parison to straining in laboratory air.

Figure 23 presents the early stages o f the stress-strain behavior o f a -b rass in the

three electrolytes and in laboratory air. It is evident that the SCC environm ents cause a

strain hardening and its extent is consistent with the SCC susceptibility shown by the

material in these environments, Table 2. It was noted that for all three electrolytes a

certain degree o f strain hardening was required prior to SCC initiation corresponding

to a stress level o f about 3 .7 x l0 3 psi (25.5 M Pa).

Figure 24 presents the FWFIM measurem ents from XRD tests for specimens

that w ere strained by £t = 5% (total strain) in the three different electrolytes. FW HM

m easurem ents are also included for specimens strained in laboratory air. These results

show that for the same strain level, all three electrolytes produce a higher dislocation

density at the near-surface region compared to that developed in laboratory air.

Com paring the three SCC electrolytes, a much higher dislocation density develops in

NH4OH. In fact, the dislocation density developed by straining at £t = 5% in NH4OH is

64

iftfCi'Jo .l m m

F igu re 20. Fracture surfaces produced by testing (a) in 5 M N H 4OH plus Cu(N 03)2, (b) in 1 M N a N 0 2) (c) in 0.1 M C u S 0 4 and (d) in laboratory air.

SBS,

jx

m65

40

35

30

25

20

5

0

5

00 5 10 15 20 25

s T X 1 0 0 0

F ig u re 21. Slip band spacing developed during SCC testing in the three electrolytes (solid signs represent crack initiation).

66

F ig u re 22. Typical surface appearance o f slip bands developed in the three electrolytes after testing for a total strain o f 0.006 in (a) 5 M N H 4OH with Cu(N 03)2, (b) 1 M N a N 0 2 and (c) 0.1 M C u S 0 4 solution.

67

• y

2000

1000Air-

500000

1500

Figure 23. The early stages o f the stress-strain behavior o f a-brass in the three electrolytes and in laboratory air.

RA/H

M68

Strain Broadening0.60

0.55 envronment

0.50

0.45

0.40

0.35

0.30

0.25

0.200 10 20 30 40 50

% strain

F ig u re 24. The full width half maximum (FW HM ) m easurem ents o f the {111} peak obtained by X-ray diffraction o f a -b rass specimens tested in different environments.

69

equivalent (no statistical difference) to the one developed at fracture by straining in

laboratory air.

C. Dislocation Configuration During TGSCC

Figure 25 (a) illustrates typical dislocation configurations from the near-fracture

surface region o f TGSCC in an a -b rass single crystal. Intensified planar slip and

dislocation pile-ups on {111} planes w ere evident in areas close to the fracture surface.

In many cases the formation o f long stacking fault ribbons was also observed. In

contrast to the dislocation arrangement o f the near-fracture surface area, the dislocation

configuration observed in the pure mechanically deformed regions (masked gauge

section, w ith no environmental contact) was significantly different, Figure 25 (b).

D ispersed dislocations and a much lower dislocation density was observed in these

areas. Some dislocation pile-ups were present, but their number and density (number

o f dislocations/unit area) w ere substantially lower compared to those observed in the

SCC areas.

Several m icrocracks w ere also present on the fracture surface, Figure 26.

E lectron diffraction analysis showed that the microcracks were following a <211 >

direction, which is the orientation o f the intersection between the {111} and the {011}

crack plane. Thus, these m icrocracks m ore than likely are associated with the

form ation process o f crystallographic steps7,12 separating parallel and displayed {011}

TGSCC facets. It is characteristic to note that the walls o f the m icrocracks were very

sm ooth suggesting that they were formed by a highly localized process. During

TGSCC the thin ligaments connecting parallel and displayed cracks sustain most o f the

applied load and thus, have high shear stresses acting upon them. On the other hand,

owing to the presence o f the environment, the material undergoes embrittlement similar

to that observed during the TGSCC initiation process, producing {111} cracking.

Figure 25. Dislocation configuration in an a-b rass SCC specimen (a) a region near the stress corrosion fracture surface showing intensified pile-ups on {111} slip planes and (b) a pure mechanically deformed region showing a more hom ogeneous dislocation distribution.

71

5 t w

0.2 (Mm

Figure 26. A microcrack in a -b rass following the <211> direction (intersection o f the {111} with the {011} crack plane).

In pure Cu, dislocation configurations in the TGSCC and pure mechanically

deform ed regions were, in general, o f a cellular structure, which is consistent with the

2 2 higher SFE o f this metal («90 mJ/m for pure Cu compared to «10 mJ\m for a-brass).

How ever, in the pure mechanically deformed areas, the dislocation density was lower,

the distribution was m ore diffused, Figure 27 (b), and indicated a rather hom ogeneous

deform ation mode. In areas close to the SCC surface, traces o f planar slip w ere found

within the dislocation cell walls, with the dislocations being m ore intensified, Figure 27

(a). Furtherm ore, m icrocracks with < 211> traces were also present on the stress

corrosion fracture surfaces. Similar to a-brass, relatively sm ooth m icrocrack walls

w ere also observed, Figure 28. Intense coplanar arrays o f dislocations w ere evident in

regions ju st ahead o f the tip o f such micro-cracks, Figure 28. The higher dislocation

density present on the m icrocrack walls in pure copper, Figure 28, com pared to that in

a-brass, Figure 26, is consistent with the significantly higher SFE o f copper. In

general, the facets o f the TG SCC steps in copper exhibit significant plastic deform ation

and a much less crystallographic character compared to those in a -b rass.95

73

0.2 f i m

ytjauMaktfftusfoiffjyiiBiaMi

Figure 27. Dislocation configuration in a pure Cu SCC specimen (a) a region near the stress corrosion fracture surface showing cellular dislocations with traces o f planar slip in the walls o f the dislocation cells and (b) a pure mechanically deformed region showing a m ore hom ogeneous dislocation substructure.

74

4, w - . y - « v y ^ n y ;y-: V ;> v T - f Y f y ' - a ' w , - y for: p, ;<■*P :%

Figure 28. Transmission electron m icrograph showing dislocation pile-ups associated with dissolution m icrocracks in pure copper. The m icrocrack tip is shown on the upper part o f the micrograph.

CHAPTER VI. DISCUSSION

A. Deformation Evolution During TGSCC Initiation

1. Environment-Induced Surface Plasticity

In view o f the localized nature o f the deformation evolved during SCC in the

present w ork, the magnitude o f the plastic strain ( £s) was estim ated at the near-surface

region. Strain due to coarse slip can be calculated from the SBH (h) and the num ber o f

slip bands per unit length (N) by using the expression.147

w here N = 1/d w here d is the SBS. It should be noted that the strain calculated from

the above equation is independent o f the grain size o f the material. Table 3 provides

the calculated values o f near-surface strain as a function o f the bulk strain, SBH and

initiation ( £ / ) is equivalent to the strain-to-fracture required to produce ductile fracture

o f the material in air ( £ / = £ /= 0.75). This is further verified by the XRD

m easurem ents shown in Figure 24, where very similar dislocation densities were

obtained for specimens strained by only £t =5% in NH4OH (crack initiation had

occurred) and strained to fracture in laboratory air, £t = 49% . Thus, a fundamental

fracture criterion apparently exists (critical strain o r dislocation density for fracture) for

these ductile FCC materials, and environmentally-induced cracking is caused by

localizing the required strain for fracture. The reason that the near-surface strain starts

decreasing after some point, is that after crack initiation, m ost o f the strain is localized

(7)

slip band density. It is striking that the localized near-surface strain required for crack

75

T ab le 3. Calculated Values o f N ear-Surface Region Strain.76

Bulk Straine

Slip Step Height h, pm

Slip Band Density N , |im"l

Near-SurfaceStrain

sc0.002 1.4 0.66 0.770.006 3.3 0.28 0.770.017 6.0 0.15 0.760.036 8.0 0.10 0.660.100 10.0 0.07 0.58

All values are based on three measurements.

77

at the crack tip and the SBH may also be reduced due to dissolution at the slip steps.

Thus, after initiation the coarsening rate o f SBH is reduced (reaching a plateau around

14 pm, Figures 12 and 13) and neighboring slip bands coalescence producing larger

SBS at the near-surface region as shown in Figure 14.

Furtherm ore, the above values o f near-surface strain can be used to obtain a

rough estimate o f the number o f dislocations involved in the slip bands and the

dislocation density at the near-surface region by using the expression.148

ss = Nnbcp (8)

For ss = 0.77 (corresponding bulk strain &p = 0.002) the above equation produces n =

3 • 1 1 2 *6x10 dislocations and a dislocation density o f about 4x10 c m '. It is im portant to

note at this point that the above calculated dislocation density is in a complete

agreement with that determined independently by K ram er et al.95 (4 .5x10 11 cm’2) for

naval brass tested in CuS04 by using X-ray diffraction methods, in spite o f differences

in alloy com position and type o f electrolyte. In fact, several previous SCC studies149' 151

have shown that a high critical dislocation density develops in the near-surface region

prior to cracking. Similar effects have been observed in materials undergoing fatigue

cracking.13,14,152 Thus, in relation to the aforementioned fracture criterion, it can be

postulated that in general, fracture occurs when the total strain energy from the

combined stress field o f dislocations in the material reaches a critical limit relating to

the lattice decohesion. In SCC the critical dislocation density is localized and in the

ductile overload fracture distributed within the entire volume o f the material.

2. Effect of Anodic Dissolution

In an effort to assess the possible contribution o f the anodic dissolution to the

crack initiation process, an estimate o f the crack initiation rate (da/di) w as obtained.

The estimated rate was com pared with the experimentally observed rate o f about

78

0.2 pm/s. Assuming that the first crack initiation event is due to anodic dissolution at

the slip bands, then

- = i — (9)dt zF p '

where i£ is the current density at the slip bands, M is the atomic weight o f the metal

(about 65 g/mole), 2 is the valence o f the solvated species ( 2 = 2), F is the Faraday's

constant (96500 C/mole) and p the density o f the material (7.13 g/cm 3 assuming Zn

dissolution). Even by assuming that the current density i6 is equivalent to the total

dissolution rate m easured during straining (6 .7 x l0 3 pA /cm 2), da/dt « 0 .3 x l0 '2 pm/s,

which is almost tw o orders o f magnitude lower than the experimentally observed crack

initiation rate for the present strain rate is obtained. The above analysis suggests that

indeed enhanced dissolution takes place at the emergent slip bands, but m ore than likely

dissolution alone cannot account for the crack initiation.

B. Effect of Electrolyte Type on TGSCC

The corrosion results from the three different electrolytes further foster the

above indication that crack initiation cannot be attributed to anodic dissolution alone.

The corrosion rate o f a -b rass in NaNCh was found to be by m ore than tw o and a half

orders o f m agnitude lower than the one in CuS04 (Table 1), however, the SCC

susceptibility in the form er solution was significantly higher (Table 2). Nevertheless,

the corrosion results clearly show that the application o f stresses can significantly

enhance anodic dissolution, especially in cases where passive films develop (Table 1).

The present results suggest that the role o f the environment is to cause strain

hardening. Figures 21, 23 and 24 show that SCC susceptibility is directly related to the

ability o f the environment to cause strain hardening and is consistent with the

79

aforementioned fracture criterion o f critical strain energy. Strain hardening is

translated into dense SBS (Figure 21) and higher slope in the stress-strain response o f

the material (Figure 23) due to the developm ent o f higher dislocation density (Figure

24). Thus, in term s o f SCC susceptibility the three environments can be ranked as

N H 40H > NaNCte > CuS04.

The present results suggest that the ability o f the environment to cause SCC

initiation is related to its effectiveness in producing dense slip. For example, among the

three SCC electrolytes, C uS04 exhibits the slowest rate o f slip band density increase

(Figure 21) and therefore the lowest SCC susceptibility. In view o f the present results,

the role o f the environment in increasing slip band density can be attributed not to the

corrosion rate itself, but probably to the corrosion instabilities and patterning in the

dissolution process.

The dissolution o f alloys often occurs via the development and m otion o f a

rugged reaction front owing to the selective dissolution o f one or more elements. This

process is known as de-alloying and in the case o f a -b rass involves mainly Zn

dissolution. Early studies by W agner153 showed that, i f the rate-limiting step for such a

process is in the phase being consumed, geometrical instabilities develop along a planar

interface and grow exponentially with time. A similar situation was observed in

aggregation processes154 and this has been an area o f intense study.155' 157 A global

theoretical approach to the morphological instability phenomenon has been developed

by Santarini158 and application o f this model to dissolution 159 predicts the development

o f patterning in the mobile interface.

Some evidence o f dissolution patterning was observed in the present study

during the immersion tests in N H 4 OH, Figure 18, and that probably can account for the

fast development o f dense slip band spacing in this environment under stress.

Immersion in CuS04 was found to cause a rather random dissolution, Figure 19, that

80

probably delayed the development o f dense slip spacing. Finally, exposure to NaNCte

w as found to cause passivation but under stress the film is ruptured selectively at

em ergent slip bands at the surface thus localizing dissolution and inducing a pattern that

consequently eases slip band development, Figure 22 (b).

C. Dislocation Arrangement During SCC

The present w ork provides further evidence that there is a difference in the

deform ation m ode betw een the areas directly in front o f the crack tip where an

environmental interaction exists and areas at larger distances (with no environmental

interaction). The deformation in both the non-ferrous FCC metals tested in the present

w ork w as found to be m ore localized (coplanar) in the TG stress corrosion crack tip

region and rather hom ogeneous in areas away from the grow ing crack, which is in

agreem ent with the previous TEM results in 304 austenitic stainless steel.117,160 In

principle, the above TEM observations are consistent with earlier studies95,151'161,152 and

with the results o f the present deform ation evolution study, where high dislocation

densities w ere observed in the near-surface region resulting from the environmental

action. It should be pointed out, that besides the higher dislocation density observed in

the vicinity o f the crack tip, the present results also indicate that the "stress-

environment interaction" modifies the dislocation configuration.

The results o f the present deform ation evolution study showed that slip is highly

localized, but on the other hand has a very coarse nature, too. The slip bands were

very sm ooth in appearance resembling that o f cleavage-like TG SCC surfaces. In

addition, from TEM observations during in situ crack tip deform ation o f FCC metals,

O hr163 reported that in the early stages o f straining, deform ation is coplanar, and a

transition to cellular arrays occurs only at larger strains. Thus, under normal

conditions, the area immediately ahead o f the stress corrosion crack tip, where the

strains are the largest, would be expected to display less dislocation planarity than areas

81

further away from the crack tip where the strains are relatively low. However, the

opposite was observed in the present study, and thus, the present evidence is difficult to

reconcile with localized dissolution proposals.5

D. Crack Initiation and Propagation

The present w ork shows that besides the enhanced dislocation activity in the

near-surface region, the environment also modifies the deform ation pattern. The

environment-induced deform ation is characterized by finely-spaced, coarse slip bands,

in agreement with the TEM observations from areas just in front o f the crack tip in

three different FCC materials, namely 304 austenitic stainless steel117, a-brass, and pure

copper.164,165 In all o f the above studies it was observed that slip in front o f the crack

tip was localized (closely spaced slip bands) and m ore coplanar (coarse slip bands).

This evidence suggests that m ore than likely the latter deformation pattern develops

prior to each crack propagation event. Thus, since under the environmental action

planar dislocation arrangements are produced, the cracking process can be attributed to

the opening o f the pile-up when a critical deform ation localization (critical strain

energy) is reached. It has been shown in earlier treatm ents, that substantial local

stresses can be developed when a sufficient number o f dislocations join a pile-up.166

Thus, the normal stresses necessary to open the pile-up can be provided from the sum

o f the externally applied stress and the local stress generated from the dislocations in

the pile-up. Also, the possibility o f highly localized dissolution assisting the mechanical

opening should not be excluded.

D ue to its localized nature, this fracture process would be expected to produce

very flat (111} facets consistent with the present experimental evidence, Figure 17.

Also, the same mechanism can produce the crystallographic steps occurring on

alternating {111} segments and separating parallel but displaced {011} propagation

facets.7,8'12'167 There is no difference in the appearance between the presently obseived

82

{111} crack initiation facets and the {111} step facets. The above reasoning seems to

provide also a plausible explanation for the observed patterning in the CIS, Figure 15.

The crack initiation patterning is developed since the combined stress field from a

num ber o f pile-ups on surrounding slip bands (one crack initiated for every three slip

bands in the present study) contribute to reach locally the critical stress required for

fracture.

The reason the crack initiation plane is different from the crack propagation

plane is probably that during the very first initiation event, slip occurs first on the

primary slip plane thus the environment-induced deform ation is pronounced on that

plane resulting in {111} crack initiation. As deform ation progresses, double (and

multiple) slip takes place and the environmental effect is occurring on all slip planes

modifying the local stress field and resulting in a new crack orientation on the {011}

plane. In the presence o f a crack obviously no {111} crack initiation is observed since

the crack tip stress intensity is high enough to produce double o r multiple slip as has

been experimentally verified.7,117 Slip in the latter case is not required to be as coarse

as in the first initiation event since a much higher stress intensity prevails. Similarly, no

{111} crack nucleation should be expected in cases where significant pre-straining

exists.

Thus, there seems to be no fundamental difference in the crack initiation and

crack propagation mechanisms but simply the differences in crack orientation and

appearance may be due to the pronounced slip on the primary slip plane during the very

first initiation event. This can explain the fact that {111} crack initiation facets are very

sm ooth and featureless (Figure 17), where crack propagation facets show evidence o f

em ergent slip traces.7,117,114,16/ However, {111} crack propagation may be possible in

materials where certain slip systems are blocked and gliding is allowed mainly on a

single slip system (i.e. martensite formation and preferential precipitation on certain

83

{111} segments in austenitic stainless steels and A1 alloys, respectively), o r in materials

where the environment is prom oting enhanced gliding on the primary slip system that is

sufficient to cause fast crack grow th prior to the activation o f additional slip

system s.168,169 Regarding the crack advance distance, it is only logical to expect that

the crack will propagate in steps about equal to the depth o f the surface layer (o r the

region in front o f the crack tip) w ith high dislocation density (for TG SCC in P-brass it

has been determined to be 20-30 pm 151). The brittle character o f the fracture surface

can be accounted for by considering the coarse character o f deform ation within its

localization (Figure 14) that on a microscale is causing very little distortion in the

adjacent atomic structure o f the material (i.e. localized coarse slip on {111} planes).

In principle, the enhanced dislocation mobility observed during SCC in the

present and previous studies is similar to liquid metal embrittlem ent170 and hydrogen-

induced micro-plasticity.26,171 Coarse slip band formation during slow rate straining o f

Al-Li-Cu alloys that were pre-exposed in NaCl solution, a stage that presumably

produces hydrogen that enters the material has also been observed.172 These

observations tend to suggest that in all o f the above environm ent-assisted cracking

phenomena the role o f the environm ent is to reduce the activation energy for

dislocation nucleation and m otion, and also to coarsen the slip pattern as observed in

the present work.

E. Proposed Environment-Induced Deformation Localization Mechanism

Even though several TG SCC m odels10,110,114,173,174 have been proposed in the

past, no consensus has been reached yet regarding the nature o f the cracking process.

In view o f the present evidence, an environment-induced deformation localization

mechanism based on a vacancy-dislocation interaction process is being proposed.

The main elem ent o f the p resen t m echanism is that the environm ent induces

84

nucleation o f dislocations from the surface region at loads well below those required

for normal yielding in air. This process is shown schematically in Figure 29. Vacancies

produced by anodic dissolution reduce the interatomic bond density at the surface

region, causing dislocation nucleation and out o f surface displacement. It has been

reported that extensive brass dealloying occurs on stressed surfaces175’176 and recently

m ore direct evidence o f anodic generation o f vacancies has been provided.177 The

present results further suggest that a periodicity in the dissolution process may be

critical in enhancing the development o f dense slip band spacing at the near-surface

region.

Very recently, the critical role o f the surface state in the dislocation nucleation

process was dem onstrated by Gerberich et al. 178 by conducting nanoindentation

studies. It w as shown that the load bearing capacity o f Fe-3wt% Si can be reduced by

tw o orders o f magnitude if the 10 nm thick native„oxide film is removed. Thus,

dislocation generation occurred at loads well below the normal plastic loading.

Similarly, the role o f the SCC environment can be thought o f as one o f creating a

surface state (corrosion patterning and periodicity) that prom otes dislocation emission

at loads below the elastic limit. In this case, vacancies reduce the interatom ic bond

density at the near-surface region facilitating nucleation o f dislocations, Figure 29. As

the process o f environment-induced dislocation emission continues, the local strain

from the accumulated dislocations reaches a critical level as described earlier and (111)

cracking occurs. It is important to point out that due to lower energy barrier for

dislocation nucleation during SCC, strain hardening (dislocation accum ulation and

interaction) is accelerated as was observed experimentally, Figures 23 and 24. After

initiation, the dislocation emission process is facilitated on more than one slip system

and the stress field is modified causing '{011} overall cracking, Figure 30.

85

ooooooooooooooooooooooooooooooooooo

oooooooooooooooooooooooooooooooooo(b )

ooooooooooooooOOOOOtOoooooooooooooo

(<=)

ooooooooooooooOOOOt Ooooooooooooooo

(d)

oooooooooooooooo

oooooooo(«) ( 0

F ig u re 29. Schematic representation o f proposed environment-induced dislocation emission (a) perfect lattice; (b) vacancy generation due to anodic dissolution; (c) generation o f a dislocation and out o f surface displacement by one inter atomic distance;(d) and (e) formation o f planar dislocation arrangement; and (f) crack nucleation by opening o f a dislocation pile-up in a slip band.

86

overa ll c rack ing p lane

d is locationnucleation

F ig u re 30. Schematic representation o f proposed multiple dislocation nucleation and fracture surface characteristics produced during crack propagation.

87

Besides the induction o f dislocations from the material surface, the possibility

exists that vacancies can further interact with dislocations and also stress-assisted

vacancy diffusion may influence the propagation process. These tw o aspects are

further explored in the next sections.

1. Vacancy-Dislocation Interaction

Additional vacancy-dislocation interactions were explored in an effort to

provide a plausible explanation for the observed localization in the deform ation mode

during SCC. It takes into consideration earlier ideas expressed by Forty179 and

Jones.174 It is advocated that besides vacancy induction o f dislocations during the

crack initiation stage, vacancies produced by the anodic dissolution process at

emergent slip steps,5 migrate in the material and interact with dislocations in the near­

crack tip region, modifying their configuration. The interaction o f vacancies with

dislocations has been previously discussed for both, the "good" lattice surrounding the

co re180 and the narrow core region consisting o f a highly distorted atomic structure.181

This subject is further reviewed by Ballufi et a l .1S2 Vacancies do not have much elastic

interaction with dislocations, but the electrical and other effects associated with the

absence o f an atom make up for this in most solids. Considering only the effect o f the

variations in charge density (z-z0)e that result from the volume dilatation, Cottrell et

al.m derived the following numerical expression for the electrical interaction energy

( Uelectr) between an edge dislocation (with a Burgers vector b) and a vacancy at

distance r :

U e le c tr * 0.02(z-zo)b/r, eV (10)

Positive vacancy binding energies ranging from 0.2 eV - 0.7 eV have been found for

both o f the above dislocation regions. Furthermore, positive vacancy binding energies

have been determined in both the compressive and tension regions o f an edge

dislocation, indicating that vacancies can easily migrate tow ards the core region.

Thus, the initial vacancy injection at the surface region o f the crack tip is

visualized to be followed by vacancy attraction to the crack tip dislocations. Vacancies

are absorbed and eliminated on dislocation jogs causing dislocation climbing and

gliding around barriers that otherwise would impede dislocation m otion191 resulting in

faster dislocation accumulation thus reducing the time for strain hardening. The

interaction o f vacancies with dislocations prom oting glide is consistent with the

emergent slip bands consisting o f a cluster o f slip lines observed in the present w ork on

the specimen surfaces and the localized dissolution taking place at these sites that serve

as source for vacancies. Intense corrosion at emergent slip steps was observed for both

a -b rass and copper specimens in the present study.

It is important to note that the vacancy generation mechanism during corrosion

requires form ation o f kink and ledge sites at the crystal surface184 which are provided

by the slip process. As has been discussed earlier in the present study, corrosion

patterning and periodicity effects can produce finely spaced slip bands which

consequently can furnish the required ledges for vacancy formation. It has been

docum ented that intense slip bands are very prone to preferential dissolution due to the

influence o f the dislocation m icrostructure on the free enthalpy o f dissolution and the

energy o f activation.185 M ore importantly, vacancy formation and their adsorption by

dislocations is expected to take place very effectively since the corrosion process is

occurring on the slip bands. Thus, the environm ent/deform ation interaction involves

localization o f deformation by the environment and localization o f corrosion by slip.

There is further evidence for the above interaction from earlier w ork on Al, Cu

and Au single crystals by K ram er186 who observed lower activation energies for plastic

deform ation in tests under metal dissolution conditions. In fact, the decrease in the

89

activation energy was consistent with the dissolution rate o f the metals and followed

the order Al > Cu > Au. Very recently, Jones and Jankow ski177 provided conclusive

evidence to an earlier proposal by Revie and Uhlig121 that anodic dissolution can

generate a supersaturation o f subsurface vacancies which form divacancies and migrate

in the material. These observations are consistent with the notion that corrosion-

induced vacancies prom ote plastic deformation. Thus, there is evidence suggesting

form ation o f corrosion-induced vacancies in FCC pure metals and alloys e.g. dealloying

o f brasses (dezincification o f SCC surfaces).177 In summary, the possible patterning in

metal dissolution seems to facilitate plastic deformation, which in turn favors formation

o f emergent slip bands that further enhance localized dissolution.

W hile the near-surface dislocation density increases, the vacancies that have

interacted with dislocations tend to prom ote a coplanar dislocation arrangement. This

can be envisioned to occur by a mechanism involving annihilation o f dislocation-

dislocation intersections and dislocation jogs by vacancies, since these sites can act as

sinks for absorbed vacancies. Furthermore, continuous supply o f vacancies to these

sites by dislocations (and diffusion) is expected due to the high subsurface vacancy

concentration produced by the corrosion process. Thus, the vacancy-induced

dislocation process, the migration o f divacancies and their interaction with dislocations

can account for the observed high dislocation density at the surface174’180 and the

localized deform ation mode observed in the present work. The sequence o f events

seems to be as follows: localization o f deformation under the environmental influence

(vacancy-induced dislocations due to corrosion patterning), localized dissolution at

emergent slip bands at the surface, vacancy generation which react to form divacancies 121

that migrate into the material, divacancy attraction by dislocations, easier dislocation

glide and build up o f high dislocation density, interaction o f divacancies with energetic

sites (dislocation intersections) and modification o f the dislocation configuration.

90

2. Vacancy Contribution to Crack Propagation

It is pertinent at this point to refer to an early suggestion by Forty179 regarding a

TG SCC mechanism for a -b rass involving the interaction o f dislocations with a thin

elongated void on {111} produced by vacancies generated from dealloying. Forty also

considered a stress analysis by Fujita187 showing that if a thin void exists on the {111}

slip plane, then slip on this plane can force dislocations to join the slit and form a

hollow dislocation with a large Burgers vector (nh). Stress field analysis around this

hollow dislocation indicated that a substantial com ponent o f normal stress is acting

across the slip plane, that can lead to fracture when a sufficient num ber o f dislocations

(ri) jo in the hollow dislocation or when

nb = (Lh/2)1/2 (11)

L and h are the length and width o f the void, respectively. A similar expression was

developed earlier by S troh166 by considering a dislocation pile-up on a Lom er-Cottrell

L ock (LCL). U nder an effective shear stress (oa-oo) the maximum normal stress (om)

at a distance r from the lock is:

om = (ocr o0) (nb/r)1/2 (12)

Oa and oo being the applied shear stress and frictional stress, respectively. This concept

was adopted later by Kramer et al.95 who suggested that TGSCC occurs when the sum

o f the applied stress and the stress due to the accumulation o f dislocations reach the

fracture strength o f the material. M ore interestingly, the above stress field analysis191,192

showed that the direction o f maximum normal stress is not vertical to the slip plane but

it forms an angle o f about 70° with this plane. The implication o f this stress field

91

geom etry is that after a certain length o f a nucleated crack, fracture may tend to extend

on a plane inclined to the {111} crack initiation plane, i.e. {011} plane.

In view o f the above discussion, we can further elaborate on the proposed

environment-induced deformation localization mechanism. Firstly, crack initiation

occurs by environmentally-induced strain localization on {111}, when a critical strain

energy is reached creating a fine slot on that plane. After crack initiation on {111}

plane, strain localization continues to occur in front o f the {111} slot and dislocations

start piling up. A gradually increasing normal stress develops in the surrounding region

as m ore dislocations join the pile up, with its maximum located along a direction

inclined at 70° from the {111}. At a certain point the stress intensity at the tip o f the

slot is high enough to activate the secondary slip plane and dislocations start piling up

on that {111} plane at 70.5° from the primary slip plane. The activation o f the

secondary slip system will modify the stress field, and under these conditions, the final

direction o f the stress (from both pile-ups) is approaching the direction normal to

{011} since this plane bisects the tw o slip planes. Thus, under the environment-

deform ation interaction localized slip and dissolution can occur simultaneously or in a

fast alternating sequence on both slip planes producing cleavage with an overall {011}

plane orientation, Figure 30, when a critical value o f strain energy is reached. Since the

m agnitude o f the critical tensile stress depends on the num ber o f dislocations in the pile

up, a direct relation between tendency to planar slip o r SFE and TGSCC susceptibility

is implied. Furthermore, tendency to planar slip will also correlate to the intensity o f

dissolution at slip steps and vacancy formation, since the latter mechanism requires

surface ledges184 that can be generated by the em ergent slip steps.

Cleavage on {011} is prom oted due to its particular crystallographic orientation

betw een the tw o {111} planes, low surface energy95 or to accum ulation o f vacancies

on this plane as suggested by Jones.177 It should be noted however, that the crack jump

92

distances predicted by Jones could not comfortably be reconciled w ith the length o f

observed crack advance distances. However, in view o f the present results, it is evident

that a higher vacancy concentration may be produced in the plastic zone due to

enhanced vacancy transport along dislocations. Indeed, high rates o f vacancy diffusion

along dislocation cores can be reached due to the substantial reductions o f the vacancy

m igration energy in the highly disturbed core region.182 Thus, vacancies may also

contribute directly to the crack propagation process.

In an effort to explore further the possible contribution o f vacancies to the

embrittlement process, the divacancy diffusion in the high triaxiality region in front o f a

blunted crack tip is considered. Divacancies instead o f vacancies w ere considered due

to their higher stability and higher diffusivity.121 The basic argum ent used is, in

principle, similar to that proposed by N abarro for creep.188 In the presence o f an

applied tensile stress o, the free energy o f vacancy form ation is reduced by an amount

o fl (Q is the atom ic volume) producing a higher equilibrium vacancy concentration C :

C = C o e x p (o Q fk T ) (13)

C o being the thermal equilibrium concentration o f vacancies in the stress free region; T

the tem perature; and k the Boltzm ann’s constant. Thus, under stress, a net flow o f

vacancies can be realized due to the high concentration gradient betw een the vacancy

source (subsurface) and the highly stressed regime just in front o f the crack tip

(triaxiality region). The process followed here is very similar to the one used recently

to describe the stress-assisted grain boundary diffusion o f atomic hydrogen in Al-Li

alloys.189 First, Hill's model was adopted to describe the stress field ahead o f a blunted

crack tip, Figure 31(a) (see Appendix B for appropriate equations and values used).

93

Then, Van Leeuwen's equation was modified to describe the diffusion o f divacancies in

the presence o f a stress field as follows:

dt v cPxd a dCv

R T dx dx (14)

v denoting divacancies and Vv the equivalent partial molar volume o f divacancies. By

assuming typical values for a-brass, such as a waiting time for crack advancement o f

60 s (experimentally observed SCC velocity «10 '7 m/s and crack jum p distance o f 6 p

m), the transient concentration profile o f divacancies was determined by solving

equation 4 numerically with Finite Difference. Figure 31(b) shows the transient

concentration profiles for stress-free (600 s) and stress-assisted (60 s and 600 s)

diffusion. For the stress-free case only the transient profile for 600 s is shown, since

the profile for 60 s was very shallow (penetration distance «2000 A) and for practical

purposes is not shown on the graph. The analysis shows that the presence o f the stress

field prom otes significantly the divacancy diffusion in the plastic zone. For a typical

waiting time o f 60 s, the divacancy diffusion zone extends almost 2 pm in front o f the

crack tip. Furtherm ore, based on the present analysis an additional beneficial effect on

divacancy m igration should be anticipated from the local stress field generated by the

dislocation pile-ups on the crack initiation slot.

It is interesting to note that recent creep experiments with planar copper-

sapphire interfaces tested in tension,190 showed the existence o f crystallographic

cavities at the interface, with cavity surfaces close to {110}. These observations

indicate that under stress, vacancies may have a tendency to accum ulate on specific

planes. The reason for that preference is not known at the moment, but it may relate to

the low atomic density o f {011}, its low surface energy and/or the particular geometry

94

a800 i

0 0 4 0 0 -

200 -Elast i cZo n e

P las t i cZ on e

2 5155 10 200x, pm

bElast i cZ on e— Plas t i c

Zone

T r a n s i en t Prof i le6 0 s 6 0 0 s

S t r e s s F r e eCJ

x

S t r e s s As s i s t e d0 . 5

0.01050

x, pjm

F ig u re 31. (a) Stress distribution in front o f a blunted crack in a-brass; (b) Transient divacancy concentration profiles o f stressfree (600 s) and stress-assisted (60 s and 600 s) diffusion (Cs denotes the divacancy concentration at the su rface ).

95

o f deformation. The {110} plane bisects the tw o {111} slip planes and vacancy

accumulation on {O il} , while these tw o planes are active, may cause maximum strain

relaxation. I f this is indeed the case, and under stress divacancies tend to accum ulate on

{110} planes in the diffusion zone, then embrittlement can be easily induced due to

highly localized alternating slip on this plane and due to the significant reduction o f

local fracture toughness in the embrittled zone. Also, the crack velocity in such an

"embrittled zone" is expected to be high enough for the crack to propagate for a

distance longer than the divacancy diffusion zone (m ore than likely up to the size o f the

plastic zone, 6 pm). Due to the stress triaxiality in this region, significant amount o f

plastic deform ation should not occur once a crack has initiated. Crack arrest can occur

at a preexisting slip band and the process is repeated. It should be noted that after the

first crack initiation event, less anodic dissolution will be required at slip bands because

as the crack gradually advances the stress intensity is increasing in front o f the crack tip

thus, a smaller number o f dislocations will be required in the pile-ups to reach the

critical stress for cracking. Finally, it should be pointed out that the environment-

induced deformation localization TGSCC mechanism described in the present work, is

compatible with the phenomenology o f cracking in non-ferrous FCC metals.

CHAPTER VII. CONCLUSIONS

The results o f the present study showed that an environment-induced

deform ation localization pattern evolves at the near-surface region o f the material prior

to TGSCC. The near-surface strain required for SCC w as found to be equivalent to

strain needed for ductile fracture o f the material in air indicating the existence o f a

fundamental fracture criterion. Also, a patterning in the crack initiation process was

observed. In view o f the present results, TG SCC can be interpreted as localized

fracture under the sum o f the externally applied stress and the stress resulting from the

neighboring dislocation pile-ups lying on different slip planes that are accumulated in

the surrounding slip bands. The role o f the environment can be thought o f as that o f

causing a local reduction in the energy barrier for dislocation nucleation, thus causing

faster strain hardening com pared to behavior in laboratory air. This environment-

induced dislocation process can be attributed to the presence o f vacancies that reduce

the density o f the interatom ic bonds at the near-surface region, facilitating dislocation

emission.

Based on the present experimental observations and the phenom enology o f

TG SCC in non-ferrous FCC metals, an environment-induced deformation localization

mechanism is presented. M ore specifically, the suggested mechanism involves the

following events:

a) D ue to corrosion instabilities and patterning, deform ation localization occurs

via a vacancy-induced dislocation process and dense but coarse emergent

slip bands are produced at the specimen surface;

b) localized anodic dissolution occurs at emergent slip bands thus generating

subsurface vacancies;

96

c) vacancies are attracted by dislocations, ease dislocation mobility and cause

a high dislocation density at the near-surface region further localizing slip;

d) a crack is nucleated on {111} slip planes when a critical strain energy is

reached from the externally applied stress and the stress from the

surrounding dislocation pile-ups;

e) in the presence o f a crack, the same process is repeated, but since relatively

high stress intensities prevail, multiple slip occurs that modifies the stress

field causing cracking on {O il} ; Stress-assisted diffusion o f vacancies may

assist the cleavage process by reducing the local fracture toughness;

f) the crack propagates in steps about equal to the depth o f the zone with high

dislocation density and finally is arrested at a pre-existing slip band. The

process is repeated and, as the crack advances, less dissolution is required

since the stress intensity is gradually increased and a smaller num ber o f

dislocations is required to reach the critical stress for cracking.

CHAPTER VIII. SUGGESTED FUTURE RESEARCH

Future research is recom mended on the initial stages o f corrosion (instability

and patterning) in the absence and presence o f stress. The in situ X -Ray diffraction or

extended X-Ray absorption fine structure (EXA FS) spectroscopy may be employed to

study the surface disorder (instability) caused by electrolytes alone and by combination

o f electrolytes and stress. TEM examinations on stress free foils after different period

o f immersion tests can reveal the dislocation density and configuration changes resulted

from pure electrochemical reactions. This will clarify the role o f the SCC environment

in enhancing the dislocation density at the near-surface region.

98

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APPENDIX A. NOMENCLATURE

110

£f strain to fracture in air

et

et

ep

es

escc

<uvw>

{hkl}

ANOVA

BCC

CAC

CF

CIS

da/dt

J-'corr

EXAFS

F

FCC

near-surface strain for crack initiation

total strain

strain for crack initiation

plastic strain

near-surface plastic strain

strain-to-failure under stress corrosion cracking

indices for a family o f crystallographic directions

M iller indices for a family o f crystallographic planes

analysis o f variance

body-centered cubic crystal structure

corrosion-assisted cleavage

corrosion fatigue

crack initiation spacing

crack initiation rate

open circuit corrosion potential

extended X-ray absorption fine structure

Faraday’s constant (96500 C/mole)

face-centered cubic crystal structure

FW HM : full-width at half maximum

GLM : general linear models procedure

H C P : hexagonal close-packed crystal structure

H E : hydrogen embrittlement

IG : intergranular

Ki : stress intensity factor

Kic : plane strain fracture toughness for m ode I loading

Kiscc : threshold stress intensity factor for stress corrosion cracking

LEED : low-energy-electron-diffraction

LM E : liquid metal embrittlement

SBH : slip band heights

SBL : slip band length

SBS : slip band spacing

SCC : stress corrosion cracking

SCE : saturated calomel electrode

SEM : scanning electron microscopy

SFE : stacking fault energy

TEM : transmission electron m icroscopy

TG : transgranular

W CC : w ear-corrosion cracking

XRD : X-ray diffraction

APPENDIX B. HYDROSTATIC STRESS AND PLASTIC ZONE CALCULATION

113

The hydrostatic stress in the plastic zone ( Op) immediately ahead o f a blunted crack o f a radius p is given by:

Op = O y[ln (l+ x /p) + 1/2]

w here O y is the material's yield strength and x is the distance in front o f the crack tip. The size o f the plastic zone (rp) is given by:

rp = (l/6 n )(k ISCC/ Oy)2

The values o f the various param eters used to determine the stress distribution and the divacancy concentration profiles (Fig. 31) are as follows: oy = 216 M Pa, p - 0.6 pm,

f p = 6 pm , V = 0.3, D v = lx lO ’16 m2/s, k ]SCC = 2.3 M Pam 1/2, Vy = 9 .16X10-6 m3/mol.

APPENDIX C. POTENTIODYNAMIC TEST RESULTS OF a-BRASS IN DIFFERENT ELECTROLYTES

115

»: .-•« m

116

LI

sm

1 0 1 0 1 0 10I < U A / C M •i ;i

Figure 32. Potentiodynamic test results of a-brass in 5 M NH4OH plus Cu(N03) 2solution.

.r"™ISO

3 0 0

- 4 5 0 i i i I i ml i i i I i ml I i l l I m ],—1,7-1, Jj l.ljJ___1—LJ-LL11JI

10- 1

100

1 0 10 10 IC U A / C M A 2 i

1 0

J I 1 .J J X1.J4

1 0

Figure 33. Potentiodynamic test results of a-brass in 1 M NaNC>2 solution.

M118

1 5 0

Ul 0

"1------ r ~l~ P 'TTTI--------- 1----- I— |—pTTTT]--------- 1---T T * f T 11 I I---------1----- 1 I | I I I I

/4 ?'/

-= = C ir'-“a..

X

I I i _ 1 j _ ui i i_L J .au

1 0

K U A / C M a Z)

Figure 34. Potentiodynamic test results of a-brass in 0.1 M CuS04 solution.

VITA

Kun Lian w as born on February 11, 1957 in W uhan, Hubei province, People’s

Republic o f China. After completing his highschool education in 1975, he w ent to the

countryside where he did three years o f field work. In February, 1982, he received the

Bachelor o f Science D egree in the Departm ent o f Chemical Industry M achinery, at

South China Institute o f Technology, Guan-Zhoung, P. R. o f China. From February,

1982 to December, 1987, he w orked at W uhan Research Institute o f M aterials

Protection, M inistry o f M achinery Industry as associate corrosion engineer and

corrosion engineer. H e began his graduate studies at Louisiana State University in the

Interdepartm ental Program in M aterials Science in the College o f Engineering in Fall

1989, and received the M aster D egree in Engineering Science in August, 1991. After

receiving his M .S., he w orked at Louisiana State University School o f Dentistry,

Biomaterial Departm ent as a research associate, and at the same time he w as a part-

time student at Louisiana State University pursuing his Ph.D. in Engineering Science.

H e received his D octor o f Philosophy D egree at the Summer commencement, 1995.

119

DOCTORAL EXAMINATION AND DISSERTATION REPORT

Candidate: Kun Lian

Major Field: Engineering Science

Title of Dissertation: a Study of the Stress Corrosion CrackInitiation Stage in Alpha-Brass

Approved:

P ' I 4 ; y >ior Professor and ChaiChairman

raduate School

EXAMINING COMMITTEE!

; t C ^ j u

Date of Examination:

0 6 / 1 9 / 1 9 9 5