Localised Corrosion of Ni-base Superalloys in Seawater

1

Localised Corrosion of Ni-base Superalloys in

Seawater

Thesis submitted to the University of Manchester for the degree of Doctor of Philosophy in the Faculty of

Science and Engineering

2019

Melissa M Keogh

School of Materials

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Contents Abbreviations ........................................................................................................... 8

Abstract .................................................................................................................. 10

Declaration ............................................................................................................. 11

Copyright ................................................................................................................ 11

Acknowledgments .................................................................................................. 13

Chapter 1; Introduction.............................................................................................. 15

1.1 Corrosion of Subsea Wellheads ....................................................................... 16

1.2 Aims and Objectives ......................................................................................... 17

1.3 Thesis Outline ................................................................................................... 19

1.4 References ........................................................................................................ 19

Chapter 2; Introduction to corrosion ......................................................................... 21

2.1 Corrosion Fundamentals .................................................................................. 21

2.1.2 Thermodynamics of Corrosion .............................................................................. 23

2.1.3 Kinetics of Corrosion ............................................................................................. 24

2.2 Types of Corrosion ........................................................................................... 26

2.2.1 Galvanic Corrosion ................................................................................................ 26

2.2.2 Pitting Corrosion ................................................................................................... 26

2.2.2.1 Passive Film Breakdown ..................................................................... 27

2.2.3 Crevice Corrosion .................................................................................................. 30

2.2.3.1 Theories of Crevice Corrosion ............................................................ 32

2.2.4 Intergranular Corrosion ........................................................................................ 34

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2.2.5 Selective Leaching ................................................................................................. 34

2.2.6 Erosion Corrosion .................................................................................................. 35

2.2.7 Stress Corrosion Cracking ..................................................................................... 35

2.2.8 Hydrogen Damage ................................................................................................ 36

2.3 Oil and Gas Environments ................................................................................ 36

2.3.1 Offshore ............................................................................................................... 36

2.3.2 Seawater ............................................................................................................... 38

2.2.3 Corrosion protection in oil and gas ....................................................................... 39

2.3.4 Material Selection ................................................................................................. 40

2.3.4.1 Superalloys ......................................................................................... 41

2.4 Nickel Superalloys ............................................................................................ 42

2.4.1 Microstructure ...................................................................................................... 44

2.4.2 Alloy 718................................................................................................................ 47

2.4.3 Alloy 625................................................................................................................ 53

2.4.4 Alloy 625+.............................................................................................................. 55

2.5. References ....................................................................................................... 56

Chapter 3; Experimental ............................................................................................ 62

3.1 Optical Microscopy........................................................................................... 63

3.2 Scanning Electron Microscopy ......................................................................... 63

3.2.1 Quanta 200............................................................................................................ 63

3.2.2 Quanta 650............................................................................................................ 64

3.2.3 EDS ........................................................................................................................ 64

3.3 Confocal Laser Microscopy .............................................................................. 65

3.3.1 Background ........................................................................................................... 65

3.3.2 Principle of Operation ........................................................................................... 65

3.3.3 Operation .............................................................................................................. 66

3.4 Metallographic Preparation ............................................................................. 67

3.4.1 Sample Fabrication................................................................................................ 67

3.4.2 Heat Treatments ................................................................................................... 68

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3.4.3 Hardness Testing ................................................................................................... 70

3.4.4 Sample Preparation............................................................................................... 70

3.5 Electrochemical Methods ................................................................................ 71

3.5.1 Reference and Counter Electrodes ....................................................................... 72

3.5.2 Artificial Crevice Environment .............................................................................. 72

3.5.3 Potentiodynamic Polarisation ............................................................................... 73

3.5.4 Tsujikawa-Hisamatsu Electrochemical Technique ................................................ 73

3.6 References ........................................................................................................ 74

Chapter 4; Materials Characterisation of Alloy 718 and Alloy 625+ .......................... 76

4.1 Abstract ............................................................................................................ 76

4.2 Introduction ..................................................................................................... 76

4.3 Experimental .................................................................................................... 77

4.4 Results .............................................................................................................. 78

4.4.1 Hardness Testing .................................................................................................. 78

4.1.2 718 ........................................................................................................................ 79

4.1.3 625+ ...................................................................................................................... 80

4.5 Discussion ......................................................................................................... 82

4.6 Conclusions ...................................................................................................... 83

4.7 References ........................................................................................................ 83

Chapter 5; Crevice corrosion behaviour of Inconel 718 and Custom Age 625+ in

sodium chloride solution............................................................................................ 85

5.1 Abstract ............................................................................................................ 85

5.2 Introduction ..................................................................................................... 85

5.3 Experimental .................................................................................................... 87

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5.4 Results .............................................................................................................. 90

5.4.1 Microstructure ...................................................................................................... 90

5.4.2 Corrosion behaviour of AR-718 ............................................................................. 91

5.4.3 Corrosion behaviour of AR-625+ ........................................................................... 96

5.5 Discussion ......................................................................................................... 99

5.6 Conclusions .................................................................................................... 102

5.7 References ...................................................................................................... 102

Chapter 6; Effects of Microstructure on localised corrosion behaviour of Inconel 718

and Custom Age 625+ .............................................................................................. 105

6.1 Abstract .......................................................................................................... 105

6.2 Introduction ................................................................................................... 106

6.3 Experimental .................................................................................................. 107

6.4 Results ............................................................................................................ 110

6.4.1 718 ...................................................................................................................... 110

6.4.2 625+ .................................................................................................................... 114

6.5 Discussion ....................................................................................................... 117

6.6 Conclusions .................................................................................................... 120

6.7 References ...................................................................................................... 120

Chapter 7; Effects of Chloride concentration on Crevice corrosion behaviour ....... 122

7.1 Abstract .......................................................................................................... 122

7.2 Introduction ................................................................................................... 123

7.3 Experimental .................................................................................................. 124

7.4 Results ............................................................................................................ 126

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7.4.1 Electrochemical Results for 718 .......................................................................... 126

7.4.2 Electrochemical Results for 625+ ........................................................................ 128

7.4.3 Post Immersion Sample Analysis ........................................................................ 132

7.5 Discussion ....................................................................................................... 137

7.6 Conclusions .................................................................................................... 139

7.7 References ...................................................................................................... 140

Chapter 8; Investigation of the galvanic crevice corrosion behaviour of Inconel 718

and Custom Age 625+ .............................................................................................. 143

8.1 Abstract .......................................................................................................... 143

8.2 Introduction ................................................................................................... 144

8.3 Experimental .................................................................................................. 145

8.4 Results ............................................................................................................ 147

8.5 Discussion ....................................................................................................... 157

8.6 Conclusions .................................................................................................... 159

8.7 References ...................................................................................................... 160

9; Discussion and Conclusions ................................................................................. 163

9.1 Corrosion Behaviour ...................................................................................... 163

9.2 Microstructure ............................................................................................... 164

9.2 Environment ................................................................................................... 165

9.4 Galvanic Crevice Corrosion ............................................................................ 166

9.4 Conclusions .................................................................................................... 168

9.5 References ...................................................................................................... 169

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Chapter 10; Future Work: ........................................................................................ 171

10.1 Microstructure ............................................................................................. 171

10.1.1 THE for other microstructures .......................................................................... 171

10.1.2 Confirm CCT through potentiostatic temperature ramping ............................. 172

10.2 Environment ................................................................................................. 173

10.2.1 Other Chloride concentrations ......................................................................... 173

10.2.2 Galvanic Corrosion of other heat treatments ................................................... 173

10.2.3 Combined effects of crevice and stress ............................................................ 174

10.3 References .................................................................................................... 174

Word Count: 29432

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Abbreviations AR As Received

ASTM American Society for Testing and Materials

BCC Body Centred Cubic

BOP Blowout Preventer

CCS Critical Crevice Solution

CCT Critical Crevice Temperature

CPP Cyclic Potentiodynamic Polarisation

DC Direct Current

E Potential of the Electrochemical Cell

EDM Electron Discharge Machining

EDS Energy-dispersive X-ray spectroscopy

F Faraday Constant

ΔG Gibbs Free Energy

FCC Face-centred-cubic

HT1 Heat Treatment 1

HT2 Heat Treatment 2

LSCM Laser Scanning Confocal Microscopy

OCP Open Circuit Potential

OM Optical Microscope

OPS Active Oxide Polishing Suspensions

PDP Potentiodynamic Polarisation

PTFE Polytetrafluroethylene

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RT Room Temperature

RPM revolutions-per-minute

SA Solution Annealed

SCC Stress Corrosion Cracking

SE Secondary Electron

SEM Scanning Electron Microscope

THE Tsujikawa-Hisamatsu Electrochemical Technique

TTT Time-Temperature-Transformation

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Abstract Nickel Superalloys are used extensively within the oil and gas industries. Theses alloys

have good resistance to a wide variety of corrosive environments in industrial

processes such as chemical and petrochemical processing, marine engineering,

oil/gas production and transport, and nuclear reactors. Similarly to general grade

steels they are readily available, but nickel-base alloys are preferred due to their

superior corrosion resistance. Within oil and gas they are often utilised as sub-sea

connectors. The geometry of these connectors makes the alloys susceptible to a

crevice corrosion, despite the alloys good resistance to general corrosion

This project focuses on determining the corrosion properties of two nickel alloys;

Inconel 718 (718) and Custom Age 625+ (625+); within environments appropriate to

the oil & gas industry; notably chloride-rich electrolytes. The work produced will aid

in determining whether these alloys are likely to suffer crevice corrosion when in

service as sub-sea connectors.

Electrochemical techniques, including potentiodynamic polarisation (PDP), and a

modified Tsujikawa-Hisamatsu Electrochemical (THE) technique were used to obtain

data on the crevice corrosion behaviour, the critical crevice temperature (CCT), and

the crevice potential (ECREV), of the materials at different temperatures.

Heat treatments were additionally employed to manipulate the microstructures of

the alloys, so that the effect of the precipitates in the microstructure on the corrosion

resistance could be investigated. Microstructures which contained the combination

of precipitates found in these alloys (γ’, γ”, δ-phase) had the highest corrosion

resistance, with alloy 625+ being more resistant to localised corrosion than alloy 718.

Three chloride concentrations were tested in combination with the microstructural

conditions and temperature effects to assess if chloride concentration had a

significant role in the crevice corrosion behaviour. Although the chloride

concentration could affect the CCT, the temperatures at which these alloys

underwent crevice corrosion are unlikely to be of concern in their current in-service

deployment.

Crevices which had formed were observed using Scanning Electron Microscopy (SEM)

and Laser Scanning Confocal Microscopy (LSCM). Both alloys suffered crevice

corrosion through an intergranular attack pathway for heat treatments where there

was no precipitation, or where the gamma precipitates were dominant. When the

delta phase was the dominant precipitate, it was the matrix which provided a

corrosion pathway.

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Declaration

No portion of the work referred to in the thesis has been submitted in support of an

application for another degree or qualification of this or any other university or other

institute of learning.

Copyright

i. The author of this thesis (including any appendices and/or schedules to this thesis)

owns certain copyright or related rights in it (the “Copyright”) and s/he has given The

University of Manchester certain rights to use such Copyright, including for

administrative purposes.

ii. Copies of this thesis, either in full or in extracts and whether in hard or electronic

copy, may be made only in accordance with the Copyright, Designs and Patents Act

1988 (as amended) and regulations issued under it or, where appropriate, in

accordance with licensing agreements which the University has from time to time.

This page must form part of any such copies made.

iii. The ownership of certain Copyright, patents, designs, trademarks and other

intellectual property (the “Intellectual Property”) and any reproductions of copyright

works in the thesis, for example graphs and tables (“Reproductions”), which may be

described in this thesis, may not be owned by the author and may be owned by third

parties. Such Intellectual Property and Reproductions cannot and must not be made

available for use without the prior written permission of the owner(s) of the relevant

Intellectual Property and/or Reproductions.

iv. Further information on the conditions under which disclosure, publication and

commercialisation of this thesis, the Copyright and any Intellectual Property and/or

Reproductions described in it may take place is available in the University IP Policy

(see http://documents.manchester.ac.uk/DocuInfo.aspx?DocID=24420), in any

relevant Thesis restriction declarations deposited in the University Library, The

University Library’s regulations

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(see http://www.library.manchester.ac.uk/about/regulations/) and in The

University’s policy on Presentation of Theses

13

Acknowledgments I would like to thank my Supervisor, Dr Dirk Engelberg, not only for his continued

support, encouragement and kindness throughout my PhD but getting me to the

finish line. Thanks to Professor Robert Akid for giving me the initial opportunity to

pursue this PhD. I would also like to thank the CDT for Advanced Metallic Systems;

Dr Claire Hinchcliffe and Professor Brad Wynne; their support throughout this

process has been unwavering, nor would I have met such a wonderful cohort of

people. Additionally I would like to thank BP for their sponsorship and guidance.

To Dr Tony Cook, Dr Rafa Leiva-Garcia and Dr David Martelo – my thanks to you is

endless. Not only have you helped with the academic aspects of this project, but you

provided a much needed friendship and comic relief in the times where everything

felt like it was going wrong.

I would like to thank everyone in the BP Group; both past and present; Clara, Chris,

Dhinakaran, Gaurav, Jake, Karyn, Yasser, Rob and Phil. You are all wonderful and I thank

you for putting up with me.

To the University of Manchester Counselling service. Your services and the support you

gave was vital for me completing my PhD – thank you.

A hug thanks to the technical staff within the School of Materials, in particular Stephen

Blatch, Harry Pickford, Paddy Hill, and Mike Faulkner who were always happy lend a

helping hand.

Thank you to Emma Lewis-Kalubowila and Olwen Richert. You were both wonderful

source of care and advice during the times I wasn’t sure I was cut out for doing a PhD,

and I wouldn’t have finished my journey without you.

Chapter 1; Introduction

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I would like to thank everyone involved at Old Bedians RUFC, especially Jennifer Tumbri.

You all made me feel so welcome, and often gave me well needed words of wisdom along

with a glass of prosecco or two to help wash them down!

Thank you Paul Whiteside and Dave Campbell for understanding the demands of writing

a thesis whilst working; and for allowing me the time to get it finished. And to my new

colleagues and graduates, in particular Sarah Mundy, for being a great source of support.

To my PhD-Sisters-In-Arms; Felicity Baxter, Emily Cooksey and Rhys Archer. Only you

could truly understand what the realities of doing a PhD were. Your friendship and

guidance, and emergency tea breaks to re-hydrate from the crying, were the light in the

darkness; and without you I would not be here with a finished thesis. I feel so lucky to

have met such strong women, and I feel honoured to call you my friends.

Sophie and Ciara, you were always there for me no matter what and our friendship has

endured and will be one that that lasts a lifetime.

Thank you to all my family and friends; in particular my Mum. I would not be me without

the support of my Mum. She has been by my side through every point in my life, and this

PhD was no different. She is the most amazing person I know and I thank her for always

being there in everything I do.

Finally to Patrick, my best-friend and partner. You had the hardest job of all; putting up

with me through the highs and all of the lows. Your support for me never wavered, and

your love and hugs were often needed at the end of the day. Without you by side I could

not have finished this journey. Thank you.


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CHAPTER 1; INTRODUCTION

As the demand for oil and gas has increased, the search for new reserves has had to

extend to deeper territories than previously accessed. These deeper territories often

include reservoirs found deep below the seabed. To access the subsea reservoirs

equipment, such as that shown in Figure 1,1 needs to be suitable to withstand the

harsh and varied conditions on the sea floor.

Part of the subsea oil and gas recovery equipment is the subsea wellhead, also known

as a Christmas tree due to its unusual geometry shown in Figure 2.2 The subsea well

head is responsible for housing important components; such as those for maintaining

Figure 1 A typical subsea production set-up1


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pressure within the pipelines called the blowout preventer (BOP) and for providing

an interface between the subsea pipelines.2

1.1 Corrosion of Subsea Wellheads The subsea environment is a very aggressive one. It has a high chloride concentration

of approximately 3.5 wt%3 and fluctuations in temperature from -2°C to 40°C.4 Not

only this, but the subsea wellheads need to withstand pressures of up to 20,000 psi5

(1361 atm).

Nickel superalloys are often employed as nuts and bolts in the subsea wellhead

construction. Despite the alloys generally high resistance to general corrosion due to

Figure 2 Typical 18¾-in.Subsea Wellhead System2


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the geometry of the components as shown in Figure 36, they are susceptible to

localised corrosion.

Figure 3 Bolt geomatrey showing crevice corrosion6

The subsea well head is usually under cathodic protection, which creates a high

concentration of hydrogen surrounding the wellhead. Should be cathodic protection

fail, the components are then susceptible to localised corrosion. When the cathodic

protection is re-established, the localised damage area allows ingress of the cathodic

hydrogen into the matrix; thus leading to hydrogen embrittlement.

Although hydrogen embrittlement is thought to be one of the leading causes of

premature failure in subsea wellheads, little is known about the initial stages in the

likelihood of localised corrosion occurring in these alloys.

1.2 Aims and Objectives The key motivation of the research carried out in this thesis is on the crevice

corrosion behaviour of alloys Inconel 718 and Custom Age 625+; which are often

utilised as the nuts and bolts in the subsea wellheads, and the aim of this thesis is to


18

investigate when localised corrosion is likely to take place to aid in improving lifetime

prediction calculations, and what role the microstructure plays on the corrosion

behaviour.

The aim of the experimental portion of this project was to simulate a crevice

corrosion environment in which the crevice corrosion properties of the alloys could

be investigated under controlled situations in order to fully understand the effects of

each changing parameter. In order to do this six main objectives are identified, and

these are listed below.

1. Reproduce a crevice corrosion environment using a ceramic crevice former in

deaerated chloride solutions.

2. Investigate microstructural changes after heat treatment.

3. Understand the role of temperature for each alloy and heat treatment by

recording the Critical Crevice Temperature (CCT) found in 3.5 wt% NaCl

solutions.

4. Understand the role of the microstructure on the localised corrosion

behaviour for both alloys.

5. Investigate effects of chloride concentration through looking at the change in

the CCT for each alloy and heat treatment in 0.1 M and 1 M chloride solutions.

6. Compare different electrochemical techniques includes PDP, THE and

Galvanostatic testing.

7. Study crevice geometries and post-corrosion sign-posts to degradation

mechanisms through using SEM, EDS and Laser Scanning Confocal

Microscopy.


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1.3 Thesis Outline This thesis contains 10 chapters. Chapter 2 contains a literature review focusing on

the alloys used and the crevice corrosion mechanism. Chapter 3 details information

regarding the main experimental processes and analytical techniques used in this

project.

Chapters 4-8 provide details of the results from the research section. Each of these

chapters contains specific introductions, results and conclusions. Chapter 4 contains

the results from the material characterisation that was carried out on the Inconel 718

and Custom Age 625+; including details of heat treatments and resultant

microstructures. Chapter 5 presents the results from the effects of temperature, and

microstructure, on the crevice corrosion behaviour of the two alloys in the as-

received microstructure in a 3.5 wt% NaCl solution. Chapter 6 furthers the results

presented in Chapter 5 by extending them to three additional microstructures,

Chapter 7 explores the effect of chloride concentration, coupled with the

temperature and microstructure effects on the crevice corrosion behaviour. The final

results chapter, Chapter 8, presents the influence that galvanic coupling may have on

the crevice corrosion behaviour of the two alloys.

Chapter 9 is a general discussion and summary of the all of work that has been

completed during this project. Chapter 10 presents any possible future work which

could be employed to prolong the investigations beyond the scope of this PhD.

1.4 References 1. Aven, T. & Pedersen, L. M. On how to understand and present the

uncertainties in production assurance analyses, with a case study related to a


20

subsea production system. Reliab. Eng. Syst. Saf. 124, 165–170 (2014).

2. Bai, Y. & Bai, Q. in Subsea Engineering Handbook 703–761 (Elsevier, 2012).

3. Millero, F. J., Feistel, R., Wright, D. G. & McDougall, T. J. The composition of

Standard Seawater and the definition of the Reference-Composition Salinity

Scale. Deep Sea Res. Part I Oceanogr. Res. Pap. 55, 50–72 (2008).

4. Pawlowicz, R. Key Physical Variables in the Ocean: Temperature, Salinity, and

Density. Nat. Educ. Knowl. 4, 13 (2013).

5. Pathak, P. D., Kocurek, C. G. & Taylor, S. L. Design Method Combining API and

ASME Codes for Subsea Equipment for HP/HT Conditions Up to 25,000-psi

Pressure and 400°F Temperature. Oil Gas Facil. 3, 47–55 (2014).

6. Kahram, M., Asnavandi, M., Koshy, P. & Sorrell, C. C. Corrosion Investigation

of Duplex Stainless Steels in Chlorinated Solutions. steel Res. Int. 86, 1022–

1027 (2015).

Chapter 2; Introduction to corrosion

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CHAPTER 2; INTRODUCTION TO

CORROSION

This chapter aims to give an overview into the background of the topics surrounding

this PhD project. There are three main sections. The goal of the first section is to

provide an understanding of the fundamentals of the corrosion process. The second

section provides detail on the different types of corrosion. The third section discusses

the environment of the oil and gas industry. The final section will include materials

used in these applications and details of earlier studies and research on the topics of

alloy development and crevice corrosion; particularly in nickel alloys related to

Inconel 718 and Custom Age 625+.

2.1 Corrosion Fundamentals Corrosion is most commonly defined as a destructive form of attack by the

environment on a material; although other definitions are available. When the

material of question is a metal, the metal is wanting to return to its lowest energy

state.1 The lowest energy state is the form the metal would be found in the earth’s

crust; usually a type of oxide or sulphide.2 The process of this return to a former state,

takes place via an electrochemical reaction. The metal oxidation reaction is known as

the anodic reaction which drives the corrosion process.3

𝑴 → 𝑴𝒏+ + 𝒏𝒆−

(1)


22

The anodic reactions must be complimented by a cathodic reaction. Oxygen

reduction and hydrogen evolution are the most common cathodic reactions during

aqueous corrosion of metals3 and these follow the form, respectively:

𝑶𝟐 + 𝟐𝑯𝟐𝑶 + 𝟒𝒆− → 𝟒𝑶𝑯−

(2)

𝟐𝑯+ + 𝟐𝒆− → 𝑯𝟐

(3)

Anodic and cathodic sites on a metal surface can occur for a variety of reasons

including; composition, grain size, impurities and localised stresses.4 For

electrochemical corrosion to take place there are four requirements; an anode, a

cathode, a conducting environment (an electrolyte), and an electrical connection

between the anode and cathode.3 The basic principle of electrochemical corrosion is

shown in Figure 1. To fully understand the corrosion kinetics, thermodynamics must

be taken into account.

Figure 1 Example of basic corrosion process


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2.1.2 Thermodynamics of Corrosion

Thermodynamics allows for an understanding of the energy change that take place

within a corrosion system. This energy change, also known as the Gibbs free energy;

ΔG, is what provides the driving force, and dictates the direction of the reaction.5 The

change in enthalpy is denoted by ΔH; temperature, T; and the change in entropy, ΔS.

∆𝑮 = ∆𝑯 − 𝑻∆𝑺

(4)

The Gibbs energy equation can also be written with respect to electric work, which is

a more potent way of calculating electrochemical reactions. Where n is the number

of stoichiometric electrons, F is Faradays constant (96,485 Coulomb Mole-1) and E is

the potential of the electrochemical cell.4

∆𝑮 = −𝒏𝑭𝑬

(5)

Corrosion will not occur spontaneously unless ΔG is less than zero for the metal

oxidation reaction. The potential of the electrochemical cell, E, can be calculated

using the Nernst Equation.6

𝑬 = 𝑬𝒐 + 𝟐. 𝟑 (

𝑹𝑻

𝒏𝑭) 𝐥𝐨𝐠

[𝒐𝒙]𝒐

[𝒓𝒆𝒅]𝒓

(6)

Where Eo is the standard electrode potential, R is the gas constant (8.314 J K−1 mol−1),

T is the temperature in Kelvin, n and F as previously stated are the number of

electrons, and the Faraday constant, (ox) is the oxidation concentration with its

stoichiometric factor o, and (red) is the reduction concentrations with its

stoichiometric factor r.


24

Using the Nernst Equation and the solubility of metals, potential-pH plots, known as

Pourbaix diagrams can be constructed. Pourbaix diagrams allow for differentiation of

regions where a metal may be immune (no corrosion), under protection of a passive

film, or active (undergoes corrosion).7 The most common Pourbaix diagram is that for

Iron8, and for this project the Pourbaix Diagram for nickel9 is of interest (Figure 2).

For each different environment, and alloy system, a different Pourbaix Diagram

needs to be constructed.

Figure 2 A comparison of Pourbaix Diagrams for a) Iron in water at 25°C8 b) Nickel

in water at 25°C9

2.1.3 Kinetics of Corrosion

In addition to corrosion thermodynamics, the kinetics of a corrosion reaction can also

be studied. Tafel’s Law gives the relationship between that the logarithm of the

current density in an electrochemical reaction varies linearly with the electrode

potential.10 The Tafel equation can predict the corrosion rate and potential according

to the kinetics and thermodynamics of all of the reactions taking place on the

electrode surface.

https://www.sciencedirect.com/topics/physics-and-astronomy/electrodes


25

𝐸 = 𝑎 + 𝑏 log 𝑖 (7)

Where a and b are constants. E is the potential and i is the current.

Tafel extrapolation is method for calculating corrosion kinetics or rates of corrosion.

When anodic and cathodic reactions occur the potential of the electrode and the

reversible or equilibrium potentials of each reaction taking place on the surface will

be dissimilar (the overpotential).11

A graph, known as a Tafel plot (Figure 3), can be drawn representing the relationship

between the overpotential and the logarithmic current density. It can, then, be

utilised to find the values of the Tafel slope, corrosion current density, and corrosion

potential utilizing the extrapolation approach.11

Figure 3 Schematic of the Tafel Extrapolation11

Extrapolation of the linear portion of the curve to Ecorr is used to gain the corrosion

current density. If uniform corrosion is assumed then Faraday's law can convert the

corrosion density into the rate of penetration or weight loss.10

https://www.sciencedirect.com/topics/physics-and-astronomy/corrosion-potential

https://www.sciencedirect.com/topics/physics-and-astronomy/corrosion-potential

https://www.sciencedirect.com/topics/physics-and-astronomy/faradays-law


26

Tafel plots are greatly affected by polorisation of the electrode surface, and localised

corrosion,12 and as such this analysis has not been applied to this project.

Despite these basic principles of thermodynamics and kinetics, corrosion is not

limited to one mechanism. Changing the conditions and environments, can alter the

corrosion mechanisms, and some of these are detailed below.

2.2 Types of Corrosion

2.2.1 Galvanic Corrosion

Galvanic Corrosion occurs when two dissimilar metals are electronically coupled to

one another, and are immersed in the same conductive or corrosive solution. The less

corrosion-resistant metal becomes anodic, while the more corrosion-resistant

becomes cathodic, creating a potential difference between the two metals which acts

as the driving force.1 The Galvanic Series can be used to find the order in which metals

will preferentially corrode, and a specific one must be used for each environment

encountered.13

Table 1 The Galvanic Series in Seawater 13

2.2.2 Pitting Corrosion

Pitting corrosion is a localised and accelerated type of dissolution of a metal which

occurs due to the breakdown of a passive film on a metal surface.

Active End Nobel End

Magnesium

AlloysZinc Aluminium

Steel or

Iron

Chromium

Steel (13%)Nickel Brasses Copper Silver Gold Platinum


27

The passive film are thin, usually in the region of nanometres, oxide or nitride layers

that naturally form on a metal surface, protecting against and reducing the rate of

corrosion. These layers, however, are susceptible to localized breakdown, which

allows the metal underneath to corrode. If the attack is initiated at the surface, this

is known as pitting corrosion. This highly localized form of corrosion can result in

premature failure of structural components.14

There are several stages of pitting, including:

1. Passive film breakdown/Pit initiation

2. Metastable pitting

3. Pit growth

There is debate, however, as to which is the most important of these stages and what

the mechanisms is for each step – in particular, the passive film breakdown. One

commonality in the breakdown theories, is that it can only occur in the presence of

an aggressive anodic species, such as chloride.

2.2.2.1 Passive Film Breakdown

A number of theories have been proposed in an attempt to explain the initiation of

pits in surfaces which are free of physical defects. These include:

The penetration mechanism - involves the transfer of anions through the oxide film

to the metal surface (Figure 4).8


28

Figure 4 A schematic of the penetration mechanism8

b. The film breakdown mechanism - requires breaks within the film giving direct

access to the anions to the unprotected metal surface (Figure 5).

Figure 5 Schematic of Film Breakdown Mechanism4

c. The adsorption mechanism – the adsorption of aggressive anions at the oxide

surface, enhancing the catalytic transfer of metal cations from the oxide to the


29

electrolyte, leading to a thinning of the passive layer and potential total removal

(Figure 6).4

Figure 6 Schematic of Adsorption Mechanism4

Metastable pits initiate and grow for a limited period before repassivating. These

metastable pits are usually of microns in size with a lifetime on the order of seconds

or less. Under certain conditions, they can continue to grow forming larger pits. Pit

growth is dependent on the maintenance of pit electrolyte composition and pit-

bottom potential. Conditions must be severe enough in order to prevent

repassivation of the dissolving metal surface at the pit bottom.15

Pitting is considered to be autocatalytic in nature; as once a pit is initiated, the

conditions within the pit become altered, allowing the propagation of further pit

growth.14

Within the pit the metal dissolves in solution to cations (e.g. Ni+2, Cr+3, Fe+2) as in eq.

1. The anodic reaction within the pit is balanced by the reduction of O2 which occurs

on adjacent surfaces surrounding the pit (eq. 2). The rapid dissolution of metals

within the pit produces an excess of positive charges in the pit. This causes migration

of Cl- ions into the pit to maintain electroneutrality. As a result of this, inside the pit


30

there will be high concentration of metal chlorides and a high concentration of

protons (H+) due to metal hydrolysis (eq. 7). Both H+ and Cl- stimulate the dissolution

of most metals and this process accelerates with time, as the aggressive hydrochloric

acid is formed.14

2.2.3 Crevice Corrosion

Crevice corrosion occurs in confined spaces within the sample as shown in Figure 7.

Generally, there is oxygen depletion within the crevice, closely followed by a

decrease in the pH, and breakdown of the passive film causing rapid corrosion.16

Figure 7 A schematic of Crevice Corrosion

aCl lectrolyte

Passive Layer

Crevice

High , H

and Cl

-

concentration. Low

concentration.

Crevice Gap

High

Concentration

etal Sample


31

In practice, two categories of crevices exist:

Naturally occurring: Those created by biofouling, sediment, debris, deposits,

etc.16

Man-made: Those created during manufacturing, fabrication, assembly, or

service.16

The onset of crevice corrosion can take place in four progressive stages:

1. Oxygen depletion from crevice. As a result of small volume of crevice, O2

concentration within the crevice decreases with time due to restricted mass

transfer. This leads to change in kinetics of electrochemical reactions and the

potential of metal within the crevice becoming more negative, and thus the

metal inside the crevice acts as anode and outside the crevice acts as

cathode.17

2. Increase of acidity and [Cl-] in the crevice. Metal cations are introduced

through the passive film into the electrolyte which has two effects: anions,

primarily Cl-, migrate into the crevice area from the bulk solution for charge

neutrality. Hydrolysis reactions involving the metal cations result in increase

in acidity.17

Mn+

(aq) + H2O(l) → M(OH) n-1(aq) + H+

(aq) (8)

3. Permanent breakdown of passivity. Onset of accelerated corrosion occurs

when the electrolyte within the crevice becomes sufficiently aggressive to


32

breakdown the passive film. This solution can be referred to as Critical Crevice

Solution (CCS).17

4. Continued propagation of crevice corrosion.17

2.2.3.1 Theories of Crevice Corrosion

There are four proposed mechanisms for the initiation of crevice corrosion, each with

their own experimental and theoretical support. For this project the metastable

pitting model was chosen as it most represented the results found. The four methods

will now be discussed.

1. Passive Dissolution Model

Developed by Oldfield and Sutton18 they describe how both anodic and cathodic

reactions begin on the metal surface, inside and outside of the crevice area. When

the oxygen is used up within the crevice it begins to act as a local anode, with metal

ions being produced, and the subsequent hydrolysis of the metal ions leading to a

decrease in the pH. When the environment within the crevice reaches a critical

crevice solution (CCS) composition, the passive film becomes unstable, and corrosion

occurs within the crevice.

2. Thiosulphate Entrapment

Whilst investigating the passive dissolution model, Lott and Alkire19 established a

model based on the dissolution of MnS inclusions within the crevice which would

produce thiosulphate ions (S2O32-). The thiosulphate ions, along with chloride ions

would aid in creating the critical condition for breakdown of the passive film. This

model, however, would only be applicable to stainless steels where MnS inclusions

are present, and is thus non-universal for describing crevice corrosion.


33

3. IR Drop

The IR drop mechanism has been proposed by Pickering and Frankenthal20 during

their investigations into the potential drop across the interface of the degrading

surface of the pit. They found that during testing, the potential at the surface of the

pit or crevice could be different to the bulk potential of the metal sample, even during

controlled potential experiments. It was assumed that hydrogen gas was being

produced inside the pits, leading to the potential drop.

4. Metastable Pitting

Stockhert and Boehini21 put forward that crevice corrosion was simply form of pitting

that was stabilised through geometry, similar to pits with a lacy cover. They suggest

that if a metal stable pit can form within the crevice area, it is more likely to become

a stable pit due to the crevice geometry. The metastable pitting theory would indicate

that crevice corrosion is not a sudden onset of corrosion, but more similar to pitting

corrosion in that there are metastable species formed first.22

The metastable pit theory has been supported with studies by Suleiman et al23 who

found that an iron oxide layer deposited on stainless steel was suitable to act as a

crevice initiation site. The crevice corrosion initiated at a similar time to the initiation

of metastable pitting of rust-free samples, thus supporting the metastable pit theory.

A study conducted by Laycock et al22 investigating the different crevice corrosion

initiation theories concluded that their results, after investigating the crevice

corrosion of 316L stainless steel in a 1M NaCl solution, supported the metastable pit

theory as the predicted initiation potentials from the model were concurrent with

the experimental data.


34

More recently, Shu et al24, found that the first corrosive attack in crevice geometry

was a pit. Additional pits then formed, eventually forming a crevice. Similarly Han et

al25, when studying the effects of environment on the crevice corrosion of 2205

duplex stainless steel, found metastable transients were visible in their

electrochemical data, and that post-immersion SEM supported that pits had formed

on the edge of the crevice area.

The evidence produced as part of this thesis also supports the metastable pit theory.

2.2.4 Intergranular Corrosion

Intergranular Corrosion is the selective breakdown of grain boundaries, and/or

closely adjacent regions, without any attack on the grains themselves. The

breakdown is a result of a potential difference between the grain boundary and any

precipitates, impurities or intermetallic phases that form at the grain boundaries.

Precipitates form due to high temperatures, often experienced during fabrication and

welding, preferentially growing at grain boundaries. If the precipitates utilise a high

content of alloying elements used for corrosion resistance, then the region adjacent

to the grain boundary is lacking in these elements, thus becomes sensitized and

susceptible to corrosion.1

2.2.5 Selective Leaching

Selective leaching is the preferential removal of one of the components of an alloy

through corrosion. The most common example of this is the dezincification of brass.

It is usually the least noble component which is susceptible to leaching.26


35

2.2.6 Erosion Corrosion

It was proposed that erosion corrosion caused by the flow of other particles, whether

liquid or solid, over the surface of the metal sample.27 It is commonly found in water

pipes, and depending on the flow of materials, the mechanism can change. It is

measured using:

𝐸 = 𝑚𝑎𝑠𝑠 𝑟𝑒𝑚𝑜𝑣𝑒𝑑 𝑓𝑟𝑜𝑚 𝑠𝑢𝑟𝑓𝑎𝑐𝑒

𝑡𝑜𝑡𝑎𝑙 𝑚𝑎𝑠𝑠 𝑜𝑓 𝑝𝑎𝑟𝑡𝑖𝑐𝑙𝑒𝑠 𝑖𝑚𝑝𝑎𝑐𝑡𝑖𝑛𝑔 𝑠𝑢𝑟𝑓𝑎𝑐𝑒

2.2.7 Stress Corrosion Cracking

Stress Corrosion Cracking (SCC) requires a susceptible metal, an applied stress and a

corrosive medium.28 The cracks nucleate at the surface, becoming highly branched

(Figure 8) during the three stages of crack growth29:

The generation of an environent which causes a crack(s) to initiate

Initiation

Propagation until failure.

Figure 8 A schematic showing Stress Corrosion Cracking

SCC


36

2.2.8 Hydrogen Damage

There is more than one type of hydrogen damage; Hydrogen embrittlement;

Hydrogen Induced Cracking; and Hydrogen Attack.

Hydrogen Embrittlement is the most prominent as it occurs when a cathodic reaction

takes place, producing hydrogen which diffuses to a tri-axial region of tension ahead

of the crack tip. The hydrogen is then within the matrix of the system, and helps in

the deformation causing a brittle fracture. The crack becomes blunt as it moves out

of the hydrogen embrittlement zone.30 Materials of higher strength, are more

susceptible to this type of fracture as they have smaller plastic zones, making them

more prone to brittle failure.30

2.3 Oil and Gas Environments Oil and gas reserves can be found all across the globe, and there are many different

types of rig in operation, both on land and offshore. The focus here will be on offshore

rigs.

2.3.1 Offshore

Offshore operations take place above the sea level on platforms, which can be either

mobile or stationary. The type of rig used is dependent on the operations taking

place, the water level and location. Components above sea-level are known to be

topside, and those beneath are downhole.31 Most offshore locations involve drilling,

of which there can be several purposes.32

Exploratory – used to find potential new reserves across several locations.33

Development – utilise a field with proven reserves to its full potential.33


37

Injection – to store hydrocarbons, usually natural gas or disposing of unwanted

production waters. The production waters can be additionally used in order to regain

pressure in the reservoir.33

Reentry – conducted on existing wells in order to tap into deeper production zones.33

Infill – replaces depleted wells or drilling addition wells in order to maintain

production value in particular area.33

A rotary drilling rig is depicted in Figure 9.33

Figure 9 A rotary drilling rig33

As the rig continues through its life cycle, more seawater is required to be injected

into the oil wells in order to maintain the pressure. This results in an oil and water

mixture being produced.33

Due to the nature of seawater, this can be detrimental to the rigs integrity and

efficiency.


38

2.3.2 Seawater

Seawater is the Earth's most abundant resource covering over 70 % of the Earth's

surface.34 It is classed as an electrolyte, and is predominately composed of a sodium

chloride solution. Its chemistry, however, is more complex as it can contain almost

all of the naturally occurring elements.

Table 2 Seawater Composition34

Seawater can act as an aggressive medium which can attack almost all common

structural materials. There are two main competing processes that operate

simultaneously in these environments:

Chloride ion activity which destroys any passive film

Dissolved oxygen which promotes the repair of the passive film.

Factors which can affect the efficiency of the seawaters corrosion potential include

1. The alloys - composition and any surface films present

2. Seawater - chemistry and composition including oxygen content

3. Environment – pH, microbiological organisms, pollution and contaminants

4. External factors – fluid velocity and temperature35

Although oxygen reduction reaction is the major cathodic reaction that occurs in

seawater, other species such as hydrogen, sulphates, nitrates, or ammonium

compounds can also partake in cathodic reactions.36

Ion or

MoleculeNa+ K+ Mg2+ Ca2+ Sr2+ Cl- Br- F- HCO3

- SO42- B(OH)3

Concentration

mmol/kg469 10.2 53.1 10.3 0.1 546 0.84 0.07 2.3 28.2 0.416


39

Due to its high electrolyte content, seawater is also a very conductive medium with

a relatively high concentration of chlorides which contribute to the development of

pitting and crevice corrosion of steels and nickel based alloys.33

2.2.3 Corrosion protection in oil and gas

Due to the corrosive nature of seawater, and other oil and gas environments,

cathodic protection is often employed to preserve the longevity of components

through protecting them from corrosion attack.

There are two common methods of cathodic protection; impressed current and

galvanic/sacrificial anodes.

The Galvanic System uses a more reactive metal as an auxiliary anode, which will

corrode preferentially. This is due to the positive current flow in the electrolyte from

the anode to the protected metal, so that the whole of the protected metal becomes

a cathode. Commonly used metals are high in the reactivity series and include; Al, Zn

and Mg.37

The Impressed Current method uses an inert anode coupled with an external DC

current to create a current from the anode to the cathode immersed in a bulk

electrolyte.37

The Impressed Current form of Cathodic Protection is a preferable corrosion

prevention treatment as it only requires the application of an external DC current,

whose effectiveness, and efficiency can be monitored constantly, and remotely. This

approach differs to paints and coatings in which defects cannot easily be rectified

once in use, and few measurements for their effectiveness can take place.37


40

2.3.4 Material Selection

There are several factors taken into consideration when selecting a material for an

oil/gas application. These include:

Resistance to service conditions – temperature, corrosion, pressure.38

Mechanical Properties – greater strength is required for downhole operations

due to the increase in demanding conditions.38

Weldability – this is avoided downhole.38

Design life – most components are designed to last for approximately 30

years, and after this upgrades are required.38

Availability and delivery – alternative materials may have to be considered if

the original material is in short supply.38

Complexity of manufacture – easy manufacture id preferred as its easier to

maintain quality control.38

Track record – materials that have been used before are more like to be used

over a new material.38

Weight – most important offshore.38

The ideal material is cheap and readily available. Due to the demanding

environments, however, these materials usually do not have the required properties.

If a more expensive material is required, then in order to reduce costs, it is considered

whether to use the material as cladding.39

If the environment is more aggressive than a clad material can withstand, then

superalloys are often employed.


41

2.3.4.1 Superalloys

The term superalloy, coined in the 1950s, describes ‘an alloy developed for elevated

temperature service, usually based on group VIII elements, where relatively severe

mechanical stressing is encountered and where high surface stability is required40’.

The three main qualities when developing a superalloy are:

Mechanical strength at elevated temperatures

Resistance to corrosion or attack by the environment

Machinability

There are three main classes of superalloy:

1. Iron-Nickel based

Iron–nickel superalloys are commonly used as gas turbine engine blades, discs, and

casings for their low thermal expansion at high temperatures, and enhanced

mechanical properties; but are generally less expensive when compared to other

superalloys.41

Iron–nickel superalloys contain 15–60% iron and 25–45% nickel and are hardened by

both solid solution, and precipitation strengthening. Alloying elements Al, Nb and C

are used to promote the formation of intermetallic precipitates, γ’Ni3(Al,Ti) and γ”

(Ni3Nb) which are similar to those in nickel-superalloys, and various types of

carbides. The precipitates provide good resistance against creep and stress rupture

at elevated temperatures.41

2. Nickel based


42

Nickel-based superalloys have been widely used as rotors, turbine discs, blades, and

bolts in high-performance combustion engines, and for power generation across the

energy industry. The alloys usually combine high strength and corrosion resistance

during service at elevated temperatures so are additionally used in the chemical

industries.42

Nickel superalloys contain at least 50% nickel, with alloying elements, including Cr

(10-20%), Al, Ti and Co as well as Mo, W and C. The alloying elements strengthen the

nickel by solid solution hardening and by forming hard intermetallic precipitates and

carbides. The Cr creates a protective passive layer for the corrosion resistance.41

3. Cobalt based

Cobalt superalloys are frequently used in components that are required to operate

under low stress, but require hot-corrosion resistance; especially in conditions

produced by jet fuel. The corrosion resistance prolongs the life of the engine and

reduces maintenance costs. Cobalt alloys generally have better hot-corrosion

resistance than nickel-based and iron–nickel alloys, but can be more expensive.

Cobalt superalloys contain about 30–60% cobalt, 10–35% nickel, 20–30% chromium,

5–10% tungsten, and less than 1% carbon. Again, the strengthening mechanisms can

be either precipitation or solid solution based.43

2.4 Nickel Superalloys Nickel-base alloys are extremely corrosion resistant materials, and can withstand

environments ranging from sub-zero to high temperatures. Generally, however, the

cobalt super alloy series have a better record in hot corrosion environments.41 In


43

addition to this, they can resist aggressive environments, high stresses and both of

these in tandem. It can be used in both reducing and oxidising environments as they

can act to promote the formation of passive films. They also have good resistance

against, caustic cracking, halide ion environments, freshwater and deaerated non-

oxidising acids.44

One class of the nickel based superalloys are those that are strengthened by

intermetallic compound precipitations in an austenitic FCC matrix. The precipitates

give the alloy enhanced creep resistance; especially when compared to cobalt super

alloys.41 If the alloy contains Titanium or Aluminium, the strengthening precipitate is

γ’. Due to this mechanism, the alloys can be used in either a wrought or cast form

due to the content of γ’ being 0-40% in wrought and as much as 60% in cast. As the

y’ content increases, so does the elevated temperature resistance. The workability of

the wrought alloys, however, decreases resulting in other fabrication methods such

as cast or powder being used. In addition to this, there is a great deal of

morphological control i.e. grain orientation and aspect ratio, which can be used in

order to optimise mechanical properties. In some example, grain boundaries have

been able to be completely eradicated.45

If the alloy contains iobium, the strengthening precipitate becomes the γ’’ phase;

NiNb3. This, however, makes the alloy susceptible to failure at high temperatures as

the γ’’ precipitate is unstable at temperatures above 650C.

The high strength and good corrosion resistance make Ni-based alloys a good

candidate for oil and gas line piping. A more in depth study, focusing specifically at

CO2/H2S environments, which are highly common within this industry found that the


44

corrosion resistance of some prominent Ni-base alloys to be of the order: Inconel 625

> G3 > Inconel 718 > Inconly 825.46

Al-Forzan and Malik47 conducted a survey to test the corrosion resistance of some

steels and Ni-base alloys in three different seawater level environments; fully

submerged, partially submerged and seawater surrounding environment. It was

found that Ni-base alloys were more corrosion resistant than carbon steels.

Being resistant to these aggressive environments allow the alloys to be classed as

Corrosion Resistant Alloys.

2.4.1 Microstructure

These precipitation hardened alloys Nickel Alloys, consist of an austenitic matrix with

the face-centred-cubic (FCC) structure as shown in Figure 10. There are a number of

secondary phases, as shown in Figure 1148, (including γ’, γ” and δ) and the volume

and distribution of each can change the materials performance.48


45

The precipitate phase γ’ follows the FCC crystal structure, is thus coherent with the

matrix, and often follows a spherical or cuboidal morphology. The γ’ composition is

Ni3((Ti,Al)Nb) and it is the prevailing strengthening precipitate despite each

precipitate only being of the nm scale.49

Figure 10 Schematic of Austenitic FCC Structure


46

In comparison, γ” has a body centred tetragonal structure, but remains semi-

coherent with the matrix. The morphology of γ” is usually disk-shaped, also nm in

size, and remains a major part in the strengthening mechanism of the alloy with a

compositions of Ni3(Nb(Ti,Al)).49

The δ phase is incoherent with the matrix as it has an orthorhombic structure and has

a platelet-like morphology often found at the grain boundaries.50 Due to its

incoherency with the matrix, and comparatively larger size (μm), it can act as a crack

initiation site, a pitting initiation site, both reducing fracture toughness51 and

promote hydrogen embrittlement.52

Other phases such as the hexagonal laves-phase, the body centred cubic (BCC) α-

phase, or the σ-phase can be found in the microstructure48 but the effects of these

have not been taken in to account during this investigation.

Figure 11 Microstructure schematic of 71848


47

Within the microstructure, there are also carbides and nitrides present which

precipitate mainly in Nb enriched areas. MC-type carbides are the predominant form

in the microstructure. Although they are not directly responsible for localised

corrosion, they can increase hydrogen embrittlement by acting as irreversible traps

for the hydrogen.53

2.4.2 Alloy 718

Inconel 718 as discovered by H.L Eiselstein is probably one of the most important

alloys in the nickel-based superalloy class.54 It has a complex composition which is

precipitate hardened, allowing it to be used for high temperature applications.

Usually, it is used in its wrought form as there is very little difference between the

mechanical properties of the cast form; as such it meets the strength, creep-rupture

and fatigue crack-propagation requirements.55

Table 3 Inconel 718 Composition

3.3.1 Improving Mechanical Properties

Grain boundary engineering is the process of improving bulk properties of a material

by manipulating the frequency of grain boundaries. This allows for improvements in

stress corrosion cracking, fatigue, weldability and creep performance of pure Ni and

Ni-base alloys.56

Boehlert et al found that by increasing the annealing temperature, and changing the

conditions of cold rolling, that the grain size also increased. The best mechanical

Ni Cr Fe Nb+Ta Mo Al Co Cu Mn Si P S B

53.4 19.6 Balance 5.22 3.09 0.48 0.01 0.01 0.01 0.01 0.004 0.004 0.003

Composition

(weight

percentage)


48

properties were found to come from the annealing at 954C, as this had a more

refined grain size (Table 4). 56

Table 4 Grain Size refinement after annealing56

954 C Annealed 1100 C Annealed

Cold Rolling % Grain Size, μm Grain Size, μm

0 36.2 92.0

10 37.5 77.5

20 30.6 89.1

30 28.9 76.4

40 28.7 83.9

60 23.0 101.8

80 20.0 99.8

Alloy 718 was initially developed for use in the aerospace industry. More recently it

has been employed across the oil and gas industry due to its high yield strength and

good corrosion resistance.

Rebak et al studied the effect of thermal treatments on the localised corrosion

behaviour of 718. They tested the alloy in the as-received (AR) condition, and two

different heat treatments. Heat Treatment 1 (HT1) consisted of a solution annealing

step at 1021°C - 1052°C for a minimum of 1 hour, but a maximum of 2.5 hours,

followed by an age hardening treatment at 774°C - 802°C for 6-8 hours. The second

heat treatment (HT2) involved solution annealing at 980°C for 1 hour, followed by

water quenching and then a two-step aging treatment of 8 hours at 720°C followed

by 18 hours at 620°C.57

The 718 samples used had dimensions of 13 x 13 x 3 mm and were mounted in epoxy

resin which were then wet ground to 600 grit. A 1 cm2 area was lacquered off to

ensure a consistent exposed area. Some samples, however, were found to develop


49

crevice corrosion due to the use of the lacquer. These samples were discarded from

the results.57

Cyclic Potentiodynamic Polarisations (CPP) took place at room temperature (RT) in a

deaerated 3.5 wt% NaCl solution and a deaerated 5% NaCl + 0.5% acetic acid. Figure

12a) shows the CPP for 718 HT2 in deaerated 3.5% NaCl at RT, and Figure 12b) the

CPP for AR 718 in deaerated 5% + 0.5% acetic acid at RT.57 These were the only

polarisation results included in the paper.

The results of these experiments showed the strengthening heat treatments

decrease the localised corrosion resistance of 718.

Figure 12 CPP for 718 HT2 in deaerated 3.5% NaCl at RT b) CPP for AR 718 in

deareated 5% + 0.5% acetic acid at RT57

Chen et al conducted a similar investigation into the effect of aging treatment on the

pitting corrosion of 718 in 3.5 wt% NaCl solutions.58

a) b)


50

The heat treatments were as detailed in Table 5. 59

Table 5 Solution Annealing and Age Hardening Conditions 60

Designation Solution Anneal Age Harden

718-SA 1026°C, 1.5h None

718-BHIA 1026°C, 1.5h 760°C, 4-5h 650°C, 4-5h

718-APIA 1026°C, 1.5h 788°C, 8h

The sample preparation again used an epoxy mount, with a 1 cm2 of working

electrode exposed. They do not mention if any of their samples developed crevice

corrosion. Before polarisation, the working electrode was cathodically polarised at -

1 V for 10 minutes. This was to remove any oxides which may have formed on the

surface60 which can reveal an almost active surface.

Figure 13 Potentiodynamic Polarisation of 718 in 3.5 wt% NaCl60

Much like the previous study, these specimens also underwent pitting corrosion,

albeit at seemingly much higher potentials (Figure 13); and were in agreement with


51

Rebak et al that the heat treatments made the alloy less resistant to localised

corrosion.57

Klapper et al investigated the effects of pH and chloride ion concentration, [Cl-], on

718 in a NaCl solution.61 Two different pH values were investigated; 6 and 10, along

with two [Cl-]; 0.5 M and 4 M. The Cyclic Potentiodynamic Polarisation (CPP) curves

for experiments at room temperature are shown in Figure 14. The samples immersed

in solutions of pH 10, regardless of [Cl-], had much lower potential at which there was

a significant increase in the current density, indicating a lower resistance to corrosion.

An anodic peak was observed at 600 mV for the high pH experiments which is said to

be the transpassive dissolution of chromium. The samples at pH 6, regardless of [Cl-]

did not undergo pitting corrosion, but transpassive dissolution. Pitting was not

experienced for samples immersed in a pH 10 solution until the temperature was

increased to 150°C.61

Figure 14 CPP curves of 718 in buffered solutions of pH 6 and 10 with 0.5 M and 4

M [Cl-] at RT 61


52

Conversely to the RT results, samples immersed in pH 10 with 0.5 M [Cl-] solution did

not undergo pitting corrosion at 150°C (Figure 15). The sample immersed in pH 10

with 4 M [Cl-] did, however, become susceptible to pitting corrosion. This susceptibly

to pitting corrosion is thought to be due to the combined effects of high temperature

and increased chloride concentrations decreasing the passive layer stability.61

Figure 15 CPP of 718 in buffer solutions of pH 6 and 10 with 0.5 M and 4M [Cl-] at

150°C61

Yin et al conducted a study of Inconel 625 and 718, Alloy G3 and Incoloy 825, at 80C

in CO2/H2S corrosion environments. Cyclic Potentiodynamic Polarization and

Electrochemical Impedance Spectroscopy (EIS) techniques were used to monitor the

electrochemical parameters. The study (Figure 16) showed that the corrosion

resistance of the Ni-alloys to CO2 or CO2/H2S corrosion followed the sequence:

Inconel 625 > G3 > Inconel 718 > Incoloy 825.46


53

Figure 16 (Left) CPP of different Ni-base alloys exposed to CO2 at 80°C. (Right)

Potentiodynamic Polarization Curves of different Ni-base alloys exposed to

CO2/H2S at 80°C46

2.4.3 Alloy 625

Development of the alloy began in the 1950s to deal with demand for high-strength

piping, non-age hardening with good weldability and creep resistance, solution

strengthened. In 1964 a patent was granted to H L Eiselstein and J.Gadbut for Alloy

625.62

During the development of Alloy 625, another alloy was born; Inconel 718; for which

the patent was granted in 1962. Inconel 718 filled a lot of the industrial demands that

625 was originally being designed for. In order to advertise to wider markets, the

corrosion resistance of alloy 625 was improved by additional Cr and Mo to the

composition.62

Table 6 Inconel 625 Composition62

Ni Cr Fe Nb+Ta Mo Al Ti Co Mn Si P S B

58 20-23 53.15-

4.158-10 0.4 0.4 1 0.5 0.5 0.0015 0.0015 0.003

Composition

( weight

percentage)


54

During a slow cool after annealing, the BCT y” is able to precipitate but not the FCC

y’ found in other Al-Ti hardened alloys.62

As an alloying element, nickel contributes corrosion resistance, particularly in neutral

salt environments and is resistant stress-corrosion cracking. Chromium, also offers

corrosion resistance, but more to acidic media. It is, however, also required for

passive film formation. Molybdenum, in addition to contributing to mechanical

strength and contributes to the corrosion resistance by providing pitting resistance.62

The corrosion resistance of alloy 625 allowed it to be marketed to other industries

than aerospace including navy and marine use. The weldability of the alloy also makes

it desirable.62

Since alloy 625 was invented, serval variations have been successfully developed and

employed in industry. Inconel 718 has been previously mentioned, which then

inspired further derivations. As demand for great depths to be reached in the oil and

gas industry, stronger and more corrosion resistant alloys were required.62 Figure 17

demonstrates the alloy development initiated by the original research into alloy

625.62

Figure 17 Alloy innovation from Alloy 625


55

Inconel 725 and Custom Age 625 PLUS were developed almost in tandem between

two companies trying to address the issues of the oil and gas industry. Alloy 725 is a

precipitation hardened alloy utilising the strength of the y”, and is particularly

resistant to sulphide stress corrosion cracking. Alloys 625 Plus is also precipitation

hardened, and will be discussed in detail in the next section.62

2.4.4 Alloy 625+

As with the other Nickel Alloys previously discussed, alloy 625+, trade name Custom

Age 625 PLUS, was developed from 625, but with a greater strength level than 625

due to being age-hardenable, and was designed to have superior corrosion resistance

than 718.63

Alloy 6 5 , as is alloy 718, is hardened by precipitation by γ” ( i3,Nb,Ti,Al). Gamma

double-prime (γ”) has a disc-centred morphology64, and is expected to also contain

γ, δ phase and carbides.

When comparing the standard compositions of the three alloys; 625+, 625, and 718,

625+ has a higher titanium content, but reduced carbon and iron contents allowing

for increased chromium, molybdenum which can increase the age-hardening

response whilst enhancing corrosion resistance. The low carbon content also reduces

the number of carbides.63

Little has been found in the literature with regards to the corrosion properties of

625+, but Schmidt et al63 determined the Critical Crevice Temperature (CCT) to be

40°C after immersion in the yellow death solution with consists of 4 wt% NaCl, 0.1

wt% Fe2(SO4). In the same solution, alloy 718 was found to have a CCT of <25°C.


56

When the pitting resistance of alloys 718 and 625+ was explored, Schmidt et al63

found that immersion the samples in the 6 wt% FeCl3 and 1 wt % HCl yielded at Critical

Pitting Temperature (CPT) of >98°C for alloys 625+ and 625, and CPT of between 50-

60°C for alloy 718. Klapper et al65 also investigated the pitting resistance of the alloys

and found 625+ to have a greater resistance than 718.

2.5. References 1. Davis, J. R. (Joseph R. . Corrosion : understanding the basics. (ASM

International, 2000).

2. Kruger, J. & Begum, S. Corrosion of Metals: Overview. Ref. Modul. Mater. Sci.

Mater. Eng. (2016). 3. Speight, J. G. & Speight, J. G. Corrosion. Subsea Deep.

Oil Gas Sci. Technol. 213–256 (2015).

4. Fontana, M. G. (Mars G. Corrosion engineering. (Tata McGraw-Hill, 2005). at

5. Pedeferri, P. in 37–56 (Springer, Cham, 2018).

6. Ciobanu, M., Wilburn, J. P., Krim, M. L. & Cliffel, D. E. Fundamentals. Handb.

Electrochem. 3–29 (2007).

7. Ahmad, Z. & Ahmad, Z. BASIC CONCEPTS IN CORROSION. Princ. Corros. Eng.

Corros. Control 9–56 (2006).

8. Marcus, P. (Philippe). Corrosion mechanisms in theory and practice. (CRC Press,

2012).

9. in Corrosion: Fundamentals, Testing, and Protection 17–30 (ASM International,

2003).


57

10. McCafferty, E. Validation of corrosion rates measured by the Tafel

extrapolation method. Corros. Sci. 47, 3202–3215 (2005).

11. Kakaei, K., Esrafili, M. D. & Ehsani, A. in Interface Science and Technology 27,

303–337 (Elsevier, 2019).

12. Buchanan, R. A. & Stansbury, E. E. Electrochemical Corrosion. Handb. Environ.

Degrad. Mater. 87–125 (2012).

13. in Corrosion: Materials 672–672 (ASM International, 2005).

14. Frankel, G. S. Pitting Corrosion of Metals. J. Electrochem. Soc. 145, 2186 (1998).

15. Frankel, G. S., Stockert, L., Hunkeler, F. & Boehni, H. Metastable Pitting of

Stainless Steel. CORROSION 43, 429–436 (1987).

16. Syrett, B. C. & Begum, S. Corrosion, Crevice. Ref. Modul. Mater. Sci. Mater. Eng.

(2016).

17. Kelly, R. G. & Lee, J. S. Localized Corrosion: Crevice Corrosion. Encycl. Interfacial

Chem. 291–301 (2018).

18. Oldfield, J. W. & Sutton, W. H. Crevice Corrosion of Stainless Steels: I. A

Mathematical Model. Br. Corros. J. 13, 13–22 (1978).

19. Lott, S. E. & Alkire, R. C. The Role of Inclusions on Initiation of Crevice Corrosion

of Stainless Steel I. Experimental Studies.

20. Pickering, H. W. & Frankenthal, R. P. On the Mechanism of Localized Corrosion

of Iron and Stainless Steel. J. Electrochem. Soc. 119, 1297 (1972).

21. Stockert, L. & Böhni, H. Susceptibility to Crevice Corrosion and Metastable


58

Pitting of Stainless Steels. Mater. Sci. Forum 44–45, 313–328 (1991).

22. Laycock, N. J. & Newman, R. C. Localised dissolution kinetics, salt films and

pitting potentials. Corros. Sci. 39, 1771–1790 (1997).

23. Suleiman, M. I. & Newman, R. C. The use of very weak galvanostatic

polarization to study localized corrosion stability in stainless steel. Corros. Sci.

36, 1657–1665 (1994).

24. Shu, H. K., Al-Faqeer, F. M. & Pickering, H. W. Pitting on the crevice wall prior

to crevice corrosion: Iron in sulfate/chromate solution. Electrochim. Acta 56,

1719–1728 (2011).

25. Han, D., Jiang, Y. M., Shi, C., Deng, B. & Li, J. Effect of temperature, chloride ion

and pH on the crevice corrosion behavior of SAF 2205 duplex stainless steel in

chloride solutions. J. Mater. Sci. 47, 1018–1025 (2012).

26. Chatterjee, U. K., Bose, S. K. & Roy, S. K. Environmental degradation of metals.

(M. Dekker, 2001).

27. Humphrey, J. A. . Fundamentals of fluid motion in erosion by solid particle

impact. Int. J. Heat Fluid Flow 11, 170–195 (1990).

28. Sieradzki, K. & Newman, R. C. Stress-corrosion cracking. J. Phys. Chem. Solids

48, 1101–1113 (1987).

29. Jones, R. H. & Ricker, R. E. Cracking.

30. Beachem, C. D. A new model for hydrogen-assisted cracking (hydrogen

“embrittlement”). Metall. Mater. Trans. B 3, 441–455 (1972).


59

31. Pinder, D. Offshore oil and gas: Global resource knowledge and technological

change. Ocean Coast. Manag. 44, 579–600 (2001).

32. Havard, D. Oil and Gas Production Handbook - An introduction to oil and gas

production, transport, refining and petrochemical industry. Power and

Productivity 3, (2013).

33. Azar, J. J. & Samuel, G. R. Drilling engineering. (PennWell Corp, 2007).

34. Boboian, R. Corrosion Tests and Standards: Application and Interpretation -

Google Books. (ASTM International, 2005). at

35. Ray, J. P. (James P. . & Engelhardt, F. R. (F. R. Produced water :

technological/environmental issues and solutions. (Plenum Press, 1992).

36. in Corrosion and Protection 65–87 (Springer London, 2004).

37. Bai, Y. & Bai, Q. Subsea pipelines and risers. (Elsevier, 2005).

38. Standard, I., Vibration, M., Standard, I., Cie, S. E. & Standard, I. International

Standard. 2004, (2004).

39. Lahiri, A. K. in 269–347 (Springer, Singapore, 2017). doi:10.1007/978-981-10-

4684-1_9

40. Whittenberger, J. D. A Review of: “SUP RALL YS II” edited by CT. Sims, N.S.

Stoloff, and W.C. Hagel A Wiley-Interscience Publication John Wiley & Sons,

New York, NY 615 pages, hardcover, 1987. Mater. Manuf. Process. 7, 463–468

(2007).

41. Superalloys, C. Superalloys for gas turbine engines. Introd. to Aerosp. Mater.


60

251–267 (2012).

42. Akca, E. & Gürsel, A. A Review on Superalloys and IN718 Nickel-Based INCONEL

Superalloy. Period. Eng. Nat. Sci. 3, (2016).

43. Coutsouradis, D., Davin, A. & Lamberigts, M. Cobalt-based superalloys for

applications in gas turbines. Mater. Sci. Eng. 88, 11–19 (1987).

44. Craig, B. D. Handbook of corrosion data. (ASM International, 1995).

45. Donachie, M. J. & Donachie, S. J. Superalloys : a technical guide. (ASM

International, 2002).

46. Yin, Z. F., Zhao, W. Z., Lai, W. Y. & Zhao, X. H. Electrochemical behaviour of Ni-

base alloys exposed under oil/gas field environments. Corros. Sci. 51, 1702–

1706 (2009).

47. Al-Fozan, S. a. & Malik, A. U. Effect of seawater level on corrosion behavior of

different alloys. Desalination 228, 61–67 (2008).

48. Engelberg, D. L., Cottis, R. A., Sherry, A. H. & Marrow, T. J. Project Report

Literature Review : Stress Corrosion Cracking of Inconel 718 in PWR

Environments. (2006).

49. Hall, R. C. The Metallurgy of Alloy 718. J. basic Eng. 511–516 (1967).

doi:10.1115/1.3609651

50. ahadevan, S., alawade, S., Singh, J. B. & Verma, A. volution of Δ Phase

Microstructure in Alloy 718. 737–750 (2010).

51. Groh, J. R. & Duvelius, R. W. Influence of Corrosion Pitting on Alloy 718 Fatigue


61

Capability. Superalloys 2001 718, 583–592 (2001).

52. Liu, L., Tanaka, K., Hirose, A. & Kobayashi, K. F. Effects of precipitation phases

on the hydrogen embrittlement sensitivity of inconel 718. Sci. Technol. Adv.

Mater. 3, 335–344 (2002).

53. Singh Handa, S. DEGREE PROJECT FOR MASTER OF SCIENCE WITH

SPECIALIZATION IN MANUFACTURING Precipitation of Carbides in a Ni-based

Superalloy.

54. Paulonis, D. F. & Schirra, J. J. Alloy 718 at Pratt & Whitney: Historical

Perspective and Future Challenges. 13–23 (2012).

55. Michel, D. J. & Smith, H. H. Mechanical properties and microstructure of

centrifugally cast alloy 718. Metall. Trans. A 16, 1295–1306 (1985).

56. Boehlert, C. J., Dickmann, D. S. & Eisinger, N. N. C. The effect of sheet

processing on the microstructure, tensile, and creep behavior of INCONEL alloy

718. Metall. Mater. Trans. A 37, 27–40 (2006).

57. Rebak, Raúl B. Iannuzzi, M. Effect of thermal treatment on the localized

corrosion behaviour of alloy 718 Alloy 718 for Oil & Gas Applications. 3, (2014).

58. Chen, T. et al. Influence of surface modifications on pitting corrosion behavior

of nickel-base alloy 718. Part 2: Effect of aging treatment. Corros. Sci. 78, 151–

161 (2014).


of nickel-base alloy 718. Part 1: Effect of machine hammer peening. Corros. Sci.

77, 230–245 (2013).

Chapter 3; Experimental

62


of nickel-base alloy 718. Part 2: Effect of aging treatment. Corros. Sci. 78, 151–

161 (2014).

61. Klapper, H., Stevens, J. & Hughes, B. Susceptibility to Pitting Corrosion of

Nickel-base Alloy 718 Exposed to Simulated Drilling Environments. Corrosion

70, 899–906 (2014).

62. Eiselstein, H. L. & Tillack, D. J. The Invention and Definition of Alloy 625. 1–14

(2012).

63. Schmidt, N. B., DeBold, T. A. & Frank, R. B. Custom age 625® plus alloy-A higher

strength alternative to alloy 625. J. Mater. Eng. Perform. 1, 483–488 (1992).

64. Voort, G. F. V., Bowman, J. W. & Frank, R. B. Microstructural Characterization

of Custom Age 625 Plus Alloy. 489–498 (2012).

doi:10.7449/1994/superalloys_1994_489_498

65. Klapper, H. S., Zadorozne, N. S. & Rebak, R. B. Localized Corrosion

Characteristics of Nickel Alloys: A Review. Acta Metall. Sin. (English Lett. 30,

296–305 (2017).

CHAPTER 3; EXPERIMENTAL


63

This chapter will detail the experimental methods and procedures followed during

this project including procedures for use of the Scanning Electron Microscope (SEM),

Laser Scanning Confocal Microscope (LSCM), Optical Microscope (OM) and

electrochemical measurements.

3.1 Optical Microscopy A Zeiss AX10 lab A1 microscope was used with a Zeiss AxioCam ERc5s camera. Image

analysis was performed using the Axiovision Rel. 4.8 optical image acquisition

software. Samples have been observed in the optical microscope (OM) to highlight

areas of interest before further investigations took place in the SEM and LSCM.

3.2 Scanning Electron Microscopy During this project, a scanning electron microscope (SEM) was used to observe the

sample surface before and after corrosion experiments, and the interior of any

crevices formed after electrochemical experimentation.

3.2.1 Quanta 200

A Quanta 200 SEM was the primary SEM used throughout this project due to its ease

of use, availability and suitable resolution for taking images of crevices and pits. The

working distance, the distance between the lens and the sample surface, was

approximately 10 mm. A spot size of 3.5 nm and an accelerating voltage of 20 kV were

used for all images.

Most images captured were using the secondary electron (SE) mode as this gave a

good contrast image in which the different topological aspects of the crevice images

could be observed.


64

3.2.2 Quanta 650

A Quanta 650 SEM was used due to its capabilities to take images at higher

magnifications and perform Energy-dispersive X-ray spectroscopy (EDS) when

compared to the Quanta 200. The working distance was approximately 10 mm. A spot

size of 3.5 nm and an accelerating voltage of 20 kV were also used for all images.

3.2.3 EDS

Energy-dispersive X-ray Spectrometers are employed to resolve surface elemental

compositions. When the electron beam interacts with the sample material, inner-

shell electrons are expelled from the material. This creates a vacancy or hole, which

can be filled by a valence electron then resulting in atomic relaxation. The relaxation

causes a photon to be released. The energy of the released photon is characteristic

of the elements found on or near the surface. The intensity of the energy released is

proportional to the amount of that individual element in the bulk of the material. 2

EDS can be used to resolve the elemental composition of a specific point on the

sample surface, across a line scan, or to create elemental maps. During this project,

an Oxford Instruments EDS Detector was used with Aztec Software to create

elemental maps of areas of interest both in and around the crevice. A spot size of 3.5

nm and an accelerating voltage of 20 kV were used. These maps were helpful in

highlighting carbides and areas of elemental depletion.


65

3.3 Confocal Laser Microscopy

3.3.1 Background

Throughout this work, a confocal laser scanning microscope has been used. The laser

scanning capabilities of this microscope have allowed for the construction of 3D maps

of the creviced areas. These maps confirm the existence of crevices on the metal

surface and precise measurements of crevice depth and volume have been able to

be taken.3

The microscope was first created by Marvin Minsky during the 1950s.4 The principle

of confocal microscopy is based on the rejection of the light from planes which are

out of focus. Coupling this principle with the scanning capabilities of the modern

microscopes allows for in-focus imaging of uneven surfaces.5

3.3.2 Principle of Operation

The modern confocal principle is shown in Figure 1.6 Coherent light is emitted by the

laser system, acting as the excitation source, which then passes through two pinhole

apertures. As the laser is reflected by a dichromatic mirror and scanned across the

specimen, fluorescence emitted from the specimen passes back through the

dichromatic mirror and is focused as a confocal point at the detector pinhole

aperture.6


66

Figure 1 Schemcatic of Laser Scanning confocal microscopy6

3.3.3 Operation

A Keyence X200K 3D Laser Microscope was used throughout this project and was

couples with VK Viewer software. Once a sample was placed on operating stage of

the confocal microscope, a magnification of 10x was typically chosen when imaging

a singular crevice. Once the focal places had been selected, the laser was scanned

over the chosen surface area. This results in multiple images being produced; optical,

laser & optical, laser, and a 3D map. Higher magnification was chosen when focussing

on a specific aspect of a crevice. Here the area scanned would be smaller due to the

increased image acquisition time. Using the VK Analyzer plus software, the 3D maps

created can be analysed to reveal volumes within the crevice.


67

3.4 Metallographic Preparation The localised corrosion behaviour of Inconel 718 and Custom Age 625+ has been

investigated and BP Plc. provided the alloys used throughout this project. The alloys

were delivered in large ingots. The large ingots were cut into smaller sections using

EDM (Electron Discharge Machining). The composition of the alloys is given in Table

1 below.

Table 1 Elemental compostion for 718 and 625+

3.4.1 Sample Fabrication

Due to the hardness of these nickel alloys, and the ability for them to undergo work-

hardening, the samples were cut using EDM into 20 mm x 20mm x 5mm cuboids and

were used as baseline reference samples. Samples required for the crevice corrosion

experiments had an addition cut made using EDM; an 8mm hole in the centre the

sample to allow for the crevice former set-up detailed in 3.5.2. These geometries gave

sample areas of 120 mm2 and 97.4 mm2 respectively and are shown in Figure 2.

Wt % Ni Cr Fe Nb+Ta Mo Al Ti

718 54.70 18.60 17.00 5.07 3.04 0.45 0.88

625+ 61.10 21.00 4.24 3.65 8.10 0.22 1.49


68

3.4.2 Heat Treatments

In order to alter the microstructure of the alloys, heat treatments were employed.

The heat treatments SA, HT1 and HT2 were designed from the Time-Temperature-

Transformation (TTT) diagram of IN718 as shown in Figure 3.7 Heat treatments took

place in an argon furnace, and samples were allowed to air cool. Each sample was

ground to remove the black oxidised layer.

In order to ensure each heat treatment was comparable, each sample was first

solution annealed at 1040°C for 1 hour and then air cooled before further heat

treatments took place. Table 2 also details the heat treatments of the as-received

ingots which took place prior to delivery to the University; condition is as-received.

Figure 2 a) working electrode for baseline and pitting experiments (120 mm2) b)

working electrode for crevice experiments (97.4 mm2)


69

Figure 3 Adapted TTT diagram for alloy 7187

Table 2 Desciptions of heat treatments

Designation Heat treatment Expected Microstructure

As received condition (AR)

718: Solution Anneal at 1021°C for 1 hour. Age hardened at 760°C for 8 hours. Air cool.

625+: Proprietary heat treatment by Foroni.

A distribution of y’. y”, delta phase and carbides.

Solution annealed (SA)

Annealed at 1040oC for 1 hour. Air cool.

A solution anneal with precipitates dissolved in microstructure.

Heat Treatment 1 (HT1)

Annealed at 1040oC for 1hr. Air cool. Age hardened at 650°C for 10 hours. Air cool.

Predominately y’ and y” which are disc shaped.


Annealed at 1040oC for 1hr. Air cool. Age hardened at 900oC for 20 hours. Air cool.

Predominant delta phase with platelet like morphology.


70

3.4.3 Hardness Testing

Vickers macro-hardness was performed on each alloy and heat treatment with

sample preparation method described below in 3.4.4, to ensure the heat treatment

was successful. A Armstrong Vickers Hardness Tester was used with a load of 10 kg

and an indentation time of 5 seconds. Indents were placed in five areas of the sample

and an average across the five measurements was used as the reported hardness

value.

3.4.4 Sample Preparation

All samples were cut using EDM into dimensions previously discussed in 3.4.1. The

samples were ground using P80 to P600 Silicon Carbide grinding papers on a rotating

machine set at 300 revolutions-per-minute (RPM) for the electrochemical

measurements. For microstructural etchings the samples were further ground to a

4000 paper. Diamond paste (3 μm to 1μm), followed by PS solution was then

applied to polishing cloths, which were then used to polish the samples to a ‘mirror

finish’; using a rotating machine, with a slower speed setting of 150 RP .

For Inconel 718, the Kallings reagent (100 ml ethanol, 100 ml HCl, 5 g CuCl) was

utilised as the etchant in order to reveal grain boundaries. The samples were

swabbed with the reagent for up to 30 seconds until a good etch had been achieved.

For Custom Age 625+, the 15-10-10 reagent (15 ml HCl, 10 ml acetic acid, 10 ml HNO3)

was used to reveal the grain boundaries. The samples were swabbed with the reagent

until the surface appeared dull; which could take up to 2 minutes.


71

3.5 Electrochemical Methods The three-electrode set-up comprises of a working electrode (Figure 4), the metal

sample; a reference electrode; and the counter electrode which provides a current-

carrying function. All tests were carried out in a deaerated 3.5 wt NaCl solution.

Deaeration was achieved through purging the solution with nitrogen for a minimum

of 30 minutes before testing, and allowing a constant nitrogen blanket throughout

testing. During elevated temperature testing the temperature of the solution was

controlled by immersing the cell in a water bath which was kept at constant

temperature, and monitored with a thermometer placed inside the cell.

Figure 4 Schematic of three-electrode set up


72

3.5.1 Reference and Counter Electrodes

For the electrochemical work carried out an 3 M Ag/AgCl reference electrode was

used alongside a platinum counter electrode.

3.5.2 Artificial Crevice Environment

As per the ASTM standard G48-118, the artificial crevice environment is created when

a non-conducting material is secured to either side of the metal sample in question.

The Crevice Formers made from Technox® 2000, a zirconia ceramic, were used to

create the artificial crevice environment on the metal specimens. The geometry of

the crevice formers consists of a cylindrical base structure, with raised pyramids at

regular intervals which create the crevice environment. The crevice formers are

carefully wrapped in PTFE tape prior to use to extend their lifetime.

The crevice formers are attached to either side of the specimen, and secured in place

using a steel washers, bolt and nut as shown in Figure 5. Plastic tubing is placed

around the bolt to ensure electrical isolation from the sample. The set-up is then

tightened to a torque of 1.5 Nm. By using a consistent torque, a consistent crevice

gap is created for the experiments, and one that is of sufficient size to create the

critical crevice gap size. A torque of 1.5 Nm also ensures the crevice formers are

secure.


73

3.5.3 Potentiodynamic Polarisation

Potentiodynamic polarisation tests were performed using a BioLogic VMP-300

multistat coupled with EC-Lab software. The open circuit potential (Ecorr) was

monitored for 5 minutes in the deaerated conditions. Potentiodynamic polarisation

was used to determine the critical crevice temperature and the crevice susceptibility

behaviour of the alloys. The scans were started at the OCP in the anodic direction at

a rate of 1 mV/s until an end point of +0.5 V or +1.5 V vs Ag/AgCl, where the scan was

then reversed to 0 V vs REF.

3.5.4 Tsujikawa-Hisamatsu Electrochemical Technique

A modified Tsujikawa-Hisamatsu Electrochemical (THE) technique was used to

investigate the protection potential for each alloy and heat treatment. A crevice is

initiated using potentiodynamic polarization, and then held at this constant current

which ensures crevice propagation takes place. The potential is then scanned in the

1 cm

Figure 5 A schematic of the the sample with the crevce formers attached


74

cathodic direction allowing the crevice to repassivate. This reveals the Crevice

Repassivation Potential (Er.crev).9

The modified-THE experiments took place on a Versastat 3F with Versa-studio

software on the adjoining computer, with the sample set up as described in 3.5.2.

The OCP was measured for 1 hour. A potentiodynamic polarisation at 0.167 mV/s was

then started from the OCP in the anodic direction until a current of 0.19 mA was

reached (equivalent of 0 μA / cm2). This current was then held for 2 hours to ensure

crevice propagation. A potentiodynamic scan was then started from the resultant

potential during the galvanostatic hold until 0 V vs OCP is reached.

3.6 References 1. Egerton, R. F. Physical principles of electron microscopy. An Introduction to

TEM, SEM and AEM (2005).

2. Goldstein, J. I. et al. Energy Dispersive X-ray Spectrometry: Physical Principles

and User-Selected Parameters. in Scanning Electron Microscopy and X-Ray

Microanalysis 209–234 (Springer New York, 2018).

3. Leiva-Garca, R., Garca-Antn, J. & Jos, M. Application of Confocal Laser Scanning

Microscopy to the In-situ and Ex-situ Study of Corrosion Processes. in Laser

Scanning, Theory and Applications (InTech, 2011).

4. Minsky, M. Memoir on inventing the confocal scanning microscope. Scanning

10, 128–138 (1988).

5. Guiñón-Pina, V. et al. Influence of temperature on the corrosion behaviour and

on the hydrogen evolution reaction on nickel and stainless steels in LiBr


75

solutions. Int. J. Electrochem. Sci. 6, 6123–6140 (2011).

6. Claxton, N. S., Fellers, T. J. & Davidson, M. W. Laser scanning confocal

microscopy. J. Opt. Microsc. … 1979, (2006).

7. Xie, X. et al. Ttt Diagram of a Newly Developed Nickel-Base Superalloy - Allvac®

718PlusTM. Superalloys 718, 625, 706 Deriv. 193–202 (2005).

8. ASTM G 48-11. Standard Test Methods for Pitting and Crevice Corrosion

Resistance of Stainless Steel and Related Alloys by Use of Ferric Chloride

Solution. 1–13 (2011).

9. Mishra, a. K. & Frankel, G. S. Crevice corrosion repassivation of Alloy 22 in

aggressive environments. Corrosion 64, 836–844 (2008).

Chapter 4; Materials Characterisation of Alloy 718 and Alloy 625+

76

CHAPTER 4; MATERIALS CHARACTERISATION OF ALLOY 718 AND ALLOY

625+

4.1 Abstract Material properties of the alloys Inconel 718 and Custom Age 625+ were

characterised through hardness testing, optical microscopy and SEM after

undergoing heat treatments to alter the microstructure. Both alloys are precipitation

hardened with an austenitic grain structure. The precipitates include γ’, γ” and a δ

phase. Alloy 718 was found to have a smaller grain size when compared to 625+ and

a lower Vickers hardness values indicating less effective strengthening precipitates.

4.2 Introduction Precipitation hardened nickel alloys are used extensively within the oil and gas

industry,1 as well as in aerospace,2 and nuclear technologies.3 They are preferable

over other common materials such as stainless steels due to their increased strength

and corrosion resistant properties4.

Alloy 718 was first used as an aerospace alloys, and was commonly employed in

turbines.5 Within oil and gas it is used as a connector or fastener;1 usually in subsea

conditions. Alloy 625+ was designed to be an improved version of alloy 625; and now

has uses across most industries.6


77

In this work, the as-received microstructures of the alloys have been revealed

through etchings. In addition, three heat treatments were employed to alter the

microstructure. The alloys are precipitation hardened they have 3 common types of

precipitation; γ’, γ” and a δ phase. The heat treatments were designed through

studying the TTT diagram for 718 (Figure 3, Chapter 3), to study the effects of the

different effects the precipitates have on the corrosion properties. The solution

anneal dissolves all precipitates. Heat Treatment 1 was designed to encourage the γ’

and γ” and precipitates; and Heat Treatment the δ phase.

This chapter will focus on resolving the microstructure after heat treatments, and

future chapters will discuss the effects of the microstructure on the localised

corrosion behaviour of the two alloys.

4.3 Experimental The heat treatments employed in order to alter the microstructures, and those of the

samples in their as received conditions, are described in Table 1.

Table 1 Heat Treatment conditions

Designation Heat treatment Expected Microstructure




A distribution of y’. y”, delta phase and carbides.



A solution anneal with precipitates dissolved in microstructure.


Annealed at 1040oC for 1hr. Air cool. Age hardened at 650°C for 10 hours. Air cool.

Predominately y; and y” which are disc shaped.


78


Annealed at 1040oC for 1hr. Air cool. Age hardened at 900oC for 20 hours. Air cool.

Predominant delta phase with platelet like morphology.

Vickers macro-hardness values were obtained, using a 10 kg load on an Armstrong

Vickers Hardness Machine, for each alloy and heat treatment to ensure the heat

treatments were successful.

Samples were successively ground and polished to a ¼ μm finish. For Inconel 718, the

Kallings reagent (100 ml ethanol, 100 ml HCl, 5 g CuCl) was swabbed on the sample

surface up to 30 seconds to reveal grain boundaries. For Custom Age 625+, the 15-

10-10 reagent (15 ml HCl, 10 ml acetic acid, 10 ml HNO3) was used to reveal the grain

boundaries. The samples were swabbed with the reagent until the surface appeared

dull; which could take up to 2 minutes.

4.4 Results

4.4.1 Hardness Testing

The results from the hardness testing are summarised below in Table 2

Table 2 Summary of Vickers Hardness for 718 and 625+

718 625+

AR 422 405

SA 175 200

HT1 305 272

HT2 216 245


79

It can be seen from Table 2, that the as-received gives the highest Vickers hardness

values for both 718 and 625+; 422 HV and 405 HV respectively. As expected due to

the lack of strengthening precipitates in the microstructure, the solution annealed

condition is the lowest hardness values for both alloys; 175 HV for 718 and 200 HV

for 625+. Heat Treatments 1 and 2 improve the hardness from solution annealed –

but it is the combined effect of the different precipitates which is giving the hardest

alloy. The resultant pattern for hardness pattern for the heat treatments is AR > HT2

> HT1 > SA.

4.1.2 718

Micrographs for the four heat treatment conditions described in Table 1 for IN718

after Kallings etching can be seen in Figure 1.

The Kallings etching revealed an austenitic grain structure for all heat treatments with

a grain size of ~85 μm 8 μm. From the micrographs above clear differences can be

a) b)

c) d)

Figure 1 Optical Micrographs of IN718 a) AR b)SA c)HT1 d)HT2


80

seen particularly with the HT2 compared with the other heat treatments. Figure 4d,

detailing the HT2 grain structure shows a high platelet-like content at the grain

boundaries which is known to be the δ-phase. The other phases, γ’ and γ”, along with

the carbides are difficult to pinpoint on the micrographs due to their small size. The

hardness testing, however, confirmed the success of the heat treatments.

4.1.3 625+

Figure 2 shows the SEM images of the 625+ after the 10-10-15 etching for the

different heat treatments. A similar grain structure to that observed for 718 can be

noticed. With the main difference being a smaller grain size (~45 μm 6 μm). Figure

3 is a high magnification image of 625+ Heat Treatment . The δ-phase is of the order

of μm and so can be observed within the S . Larger δ-phase precipitates are found

along the grain boundaries, whereas the smaller δ-phase precipitates can be found

dispersed within the matrix.


81

Figure 2 SEM images of 625+ a) AR b)SA c)HT1 d)HT2


82

Figure 3 High magnifcation SEM image of 625+ HT2

4.5 Discussion Both alloys 718 and 6 5 have an austenitic grain structure strengthened by the γ’,

γ” and a δ phases. Although not further investigated the literature shows the γ’ and

γ” are semi-coherent and the δ phase is BCC and thus incoherent with the matrix.

When predicting the effect of the microstructure on the corrosion behaviour, it is

thought the HT which is dominated by the δ phase would have the least corrosion

resistance. As the δ phase is incoherent with the matrix, it was assumed that

δ-phase precipitates – larger at grain boundaries; and smaller intragranular precipitates.


83

differences between the precipitate and the matrix would act as initiation points for

corrosion; in a similar way to MnS inclusions in stainless steels.

The solution annealed microstructure is predicted to have the highest corrosion

resistance. The SA heat treatment should give the alloy the most uniform

microstructure and distribution of precipitates should they be able to form. There are

no incoherencies in the matrix to act as initiation points.

4.6 Conclusions It can be observed from the resultant micrographs from the etchings, and the

hardness testing that there is a defined difference in microstructures between the

different heat treatments. Predictions have been made with regards to the corrosion

behaviour of the microstructures and their effects of the corrosion behaviour; with

HT2 expecting to have the least corrosion resistance, and SA the highest for both

alloys. The effects of the change in microstructure on the localised corrosion

behaviour are discussed in the following chapters.

4.7 References 1. Iannuzzi, M., Barnoush, A. & Johnsen, R. Materials and corrosion trends in

offshore and subsea oil and gas production. npj Mater. Degrad. 1, (2017).

2. Dul, I. Application and processing of nickel alloys in the aviation industry. Weld.

Int. 27, 48–56 (2013).

3. Zinkle, S. J. & Was, G. S. Materials challenges in nuclear energy. Acta Mater.

61, 735–758 (2013).

4. Reed, R. C. & Rae, C. M. F. Physical Metallurgy of the Nickel-Based Superalloys.


84

Physical Metallurgy: Fifth Edition 1, (Elsevier, 2014).

5. Krueger, D. D. Disk Applications. Direct 279–296 (1989).

6. Voort, G. F. V., Bowman, J. W. & Frank, R. B. Microstructural Characterization

of Custom Age 625 Plus Alloy. 489–498 (2012).

7. Acharya, V., Basa, D. K. & Murthy, G. V. S. Microstructural Characterization Of

Intermettalic Precipitates In Inconel 718 Alloy By DC Electrical Resistivity

Measurements. 2–9

Chapter 5; Crevice corrosion behaviour of Inconel 718 and Custom Age 625+ in sodium chloride solution

85

CHAPTER 5; CREVICE CORROSION BEHAVIOUR

OF INCONEL 718 AND CUSTOM AGE 625+ IN

SODIUM CHLORIDE SOLUTION

M. Keogh, T. Cook, R.Akid, D.Engelberg

5.1 Abstract The crevice corrosion behaviour of alloy 718 and alloy 625+ were investigated

electrochemically via potentiodynamic polarisation with specimens fitted with

crevice formers. These alloys were investigated due to their use as subsea

connections in the oil and gas industry, making their geometry one that is susceptible

to crevice corrosion. It was found that alloy 718 is susceptible to crevice corrosion at

temperatures as low as 30°C in de-aerated 3.5 wt% NaCl solution; whereas the 625+

does not undergo crevice corrosion until temperatures reach 60°C.

5.2 Introduction Ni-base superalloys are used extensively within the oil & gas, and aerospace

industries. Their high strength and superior corrosion properties make them a

desirable alloy in service.1 These alloys have good resistance to a wide variety of


86

corrosive environments, and as such are used within oil recovery pipelines. The main

function of these alloys is to act as connectors in the subsea well heads. They are

often subject to aggressive environments and the application of static and cyclic

stress. New sources of oil & gas require that the industry has to access the resources

from greater depths, meaning more aggressive environments for these alloys.

Well systems and pipelines are cathodically protected in order to improve their

operating lifetime. If the cathodic protection should fail for any reason, the Ni-base

alloys may be subject to localised corrosion, in the form of crevice corrosion and

pitting, due to the geometry of the components. The crevice corrosion is further

encouraged due to the chloride environment of seawater; an aggressive electrolyte

which can attack almost all common structural materials. In aerated, neutral pH

marine environments, the corrosion process is defined by two competing factors:

Chloride ion activity; which can lead to passive film breakdown.2

Dissolved oxygen; which acts to promote repair of the passive film.2

Both of these factors are detrimental to the localised corrosion resistance of Ni-base

alloys. Klapper et al record that 718 has a CCT of ≤ 5°C in similar conditions to that

discussed in this paper, although slightly harsher as their test solution contained FeCl3

and HCl.3 Hornus et al reported that Alloy 22 has a CCT between 30°C and 40°C,4 with

Mishra et al giving a crevice protection potential (ER.CREV) of -175 mVSCE5

for the same

alloy. Although there is little in the literature about 625+, Schmidt et al reported 625+

having a CCT of 40°C in the yellow death solution (4 wt % NaCl, 0.1 wt % Fe2(S04)3 +

0.01 M HCl).6 Alloy 625, the precursor to 625+ can also be used as a good standpoint

of what to expect from 625+ due to their similar compositions. Alloy 625 has a CCT of


87

between 30-35 °C, and 725, another nickel alloy has a CCT of 35°C.7 Values of 25°C to

50°C are expected in this work which is in keeping with other nickel superalloys with

similar compositions.

The purpose of this work was to investigate the difference in crevice corrosion

behaviour at different temperatures in sodium chloride solution between two alloys,

718 and 625+. The electrochemical behaviour of the alloys was studied using

potentiodynamic polarisations in order to obtain the Critical Crevice Temperature

(CCT).

5.3 Experimental Test specimens of Inconel 718 and Custom Age 625+ were cut using EDM (Electron

Discharge Machining) from annealed ingots into 20 mm x 20mm x 5mm cuboids and

with an 8mm hole in the centre. The chemical composition and heat treatment

histories are provided in tables 1 and 2, respectively. After heat treatment, the black

oxide layer was removed through SiC grounding.

For Alloy 718, the Kallings reagent (100 ml ethanol, 100 ml HCl, 5 g CuCl) was utilised

as the etchant of choice in order to reveal grain boundaries. The samples were


For Alloy 625+, the 15-10-10 reagent (15 ml HCl, 10 ml acetic acid, 10 ml HNO3) was

used to reveal the grain boundaries. The samples were swabbed with the reagent



88

Table 1 Chemical composition of 718 and 625+

Ni Cr Fe Nb+Ta Mo Al Ti

718 54.70 18.60 17.00 5.07 3.04 0.45 0.88

625+ 61.10 21.00 4.24 3.65 8.10 0.22 1.49

Table 2 Heat treatment of alloys

Alloy Heat Treatment History

718 Solution Anneal at 1021°C for 1 hour. Age hardened at 760°C - 8 hours. Air

cool.

625+ Proprietary heat treatment by Foroni.

Figure 1 shows the crevice former set-up used, based on the ASTM G488 which utilises

ceramic crevice formers made from Technox® 2000, a zirconia ceramic, which are

wrapped with PTFE tape and then attached to either side of the specimen with steel

washers, bolt and nut which secure the set-up. Plastic tubing is placed around the

bolt to ensure electrical isolation from the sample. The torque applied to was 1.5 N

m. The total surface area of the specimen was 9.74 cm2. The specimens were

successively ground to a finish of SiC paper 600, and then washed with acetone and

distilled water. Samples were allowed to air-dry for a minimum of 24 hours before

testing.


89

All the electrochemical tests were conducted in a one-litre three-electrode cell. The

reference electrode used was a 3M Ag/AgCl, and a platinum counter electrode.

Nitrogen was purged for a minimum of 30 minutes to ensure deaeration before

testing, and there was a continuous flow of nitrogen throughout testing. The

temperature of the solution was controlled by immersing the cell in a water bath

which was kept at constant temperature, and monitored with a thermometer placed

inside the cell.

Tests were performed in a 3.5 wt% NaCl solution. The open circuit potential (ECORR)

was monitored for 5 minutes in the deaerated conditions. Potentiodynamic

polarisation was used to determine the critical crevice temperature and the crevice

susceptibility behaviour of the alloys. The scans were started at the OCP in the anodic

direction at a rate of 1 mV/s until an end point of +0.5 V or +1.5 V vs Ag/AgCl. The

Crevice Potential (ECREV) is defined at the point when then current density remains at

least 20 μA above the passive current density.

Figure 1 Schematic of a sample with attched crevice formers

1 cm


90

A modified THE method was then used to determine the protection potential (ER.CREV).

Specimens were characterised after testing with both a Scanning Electron

Microscopy (SEM) and a Laser Scanning Confocal Microscope (LSCM).

5.4 Results

5.4.1 Microstructure

Figure 2 shows the microstructure of the two alloys investigated in this work. Both

alloys are precipitation hardened, consisting of an austenitic matrix with the face-

centred-cubic (FCC) structure. There are a number of secondary phases, including γ’,

γ” and δ and the volume and distribution of each can change the materials

performance9.

It is expected that these alloys will contain the precipitate phase γ’ which follows the

FCC crystal structure and is thus coherent with the matrix, is the prevailing

strengthening precipitate.10 The γ” phase may also be present has a body centred

tetragonal structure, but remains semi-coherent with the matrix and remains a major

part in the strengthening mechanism of the alloys.10 The δ phase is the least coherent

with the matrix as it has an orthorhombic structure and has a platelet-like

Figure 2 Microstructure; Optical Micrograph a) 718 and SEM image of b) 625+

a) b)


91

morphology often found at the grain boundaries.11 Due to its incoherency with the

matrix it could act as a crack initiation site, or a pitting initiation site, both reducing

fracture toughness12 and promoting hydrogen embrittlement.13

5.4.2 Corrosion behaviour of AR-718

Figure 3 shows the potentiodynamic polarisation curves for 718 in 3.5 wt% NaCl

solution at increasing temperatures. Curves at room temperature and 30°C show a

wide passive range up to 0.6V where the current increases due to transpassivity and

again at 1.2 V due to oxygen evolution. There is a clear difference shape of the curves

30°C and 40°C as highlighted with the orange box. The sharp increase in current

density at 40°C seen around 0.4 V indicates that a localised corrosion event has taken

place, and is the boundary of the Critical Crevice Temperature (CCT). It was

confirmed by visual inspection of the sample at 40°C that crevice corrosion had taken

place under the occluded regions of the crevice former as shown in Figure 4.


92

Figure 3 Potentiodynamic Polarisation Curves of Inconel 718 in deaerated 3.5 wt%

NaCl solution

Figure 4 comparison of 718 speciemens after corrosion testing at a) room

temperature showing only transpassive film b) 40°C showing crevice corrosion

and transpassive film c) 50°C crevice corrosion only

Room Temperature

40°C 50°C

718 AR


93

Figure 4 a) and b) also shows a black nickel-oxide film as a result of the transpassivity

that occurs when taking the polarisation curves to high potentials of 1.5 V. The film

is not present on Figure 4c) as the potential finish point was lower at 0.5V.

Figure 5 shows SEM images of a post-immersion 718 specimens at room temperature

and 50°C. Figure 5a/b) shows the cracked transpassive nickel-oxide film, but no sign

of crevice corrosion. Figure 5c/d) shows crevice corrosion has taken place and that

the attack follows an intergranular pathway.

Figure 5 SEM images of 718 after immersion a/b) at room temperature c/d) at

50°C

a)

b)

c) d)

a)


94

Figure 6 shows the modified THE results for 718 at room temperature and 40°C; with

6a) shown as potential and current density with respect to time, and b) the resultant

polarisation curves when current is plotted as a function of potential. The polarisation

curves show a difference in shape due to the sample at 40°C undergoing crevice

corrosion, whereas the sample at room temperature did not suffer crevice attack;

which was confirmed with visual observation and SEM images (Figure 7).


95

Figure 6 Polarisation curves from modified THE experiment for 718 at room

temperature and 40°C


96

5.4.3 Corrosion behaviour of AR-625+

Figure 7 shows the potentiodynamic polarisation curves for 625+ at increasing

temperatures. Curves show a shallow passive area followed by current increases due

to transpassivity and oxygen evolution. There is a clear difference shape of the curves

50°C and 60°C as highlighted in Figure 7 with an orange box. The sharp increase in

current density at 60°C seen around 0.3V indicates that a localised corrosion event

has taken place, and is the boundary of the Critical Crevice Temperature (CCT). It was

confirmed by visual inspection of the sample at 60°C that crevice corrosion had taken

place under the occluded regions of the crevice former as shown in Figure 8.

Additional surface defects can also be seen in Figure 8a. The black film present is a

transpassive nickel oxide film formed when experiments are taken to high potentials

of 1.5 V.


97

Figure 7 Potentiodynamic Polarisation Curves of Custom Age 625+ in deaerated

3.5 wt% NaCl solution

Figure 8 Photographs of 625+ specimens after electrochemical testing at a) 30°C b)

70°c and c) 90°C

2D Graph 2

E / V v Ag/AgCl

-0.6 -0.4 -0.2 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6

i /

mA

cm

-2

-5

-4

-3

-2

-1

0

1

2

30

40

50

60

70

90

30°C 70°C 90°C

625+ AR


98

Figure 9 shows SEM images of a post-immersion 625+ specimen at 90°C where,

similar 718 shows the crevice corrosion has taken place and that the attack follows

an intergranular pathway.

Figure 9 SEM images of 625+ at 90°C showing crevice corrosion

Figure 10 shows the resultant polarisation from the modified THE results when

current is plotted as a function of potential. The polarisation curves show a difference

in shape due to the sample at 60°C undergoing crevice corrosion, whereas the sample

at room temperature did not suffer crevice attack, which was confirmed by visual

inspection.

a) b)


99

Figure 10 Resultant polarisation curves from modified THE experiment of 625+ at

room temperature and 60°C

5.5 Discussion The crevice corrosion resistance of Alloys 718 and 625+ in deaerated 3.5 wt% NaCl

solution was investigated through electrochemical testing at temperatures ranging

between room temperature and 90°C. Alloy 625+ showed a greater resistance to

crevice corrosion, than 718; both in the post-immersion sample inspections and

through the electrochemical behaviour.

Custom Age 625+ appears to exhibit a first crevice potential at 60°C, indicating it is

more resistant to crevice corrosion than Inconel 718 which crevices at 40°C as shown

in Figure 9. It can be seen from Figure 11, showing ETRANS and ECREV as a function of

temperature, that the interchange between just transpassivity and both crevice and

transpassivity comes between the temperatures of 30°C and 40°C for 718, and 50°C


100

and 60°C for 625+, indicating that the CCT also falls within these ranges. As the

microstructures between these two alloys are similar with regards to precipitates,

the difference in corrosion behaviour is attributed to differences in alloy composition.

The main difference in composition between the two alloys is an increased Mo

content for alloy 625+ when compared to 718. Molybdenum additions have been

known to restrict the initiation of crevices and reduce propagation rates14. Although

one mechanism has not been agreed for how the molybdenum can affect the crevice

initiation and propagation, it is thought to be due to the production of a MoO42- ion

which can provide a protective oxide layer.14

Both alloys show that the crevice corrosion follows an intergranular attack pathway,

which indicates the precipitates found at the grain boundaries may have a role in

facilitating the corrosion pathway. The gamma precipitates have been known to be

more electrochemically active than the matrix, and congregate at grain boundaries15.

All the scans for 625+, prior to the onset of crevice corrosion yielded an anodic peak

at 0.4 V. In the literature, this has been previously attributed to the transpassive

dissolution of chromium as similar peaks were observed by Klapper et al16. The

dissolution of chromium phenomenon may be able to explain the additional visual

surface defects, as selective dissolution of chromium, observed on the sample surface

which are unrelated to the crevice corrosion.

Both alloys show a similar trend with the increase in temperature causing a decrease

in the ECREV value as summarised in Table 3; indicating that a higher temperature

allows for an easier initiation and propagation of crevice behaviour. Figure 11

demonstrates than there is a linear relationship between temperature and


101

breakdown potentials, with the increasing temperature causing a linear decrease in

ECREV. Evidence that temperature affects crevice corrosion behaviour is further

supported by Figures 5 and 9 where the extent of crevice corrosion is seen to become

more severe for specimens immersed at higher temperatures, indicating that

temperature can also have an effect on the severity of crevice corrosion.

Figure 11 ECREV and ETRANS as a fuction of temperature for 718 and 625+ with dashed lines indicating the tranpassive to crevice transition

Table 3 Summary of ECREV values for 718 and 625+

Temp °C 40 50 60 70 90

718 0.28 V 0.08 V -0.19 V -0.21 V /

625+ / / 0.35 V 0.31 V 0.19 V


102

These results from electrochemical testing of the alloys 718 and 625+ in this work

indicate that it may be possible to eliminate the possibility of crevice corrosion

occurring by operating at a sufficiently low temperature; i.e. one that is below the

CCT.

5.6 Conclusions From the results obtained regarding the effect of temperature on the localised

crevice corrosion, and the effects of changing the start point of the potentiodynamic

scans of the Ni-base superalloys 718 and 625+, the following conclusions can be

drawn:

• Alloy 718 is more susceptible to crevice corrosion than 625+ as it experienced

crevice corrosion more readily at temperatures less than 50oC. The crevice

potentials for 718 were lower than those for 625+ for the first instance of

crevice corrosion occurring (0.28V(SCE) vs 0.35V(SCE) respectively).

• The Critical Crevice Temperature (CCT) for 718 was recorded as being in the

range of 30-40°C with crevices being visually observed at all temperatures

investigated above this range.

• CCT for 625+ has been recorded as being in the range of 50°C-60°C.

5.7 References 1. Rhodes, P. R. Environment-assisted cracking of corrosion-resistant alloys in oil

and gas production environments: A review. Corrosion 57, 923–966 (2001).

2. Ahn, T., Jung, H., Shukla, P., He, X. & Introduction, I. Criteria For Crevice

Corrosion In Concentrated Chloride Solutions Nuclear Systems. (2012).


103



296–305 (2017).

4. Hornus, E. C., Rodríguez, M. A., Carranza, R. M., Giordano, C. M. & Rebak, R. B.

Effect of Environmental Variables on Crevice Corrosion Susceptibility of Ni–Cr–

Mo Alloys for Nuclear Repositories. Procedia Mater. Sci. 8, 11–20 (2015).



6. Schmidt, N. B., DeBold, T. a. & Frank, R. B. Custom age 625® plus alloy-A higher

strength alternative to alloy 625. J. Mater. Eng. Perform. 1, 483–488 (1992).

7. McCoy, S., Hereford, U. & Puckett, B. High performance age-hardenable nickel

alloys solve problems in sour oil and gas service. Balance (2002).



Solution. 1–13 (2011).

9. Engelberg, D. L., Cottis, R. A., Sherry, A. H. & Marrow, T. J. Project Report

Literature Review : Stress Corrosion Cracking of Inconel 718 in PWR

Environments. (2006).





104





Mater. 3, 335–344 (2002).

14. Lillard, R. S., Jurinski, M. P. & Scully, J. R. Crevice Corrosion of Alloy-625 in

Chlorinated Astm Artificial Ocean Water. Corrosion 50, 251–265 (1994).

15. Hwang, I. S. Electrochemistry of Multiphase Nickel-Base Alloys in Aqueous

Systems. J. Electrochem. Soc. 136, 1874 (2006).



70, 899–906 (2014).

Chapter 6; Effects of Microstructure on localised corrosion behaviour of Inconel 718 and Custom Age 625+

105

CHAPTER 6; EFFECTS OF MICROSTRUCTURE ON

LOCALISED CORROSION BEHAVIOUR OF INCONEL 718 AND CUSTOM AGE

625+

M. Keogh, T. Cook, D.Martelo-Guarin, R. Akid, D. Engelberg

6.1 Abstract

The crevice corrosion behaviour properties of the alloys 718 and 625+ have been

investigated. Three microstructures were tested in order to reveal microstructural

impacts on crevice corrosion; in particular the roles of the precipitates γ’, γ” and δ

phase. Each heat treatment of the alloy underwent potentiodynamic polarisation at

increasing temperatures to establish the Critical Crevice Temperature (CCT) and the

crevice potential (ECREV). For both alloys it was found that the solution annealed

condition was the most susceptible to crevice corrosion, as it had the lowest CCT and

ECREV values, indicating the strengthening precipitates play a role in the corrosion

resistance behaviour.


106

6.2 Introduction

As nickel alloys have become more sophisticated in their design, the severity of the

environments that they are utilised in has increased.1 Many of these environments,

particularly marine, can be detrimental to the corrosion behaviour of the alloys due

to the presence of chloride ions. Chloride ions can cause pitting and crevice corrosion

which can act as initiation points for Stress Corrosion Cracking.2

It is important to be able to understand the role that the microstructure plays on the

corrosion properties of an alloy as there may be a heat treatment available which can

maximise the microstructure to give the greatest corrosion resistance.

In Nickel-alloys there are 3 main types of precipitate which can occur; γ’, γ” and δ

phase. The precipitate phase γ’, i3((Ti,Al)Nb), follows the FCC crystal structure, is

coherent with the matrix is and is the prevailing strengthening precipitate.3

In comparison, γ” has a body centred tetragonal structure, but remains semi-

coherent with the matrix. The morphology of γ” is usually disk-shaped, and remains

a major part in the strengthening mechanism of the alloy with a compositions of

Ni3(Nb(Ti,Al)).3

The δ phase incoherent with the matrix as it has an orthorhombic structure and has

a platelet-like morphology often found at the grain boundaries.4 Due to its

incoherency with the matrix it can act as a crack initiation site, a pitting initiation site,

both reducing fracture toughness5 and promote hydrogen embrittlement.6 Due to

these factors it was predicted that Heat Treatment which contained only δ phase

would be the least corrosion resistant.


107

In this work, three heat treatments were developed to study the effects of the

precipitates on the corrosion resistance of the alloys 718 and 625+. Each heat

treatment highlights different precipitates. The microstructures can be seen in

Chapter 4. The solution anneal dissolves all the precipitates into the matrix, and was

expected to have the greatest corrosion resistance as there would be no precipitates

present to act as initiation points for corrosion. Heat Treatment 1 reveals the primary

strengthening precipitates y’ and y”, and Heat Treatment gives the needle-like delta

phase and was expected to have the least corrosion resistance as the delta phase in

incoherent with the matrix.

6.3 Experimental

Test specimens of Inconel 718 and Custom Age 625+ were cut using EDM (Electron


with an 8mm hole in the centre. The chemical composition and heat treatments are

provided in tables 1 and 2, respectively.

For Alloy 718, the Kallings reagent (100 ml ethanol, 100 ml HCl, 5 g CuCl) was utilised

as the etchant of choice in order to reveal grain boundaries. The samples were


For Alloy 625+, the 15-10-10 reagent (15 ml HCl, 10 ml acetic acid, 10 ml HNO3) was

used to reveal the grain boundaries. The samples were swabbed with the reagent



108


Ni Cr Fe Nb+Ta Mo Al Ti

718 54.70 18.60 17.00 5.07 3.04 0.45 0.88

625+ 61.10 21.00 4.24 3.65 8.10 0.22 1.49


Designation Heat treatment Expected microstructure

Solution

annealed

(SA)

Annealed at 1040oC for 1 hour.

Air cool.

Solution annealed (carbides and

nitrides)

HT1 Annealed at 1040oC for 1hr. Air

cool. Age hardened at 650°C for

10 hours. Air cool.

Precipitation hardened (γ` + γ``

phases – dominant precipitate

γ`)

HT2 Annealed at 1040oC for 1hr. Air

cool. Age hardened at 900oC for

20 hours. Air cool.

Precipitation hardened

(dominant precipitate δ phase)

Figure 1 shows the crevice former set-up used, based on the ASTM G487 which utilises

ceramic crevice formers made from Technox® 2000, a zirconia ceramic, which are

wrapped with PTFE tape and then attached to either side of the specimen with a

torque of 1.5 N m, and the electrical connection isolated. The specimens were

successively ground to a finish of SiC paper 600, and then washed with acetone and


109

distilled water. Samples were allowed to air-dry for a minimum of 24 hours before

testing.


All the electrochemical tests were conducted in a one-litre three-electrode cell with

a used a 3M Ag/AgCl reference electrode, and platinum counter electrode. Nitrogen

was purged for a minimum of 30 minutes prior to testing, and there was a continuous

flow of nitrogen throughout testing. The temperature of the solution was controlled

by immersing the cell in a water bath and monitored with a thermometer placed

inside the cell.

Tests were performed in a 3.5 wt% NaCl solution. The open circuit potential (ECORR)

was monitored for 5 minutes, from which scanning started in the anodic direction at

a rate of 1 mV/s until an end point of +0.5 V or +1.5 V vs Ag/AgCl. The Crevice

Potential (ECREV) is defined as the point where the current density remains above 4

μA cm-2. Specimens were characterised after testing with both a Scanning Electron

Microscopy (SEM) and a Laser Scanning Confocal Microscope (LSCM).

1


110

6.4 Results

6.4.1 718

Figure 2 shows the potentiodynamic polarisation curves for the three heat

treatments; SA, HT1 and HT2; for 718 in 3.5 wt% NaCl solution at increasing

temperatures. The curves at room temperature have a passive range up to

approximately 0.7 V where the current increases due to transpassivity and again at

1.2 V due to oxygen evolution.

Figure 2 Potentiodynamic Polarisation Curves of solution-annealed 718 in

deaerated 3.5 wt% NaCl solution a) solution-annealed; b) Heat Treatment 1; c)

Heat Treatment 2


111

In Figure 2a) for the solution annealed condition there is a clear difference shape of

the curves at room temperature and 30°C. The sharp increase in current density at

30°C seen around -0.1 V indicates that a localised corrosion event has taken place,

and is the boundary of the Critical Crevice Temperature (CCT). It was confirmed by

visual inspection of the sample at 30°C that crevice corrosion had taken place under

the occluded regions of the crevice former as shown in Figure 3. Figure 3 a) shows

only evidence of film as a result of the transpassivity that occurs when taking the

polarisation curves to high potentials of 1.5 V. The film is not present on Figure 3 b)

as the potential finish point was lower at 0.5 V, but crevices are clearly evident.

A similar result was found for Heat Treatment 1, with the boundary of the CCT also

being 30°C (Figure 2b), and crevice corrosion being observed in the post immersion

analysis (Figure 3d).

For Heat Treatment 2, however, the CCT has been elevated to 40°C. It was confirmed

by visual inspection of the sample at 40°C that crevice corrosion had taken place

under the occluded regions of the crevice former as shown in Figure 3f). Figure 5b) is

the same sample shown as an SEM image.


112

Figure 3 Photographs of 718 specimens after electrochemical testing at

a) SA RT b) SA 40°C c) HT1 30°C d) HT1 70°C e) HT2 RT f) HT1 50°C

Figure 4 exemplifies ETRANS and ECREV as a function of temperature for all heat

treatment conditions of 718. Here the transition between transpassive corrosion and

crevice corrosion can be easily identified.


113

Figure 4 ECREV and ETRANS as a function of temperature for 718 a) Solution

Annealed; b) Heat Treatment 1; c) Heat Treatment 2

Figure 5 SEM images of a post-immersion crevice from Heat Treatment 1 718

specimen at 50°C.


114

6.4.2 625+


temperatures for all heat-treatments. Curves that are extended to 1.5 V show a

shallow passive area followed by current increases due to transpassivity and oxygen

evolution.

For solution annealed 625+ (Figure 6a) there is a clear difference shape of the curves

30°C and 40°C which is the lowest CCT observed for 625+. The sharp increase in

current density at 40°C seen around 0.1V indicates that a localised corrosion event

has taken place. When compared to both 625+ Heat Treatments 1 and 2(Figure 6b/c),

however, the Critical Crevice temperature can be seen to occur between 40°C and

50°C. The CCT data, with ECREV and ETRANS values, is summarised in Figure 7.


115

Figure 6 Potentiodynamic Polarisation Curves of Custom Age 625+ in deaerated

3.5 wt% NaCl solution a) Solution-Annealed; b) Heat Treatment 1; d) Heat

Treatment 2

Figure 7 ECREV and ETRANS a function of temperature for heat treatments of 625+ a)

Solution Annealed; b) Heat Treatment 1; c) Heat Treatment 2

After samples have undergone electrochemical testing they are photographed, and

observed using SEM. Figure 8 shows example of 625+ specimens for all heat

treatments; SA, HT1, and HT2; after corrosion testing. Images on the left show only a

transpassive film (orange circle), whereas those on the right which have exceeded

the CCT have undergone crevice corrosion (blue circle).


116

Figure 8 Photographs of 625+ specimens after electrochemical testing at a) SA RT

b) SA 60°C c) HT1 RT d) HT1 50°C e) HT2 RT f) HT2 70°C

Figure 9 shows SEM images of a post-immersion 625+ specimen at 60°C. The image

on the left shows the whole crevice and on the right highlights the corner of the

crevice, where it appears from these images that the attack follows an intergranular

pathway.


117

Figure 9 SEM images of a crevice from Heat Treatment 1 625+ after immersion at

60°C

6.5 Discussion The effects of microstructure on the crevice corrosion resistance of Alloys 718 and

625+ in deaerated 3.5 wt% NaCl solution were investigated through electrochemical

testing at temperatures ranging between room temperature (range of 20°C to 28°C)

and 60°C. Alloy 625+ showed a greater resistance to crevice corrosion, for all heat

treatments, than 718 as summarised in Figure 10; which also includes values from

Chapter 4 about the as-received heat treatment conditions. The data shows that the

as-received heat treatments for both alloys maintain to be the heat treatments with

the highest corrosion resistance for each respective alloy. The Critical Crevice

Temperature for AR 718 is 40°C with an ECREV of 0.369 V; the CCT for AR 625+ is 60°C

with an ECREV of 0.308 V. The CCT for 718 and 625+ are similar to those quoted by

McCoy et al for 625 being between 30-35 °C and for 725 being 35°C8. The CCT value

for 625+ is higher than that of 625 which is to be expected to due to 625+ being


118

engineered to have superior qualities when compared to 625. The CCT value for 725

is similar to that of 718 due to their similar compositions.

The solution annealed condition, again for 718 and 625+, shows to give the least

localised corrosion resistance with CCTs of 30°C and 40°C respectively. Although the

CCT for the solution annealed condition of also shared with Heat Treatment 1 of 718,

the ECREV is significantly lower; 0.055 V for SA compared to 0.395 V for HT1. It was not

expected that the solution annealed condition had the lowest corrosion resistance. A

possible explanation for the behaviour of the solution annealed conditions is that it

experiences a great rate of stress when in the crevice set-up. It is know that increased

stress on materials can accelerate the corrosion rate, as often found in stress

corrosion cracking.9 As the solution annealed condition has the lowest Vickers

hardness values it is susceptible to higher strain levels when compared to other heat

treatment conditions.10

It is however, the combination of precipitates which provide the greatest corrosion

resistance, with δ phase seemingly the most contributing factor, despite it being

incoherent with the matrix. It was hypothesised that the heat treatments containing

precipitates, especially Heat Treatment 2 which contains the incoherent delta phase,

would have a reduced corrosion resistance. The basis for this hypothesis was thinking

that the precipitates would act as initiation points for corrosion to occur; similar to

MnS inclusions in stainless steel.11

In conditions where the y’ and y” are the most abundant the microstructure is

revealed. Hwang et al found that the gamma-phases can be more electrochemically

active than the matrix thus acting as a localised anode.12 As these precipitates can be


119

found at the grain boundaries, it may rationalise the intergranular pathway seen on

SEM images during post-immersion analysis.

The corrosion resistance follows the pattern AR > HT2 > HT1 > SA for both alloys.

Figure 10 A graph summarising the Critical Crevice Temperature and ECREV values

for all heat treatments of 718 and 625+

Some scans for 625+, prior to the onset of corrosion, yielded an anodic peak at 0.4 V.

In the literature, this has been previously attributed to the transpassive dissolution

of chromium as similar peaks were observed by Klapper et al.13 This phenomenon

may be able to explain the additional visual surface defects, unrelated to the crevice

corrosion, observed on some samples as selective dissolution of chromium.


120

6.6 Conclusions From the results obtained regarding the effect of temperature on the localised

crevice corrosion, and the effects of changing the start point of the potentiodynamic

scans of the Ni-base superalloys 718 and 625+, the following conclusions can be

drawn:

Alloy 718 is more susceptible to crevice corrosion than 625+ for all heat

treatments as it experienced crevice corrosion more readily; both with

temperature and ECREV.

The corrosion resistance follows the pattern HT2 > HT1 > SA for both alloys.

6.7 References

1. Pint, B. A. Performance of Wrought Superalloys in Extreme Environments. 740,

165–178 (2018).

2. Ma, F.-Y. Corrosive Effects of Chlorides on Metals. Pitting Corros. (2012).








Mater. 3, 335–344 (2002).


121



Solution. 1–13 (2011).

8. McCoy, S., Hereford, U. & Puckett, B. High performance age-hardenable nickel

alloys solve problems in sour oil and gas service. Balance (2002).

9. Christman, T. K. Relationships Between Pitting, Stress, and Stress Corrosion

Cracking of Line Pipe Steels. CORROSION 46, 450–453 (1990).

10. Busby, J. T., Hash, M. C. & Was, G. S. The relationship between hardness and

yield stress in irradiated austenitic and ferritic steels. J. Nucl. Mater. 336, 267–

278 (2005).

11. Maciejewski, J. The Effects of Sulfide Inclusions on Mechanical Properties and

Failures of Steel Components. J. Fail. Anal. Prev. 15, 169–178 (2015).





70, 899–906 (2014).

Chapter 7; Effects of Chloride concentration on Crevice corrosion behaviour

122

CHAPTER 7; EFFECTS OF CHLORIDE

CONCENTRATION ON CREVICE CORROSION

BEHAVIOUR

7.1 Abstract The localised corrosion behaviour properties of two nickel alloys, 718 and 625+, were

investigated after undergoing heat treatments to change the microstructure. These

alloys were chosen as they are commonly used in the petroleum industry for sub-sea

conditions. Each heat treatment of the alloy underwent potentiodynamic

polarisation at increasing temperatures to establish the Critical Crevice Temperature

(CCT) in 0.1M NaCl and 1M NaCl solution. Results were compared to a previous study

which took place in 0.6 M (3.5 wt%) NaCl solution. The effects of chloride was

investigated as ocean chloride concentrations can differ worldwide, and these alloys

are being used in other industries where chloride concentrations may differ to those

within the oil and gas industry. It was found that a reduction in chloride concentration

to 0.1M can increase the CCT when compared to 0.6 M (3.5wt%), and an increase

in chloride concentration can decrease the CCT when also compared to 0.6M; but

both have not occurred for the same alloy and heat treatment.


123

7.2 Introduction Subsea oil and gas production wells are in use all over the globe with their locations

ranging from the North Sea to the Gulf of Mexico.1 The different locations each pose

their own challenges; not only with temperature variations but also with differences

in salinity.2 The differences in salinity can be affected by different temperatures3 and

depths.4

Despite the possible differences in the environmental conditions, the components

used in sub-sea wells are uniform across deployed fleet globally;5 meaning that the

alloys utilised must be able to withstand the differing conditions.

Chloride content is of interest due to the possibility of chloride initiating localised

corrosion if the material is susceptible;6 which can then possible act as an inlet for

hydrogen, causing hydrogen embrittlement;7 or stress corrosion cracking.8 All these

degradation mechanisms can be a possible cause of the premature failure of

components.9 Nickel alloys are commonly used as connectors and fasteners10 for the

sub-sea well heads which are responsible for maintaining and monitoring the

pressures of the pipelines.11

It has become widely known that Ni-Superalloys are susceptible to localised

corrosion.12 The effects of temperature, and microstructure have previously been

discussed in Chapter 5; this chapter will focus on the effects that chloride

concentration has on the localised corrosion behaviour of alloy 718 and 625+. The

temperature range surrounding the reported CCT from the previous chapters have

acted as a starting point for the chloride concentration investigations.


124

Two chloride concentrations [0.1 M] and [1M] were tested, and the results from

these were compared to those in Chapter 5, to provide an indication of the effect of

[Cl-] on the localised corrosion behaviour.



with an 8 mm hole in the centre. The chemical composition and heat treatment

histories are provided in Tables 1 and 2, respectively.



718 54.70 18.60 17.00 5.07 3.04 0.45 0.88

625+ 61.10 21.00 4.24 3.65 8.10 0.22 1.49


Designation Heat treatment Expected microstructure




A distribution of y’. y”, δ phase and carbides.



Solution annealed (carbides and nitrides dissolved in matrix)

HT1 Annealed at 1040oC for 1hr. Air cool. Age hardened at 650°C for 10 hours. Air cool.

Precipitation hardened (γ` + γ`` phases – dominant precipitate γ – disc morphology`)

HT2 Annealed at 1040oC for 1hr. Air cool. Age hardened at 900oC for 20 hours. Air cool.

Precipitation hardened (dominant precipitate δ phase found at grain boundaries with platelet morphology)


125

Figure 1 shows the crevice former set-up used, based on the ASTM G4813 which

utilises ceramic crevice formers made from Technox® 2000, a zirconia ceramic, which

are wrapped with PTFE tape and then attached to either side of the specimen with

steel washers, bolt and nut which secure the set-up. Plastic tubing is placed around

the bolt to ensure electrical isolation from the sample. The torque applied to was 1.5

N m. The total surface area of the specimen was 9.74 cm2.

The specimens were successively ground to a finish of SiC paper 600, and then

washed with acetone and distilled water. Samples were allowed to air-dry for a

minimum of 24 hours before testing.

All the electrochemical tests were conducted in a one-litre three-electrode cell. The

reference electrode used was a 3M Ag/AgCl, and a platinum counter electrode.

Nitrogen was purged for a minimum of 30 minutes to ensure deaeration before

1cm



126



which was kept at constant temperature, and monitored with a thermometer placed

inside the cell.

Tests were performed in a either 0.1M or 1M NaCl solutions. The open circuit

potential (ECORR) was monitored for 5 minutes in the deaerated conditions.

Potentiodynamic polarisation was used to determine the critical crevice temperature

and the crevice susceptibility behaviour of the alloys. The scans were started at the

OCP in the anodic direction at a rate of 1 mV/s until an end point of +0.5 V or +1.5 V

vs Ag/AgCl.

7.4 Results

7.4.1 Electrochemical Results for 718

Figure 2 shows the potentiodynamic polarisation curves for 718 at increasing

temperatures for all heat-treatments in 0.1M NaCl solution. Curves that are extended

to 1.5 V show a shallow passive area followed by current increases due to

transpassivity and oxygen evolution. For as-received 718 (Figure 2a) there is a clear

difference shape of the curves 40°C and 50°C. The sharp increase in current density

at 50°C seen around 0.3V indicates that a localised corrosion event has taken place.

When compared to 718 Solution Annealed (SA) (Figure 2b) however, the increase in

current relating to a crevice corrosion event can be seen to present in the curve at

room temperature at around 0.1 V. Heat Treatment 1 (Figure 2c) has clear changes

in curve geometry between room temperature and 30°C, whereas for Heat

Treatment 2 (Figure 2d) has the CCT occurring between 40°C and 50°C.


127

Figure 2 Potentiodynamic polarisation curves of 718 in 0.1M NaCl solution at

increaseing temperatures for different heat treatments a) As-Recieved b) Solution

Annealed c) Heat Treatment 1 d) Heat Treatment 2

Figure 3 shows the potentiodynamic polarisation curves for 718 at increasing

temperatures for all heat-treatments in 1M NaCl solution. Curves that are extended


transpassivity and oxygen evolution. For as-received 718 (Figure 3a) there is a clear

difference shape of the curves 30°C and 40°C. The sharp increase in current density

at 40°C seen around 0.3V indicates that a localised corrosion event has taken place.




128

room temperature. Both Heat Treatment 1 and Heat Treatment 2 (Figure 3 c/d) have

their CCT occurring between 40°C and 50°C.

Figure 3 Potentiodynamic polarisation curves of 718 in 1M NaCl solution at



7.4.2 Electrochemical Results for 625+


temperatures for all heat-treatments in 0.1M NaCl solution. Curves that are extended


transpassivity and oxygen evolution. For as-received 625+ (Figure 4a) the curves for


129

40 °C and 50 °C do not follow the expected path. The curves have high initial current

densities, and do not have a clear passive range. If a crevice were to occur close to

the Open Circuit Potential (OCP) then an immediate increase in current density would

be expected. The current densities measured here, however, are much higher than

for the other heat treatments when a crevice has initiated, and for the as-received at

60 °C. In addition, upon inspection of the samples there was no crevice corrosion

observed. It is hypothesised that the high initial current could be due to corrosion of

the electrode wire due to an improper seal between the plastic tubing protecting the

wire and the sample. The experiments for 625+ AR would have to be repeated for

them to be a reliable source. At the current time a CCT of 60°C has been assumed.



room temperature at 0.2V. Heat Treatment 1 has its CCT occurring between 40°C and

50°C, but it occurs 10°C below this for Heat Treatment 2.


130

Figure 4 Potentiodynamic polarisation curves of 625+ in 0.1M NaCl solution at




temperatures for all heat-treatments in 1M NaCl solution. Curves that are extended


transpassivity and oxygen evolution. For as-received 625+ (Figure 5a) there are

similar anomalous results for 50°C and 60°C. A CCT of 40°C has been assumed. When

compared to 625+ Solution Annealed (Figure 5b) however, the increase in current

relating to a crevice corrosion event can be seen to present in the curve at room


131

temperature. Both Heat Treatment 1 and Heat Treatment 2 (Figure 6 c/d) have their

CCT occurring between 40°C and 50°C.

Figure 5 Potentiodynamic polarisation curves of 625+ in 1M NaCl solution at



Figure 6 shows comparisons of CCT for 718 when compared to 625+ across all heat

treatments and experimental chloride concentrations. 718 has lower CCT for all Heat

Treatments and [Cl-] when compared to 625+. For both alloys the as-received

condition has the highest CCT, so the highest resistance to crevice corrosion. The

solution annealed heat treatment is unaffected by the change in chloride

concentration; its CCT remains constant at 30°C for 718 and 40°C for 625+. For Heat


132

Treatment 1, the CCT decreases for 718 in 1 M NaCl solution, but remains constant

for 625+. Heat Treatment 2 has an increase in CCT from 40 °C to 50 °C with the

decrease in [Cl-] from 0.6M to 0.1M for 718, but the same is not observed for 625+.

Instead the CCT for HT2 remains at 50 °C for 0.1 M and 0.6 M, but decreases to 40 °C

when chloride concentration increased to 1 M.

7.4.3 Post Immersion Sample Analysis

After samples have undergone electrochemical testing they are photographed, and

observed using SEM. Figure 7 shows samples from 718 after immersion. The top row

shows 718 samples after immersion in 0.1M NaCl, in the order a) AR, b) SA, c) HT1

and d) HT2; after immersion at 50°C. Crevices are visible on all samples. The bottom

row follows the sample order also shown are after immersion at 50°C. All samples in

the bottom row underwent crevice corrosion. The solution annealed sampled after

immersion in 1M NaCl at 50°C was the most attacked as it has the highest number of

crevices.

2D Graph 1

Molar concentration of NaCl solution

0.1 M 0.6 M 1M

Tem

pera

ture

(oC

)

10

20

30

40

50

60

70

718 AR

718 SA

718 HT1

718 HT2

2D Graph 3

Molar concentration of NaCl solution

0.1 M 0.6 M 1M

Tem

pera

ture

(oC

)

10

20

30

40

50

60

70

625+ AR

625+ SA

625+ HT1

625+ HT2

Figure 6 Summary graphs showing CCT for 718 and 625+ across all heat treatments

and chloride concentrations tested.


133

Figure 8 shows SEM images for two 718 samples after immersion at 50°C in 1M NaCl

for HT1 and HT2. Both images show clear crevices, with indicators of intergranular

attack as the microstructure has been releveled in both; especially clear for Heat

Treatment 2 (Figure 8b) where the needle-like delta phase has been exposed.

The SEM images also reveal that the crevice attack is not uniform beneath the crevice

former. The attack appears to be more severe at the edges of the crevices. Similar

attack pathways were seen for the other microstructures and temperatures.

718 AR 718 SA 718 HT1 718 HT2

0.1 M

50°C

1M

50°C

Figure 7 Photographs of 718 specimens after electrochemical testing at 50°C. Top Row (left to right) 0.1M AR SA HT1 HT2; Bottow Row 1M (left to right) 0.1M AR

SA HT1 HT2 b) 0.1M SA 50°C c) 0.1M HT1 RT d) 0.1M HT2 RT e) 1M AR 50°C f) 1M SA 50°C g)

1M HT1 50°C h) 1M HT2 50°c


134

Figure 8 SEM images of 718 crevice tips after immersion at 50°C in 1M NaCl

Solution for a) and b) HT1 and c) and d) HT2

Figure 9 shows samples from 625+ after immersion. For the as-received samples the

photos show crevice corrosion for both 0.1 M and 1 M after immersion at 60°C. The

samples for SA, HT1, and HT2, are all shown after immersion at room temperature

for both 0.1M and 1M. Immersion at room temperatures reveals a transpassive film,

but no observable crevice corrosion.

a) b)

c) d)


135

Figure 9 Photographs of 625+ specimens after electrochemical testing at 60° C for

AR, and RT for SA, HT1 and HT2 in both 0.1M and 1M NaCl solutions.

Figure 10 shows SEM images for two samples after immersion at room temperature

and 50°C in 1M NaCl for HT2. Figure 10 a/b) show the sample surface of 625+ HT2

after immersion at room temperature. No crevice corrosion was observed for this

sample, but as can be seen on Figure 9, a black film has coated the surface. Figure

10a) shows this black film, and where no film has formed under the crevice former.

As the film has a black colour to it, this would indicate it is thicker than the passive

film. The colour is thought to derive from a nickel oxide formation. Figure 10 c/d)


136

show clear crevices after sample immersion at 50 °C, with indicators of intergranular

attack as the microstructure has been releveled in both and the needle-like delta

phase has been exposed. The SEM images also reveal that the crevice attack is not

uniform beneath the crevice former with the attack being more acute at the

boundaries of the crevice former.

a) b)

c) d)

Figure 10 SEM images of 625+ HT2 after immersion in 1M NaCl a) and b) film on

sample surface at room temperature where no crevice occurred and c) and d) crevice

tip at 60°C


137

7.5 Discussion The crevice corrosion resistance of Alloys 718 and 625+ in deaerated 0.1M and 1M

NaCl solutions were investigated through electrochemical testing at temperatures

ranging between room temperature (range of 20°C to 28°C) and 60°C. Alloy 625+

showed a greater resistance to crevice corrosion, for all heat treatments, than 718 as

summarised in Figure 6. The data shows that the as-received heat treatments for

both alloys maintain to be the heat treatments with the highest corrosion resistance

for each respective alloy, and effects of microstructure and differences in the

composition of the alloys have been previously discussed in Chapters 5 and 6.

During the crevice corrosion process, there is deoxygenation of the electrolyte due

to oxygen reduction occurring faster than diffusion which is hindered by the crevice

geometry. Once the oxygen has been depleted within the crevice, oxygen reduction

only occurs on the metal surface outside of the crevice. Metal dissolution ions are

hydrolysed, causing a reduction in the pH, increasing the rate of metal dissolution.

The process is now autocatalytic in nature.14

As part of the degradation mechanism, there is also the breakdown of the passive

film. Hoar et al15 suggests that pitting; and if following the meta-stable pit-theory of

crevices formation, also crevices; occurs as a result of adsorption of aggressive anions

such as Cl on an oxide film, followed by penetration of this film.

When looking at the effects of chloride in this study there is sometimes an increase

in CCT with a reduction of chloride concentration from 3.5 wt% to 0.1M. There can,

however, also be a decrease in CCT with an increased chloride concentration to 1M.

Indicating there may be a cumulative effect of chloride and temperature on the

b)


138

crevice corrosion behaviour, and agreeing with results Pardo presented in 2000 on

two high-alloyed stainless steels. Han 201216 found the CCT of SAF 2205 DSS

decreased linearly with increasing the chloride ion concentration. In order to find out

if this was true for the investigated nickel alloys, a more sensitive temperature step

would have to be taken. Although these two studies were on high-alloyed steels the

compositions include similar components including Cr and Mo which will have similar

corrosion resistance effects.

Not only is the CCT affected, but the polarisation curves reveal a shortening of the

passive region with increasing chloride concentration. An indication that the crevice

is initiating and propagating more readily, which could be accredited to the increase

in chloride ions aiding in the breakdown of the passive film. Abdallah et al17 also

observed the shortening of the passive region during electrochemical testing of 316

SS when chloride concentrations were increased. Generally the severity of the

crevices increased with increasing chloride concertation; but due to the stochastic

nature of crevice corrosion, or a human error in the set-up, this cannot be said for all

samples. Figure 7 a) & e) are the prime example; where 7a) is an AR 718 sample which

has undergone crevice corrosion in 0.1M NaCl solution at 50°C, but the crevices are

more pronounced here than for 7e) which was immersed in the 1M NaCl, also at 50°C.

Zhou et al18 came to a similar conclusion when investigating the effects of chloride

on X80 pipeline steel. They observed a greater number of pits with an increase in

chloride concentration.

The increasing temperature may also be having an effect on stabilising the crevices

and hindering repassivation. Park et al19 identified at high temperatures and high


139

chloride concentrations, numerous additional initiation processes lead to an etching-

type corrosion of their samples. A similar etching-type corrosion where the

microstructure is revealed for all samples that underwent crevice corrosion, and for

those where an obvious film was found on the surface. The additional initiation

processes that may be possible in the presence of chloride could be the cause of the

intergranular attack observed.

7.6 Conclusions From the results obtained regarding the effect of chloride concentration on the

localised crevice corrosion, and the effects of changing the start point of the

potentiodynamic scans of the Ni-base superalloys 718 and 625+, the following

conclusions can be drawn:

Alloy 718 is more susceptible to crevice corrosion than 625+ for all heat

treatments and chloride concentrations as it experienced crevice corrosion

more readily at lower temperatures.

No discernible pattern for the effects of [Cl-] could be formed. A reduction in

chloride concentration can result in an increase in the CCT when going from

0.6 M to 0.1 M, but the CCT will then remain constant to 1M. An increase in

the chloride concentration did result in a decrease in CCT, if the CCT had not

changed with previous chloride concentration changes.

The as-received heat treatments for both alloys have been shown to be the

most corrosion resistant when compared to the other heat treatments.

The corrosion resistance follows the pattern AR > HT2 > HT1 > SA for both

alloys in both chloride conditions.


140

7.7 References 1. Pinder, D. Offshore oil and gas: global resource knowledge and technological

change. Ocean Coast. Manag. 44, 579–600 (2001).

2. Ocean Salinity. Environmental Science: In Context 2, (Gale, a Cengage

Company, 2008).

3. Talley, L. D. et al. Typical Distributions of Water Characteristics. Descr. Phys.

Oceanogr. 67–110 (2011).

4. Ozawa, M., Yamaguchi, A., Ikeda, T., Watanabe, Y. & Ishizaka, J. Abundance

and community structure of chaetognaths from the epipelagic through

abyssopelagic zones in the western North Pacific and its adjacent seas. Plankt.

Benthos Res. 2, 184–197 (2007).

5. Speight, J. G. & Speight, J. G. Corrosion. Subsea Deep. Oil Gas Sci. Technol. 213–

256 (2015).

6. Burstein, G. T., Liu, C., Souto, R. M. & Ines, S. P. V. Origins of pitting corrosion.

(2004).

7. Hartt, W. H., Kumria, C. C. & Kessler, R. J. Influence of Potential, Chlorides, pH,

and Precharging Time on Embrittlement of Cathodically Polarized Prestressing

Steel. CORROSION 49, 377–385 (1993).

8. Truman, J. E. The influence of chloride content, pH and temperature of test

solution on the occurrence of stress corrosion cracking with austenitic stainless

steel. Corros. Sci. 17, 737–746 (1977).

9. Yang, Y., Khan, F., Thodi, P. & Abbassi, R. Corrosion induced failure analysis of


141

subsea pipelines. Reliab. Eng. Syst. Saf. 159, 214–222 (2017).

10. Mccoy, S. A., Puckett, B. C. & Hibner, E. L. High Performance Age-Hardenable

Nickel Alloys Solve Problems in Sour Oil and Gas Service. J E Ward Corrotherm

Int. Ltd (2011).

11. Bai, Y. & Bai, Q. in Subsea Engineering Handbook 703–761 (Elsevier, 2012).



296–305 (2017).



Solution. 1–13 (2011).

14. Heppner, K. L., Evitts, R. W. & Postlethwaite, J. Effect of Ionic Interactions on

the Initiation of Crevice Corrosion in Passive Metals. J. Electrochem. Soc. 152,

B89 (2005).

15. T. P. Hoar. The production and breakdown of the passivity of metals. Corros.

Sci. 7, 341–355 (1967).

16. Ebrahimi, N., Jakupi, P., Noël, J. J. & Shoesmith, D. W. The role of alloying

elements on the crevice corrosion behavior of Ni-Cr-Mo alloys. Corrosion 71,

1441–1451 (2015).

17. M. Abdallah, B. A. AL Jahdaly, M. M. Salem, A. Fawzy, A. A. A. F. Pitting

Corrosion of Nickel Alloys and Stainless Steel in Chloride Solutions... J. Mater.

Environ. Sci. ISSN 2028-2508 8, Page 2599-2607 (2017).


142

18. Zhou, W. et al. Effect of a high concentration of chloride ions on the corrosion

behaviour of X80 pipeline steel in 0.5 mol L-1NaHCO3solutions. Int. J.

Electrochem. Sci. 13, 1283–1292 (2018).

19. Park, J. O., Matsch, S. & Böhni, H. Effects of Temperature and Chloride

Concentration on Pit Initiation and Early Pit Growth of Stainless Steel. J.

Electrochem. Soc. 149, B34 (2002).

Chapter 8; Investigation of the galvanic crevice corrosion behaviour of Inconel 718 and Custom Age 625+

143

CHAPTER 8; INVESTIGATION OF THE

GALVANIC CREVICE CORROSION BEHAVIOUR

OF INCONEL 718 AND CUSTOM AGE 625+

M. Keogh, D. Engelberg

8.1 Abstract The galvanic crevice corrosion behaviour of the alloys 718 and 625+ was investigated

via potentiostatic polarisation. Both the alloys were coupled together, fitted with

ceramic crevice formers, and exposed to deaerated 3.5 wt% NaCl at 45°C. Alloy 718

was found to be more susceptible to crevice corrosion, as the number of crevices and

total volume within the crevice was higher for 718 than 625+. Alloy 625+, however,

also undergoes crevice corrosion below far its previously reported Critical Crevice

Temperature (CCT) of 60°C. It is thought that galvanic effects influence the crevice

corrosion behaviour; making alloys more susceptible to crevice corrosion when a

galvanic effect is at play.


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8.2 Introduction With the development of advanced marine engineering to combat pressures of oil

and gas demand, more materials; in particular new alloys, are being introduced into

the ocean engineering systems.1 It is therefore possible that two differing alloys may

come into electrical contact with one another; with the different electrical potentials

between the alloys allowing for the possibility of galvanic polarisation and corrosion

to occur2. Often, steps are taken to minimise galvanic corrosion such as using metals

with similar electrical potentials, coatings3, or cathodic protection.4 These protection

systems, however, are not always guaranteed and there may be lapses which leave

the alloy vulnerable for corrosion to occur.5 Understanding the possible corrosion

outcomes is thus vital until fail-safe protection mechanisms are in place.

Literature is available to understand galvanic corrosion mechanisms in commonly

used steels6 and copper-alloys.7 The summary in this paper will focus mainly on alloys

with similar properties to those of the nickel alloys employed in the energy industry.

A study by Barik et al8 on y’-precipitation-hardened Cu-Ni alloys with similar

mechanical and corrosion properties to the alloys utilised in this study, indicated that

alloys with a lower Ni content act as the anode, whilst the higher Ni-content alloy,

Cu-19Ni, acted as the cathode. The galvanic corrosion susceptibility lessened as the

protective films on the alloys become more similar in composition, with the

overarching factor affecting extent of galvanic corrosion being time of immersion.

Mansfield et al9, also found that Cu-Ni alloys were susceptible to galvanic corrosion,

with the main contributing factor to severity, again, the immersion time. These alloys

showed an intergranular corrosion, thought to be due to nickel at the grain


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boundaries being the preferred inlet for attack. It was proposed that the corrosion of

the copper alloys was under both mass and charge transfer control. All other alloys

in this study, stainless steels and titanium, however, were found to only be under

charge transfer control due to the passive layer.

A review by Francis10 highlighted some of the experimental findings of galvanic

corrosion of stainless steels in seawater environments. Findings from this review

include type 300 stainless steels becoming more susceptible to pitting and crevice

corrosion when they are coupled to more corrosion resistant alloys. The ease of

initiation is unaffected, but the rate of propagation is increased when an increase in

the cathodic area is available. The susceptibility to localised corrosion is pertinent in

the results presented in this study. Shone et al11 conducted a study in which high

alloyed stainless steels were coupled to corrosion-resistant nickel alloys and titanium.

The results found that no galvanic corrosion took place between the steels and the

nickel alloys or titanium.

Ni-base superalloys are used extensively within the oil & gas industry due to their

high strength and superior corrosion.12 The main function of these alloys is to act as

connectors in the subsea well heads. Previous work (reference other chapters) has

identified that these alloys may be susceptible to localised corrosion. The purpose of

this work was to investigate whether the alloys would undergo galvanic crevice

corrosion in seawater environment.


Discharge Machining) from annealed ingots into 20 mm x 20mm x 5mm cuboids, with


146

an 8mm hole in the centre. The chemical composition is provided in Tables 1. Both

alloys were supplied in a solution annealed at 1020 °C for 1 hour, with alloy 718

having then undergone an age-hardening at 760 °C for 8 hours, followed by an air

cool.



718 54.70 18.60 17.00 5.07 3.04 0.45 0.88

625+ 61.10 21.00 4.24 3.65 8.10 0.22 1.49

Figure 1 shows the crevice former set-up used, based on the ASTM G4813 which

utilises ceramic crevice formers made from Technox® 2000, a zirconia ceramic, which

are wrapped with PTFE tape. Two sample specimens, one of 718 and one of 625+ are

coupled, and then crevice formers are attached to either side of the exposed

specimen side with steel washers, bolt and nut which secure the set-up. The crevice

former set-up is electrically isolated from the samples. The alloy 718/625+ set-up is

compared to individual metal samples under the same exposure conditions. The

torque applied to the crevice former was 1.5 N m. Specimens were ground to a SiC

paper 600 finish, and then washed with acetone and distilled water. Samples were

allowed to air-dry for a minimum of 24 hours before testing.

All the tests were conducted in a one-litre three-electrode cell. The reference

electrode used was a 3M Ag/AgCl, and a platinum counter electrode. Nitrogen was

purged for a minimum of 30 minutes to ensure deaeration of the solution before




147

which was kept at constant temperature, and monitored throughout the test with a

thermometer placed inside the cell.

Crevice Corrosion tests were performed in a 3.5 wt% NaCl solution at 45 °C. The open

circuit potential (ECORR) was monitored for 5 minutes, potentiodynamic polarisation

scans were carried out from OCP in the anodic direction at a rate of 1 mV/s until end

points of; +150 mV, +200 mV and +250 mV vs Ag/AgCl where the potential was kept

constant for 24 hours. Specimens were characterised after testing with Quanta 200

Scanning Electron Microscopy (SEM) and a Keyence Laser Scanning Confocal

Microscope (LSCM).

Figure 1 Coupled crevice former experimtenal set up

8.4 Results Figure 2 shows the potentiostatic polarisation curves for 718 coupled with 625+, and

compared to both individual metal samples at 45°C during a hold at +150 mV vs

Ag/AgCl for 24 hours. There is an increase in current density observed shortly after

experiment initiation for all samples, which indicates that a localised corrosion event

has taken place. Localised corrosion was confirmed by visual inspection (Figure 3),


148

showing that both 718 and 625+, and the coupled samples had undergone crevice

corrosion underneath the crevice formers. No corrosion was observed at the metal

to metal interface of the coupled sample.

Figure 2 Electrochemical data for 718, 625+ and coupled set-ups showing current

response from +150 mV hold for 24hours

Table 2 OCP values for 718, 625+ and coupled samples prior to current hold at 150

mV for 24 hours


149

Figure 3 Photographs of samples post-testing; 625+ only, 718 only, & 718/625+

coupled samples after electrochemcial testing at 150 mV for 24 hours at 45°C

Crevices formed under the crevice former in the coupled samples were observed

under the SEM as shown in Figure 4. Both alloys show indication of intergranular

corrosion which is observed underneath the corrosion product.

Figure 4 SEM images of coupled 718 (a, b, c) and 625+ (c, d, e) crevice corrosion


150

LSCM was used to quantify the volume of the crevices on the sample surfaces

through the creation of a contour map of the crevice and the surorunding unaffected

sample surface (Figure 5).

Figure 5 LSCM images of individual samples showing signs of crevice corrosion

after 150 mV potentiostatic hold for 24 hours at 45°C in deaerated 3.5 wt% NaCl

solution a) 625+ b)718

Testing at 200 mV showed similar results (Figure 6) to those obtained for +150 mV

with the 625+ having the lowest current response, followed by 718, followed by three

coupled experiments. One experiment coupled 718 with 625+, the second with 718

coupled to 718, and the third 625+ coupled to 625+.Increases in current are seen for

all experiments, with again, the current increase being higher for 718 than 625+. The

highest current observed was that for the individual 718 sample. The coupled

experiments which contained a 718 sample showed similar electrochemical current

responses. The experiments which contained only 625+; both individual and coupled;


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also showed similar current response to one another. Visual inspection (Figure 7)

showed that all samples had experienced crevice corrosion underneath the crevice

formers, but the metal-metal interface of the coupled samples showed no sign of

attack.

Figure 6 Electrochemical data for 718, 625+ and coupled set-ups showng resultant

current from +200 mV hold for 24 hours

Table 3 OCP values for 718, 625+ and coupled samples prior to current hold at 200

mV for 24 hours

2D Graph 2

Time / s

0 20000 40000 60000 80000 100000

i /

mA

cm

-2

0.0

0.2

0.4

0.6

0.8

1.0

1.2

625+

718

718 & 625+

718 & 718

625+ & 625+


152

Figure 7 Photographs of sampes post-tesing; 625+ only, 718 only, & 718/625+

coupled samples after electrochemcial testing at +200 mV for 24 hours at 45°C

.

LSCM confirmed the presence of crevice corrosion as shown in Figure 8. Analysis from

the volume within side the crevice showed that there was a greater volume inside

the crevice of 718 compared to that of 625+.


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Figure 8 LSCM images of individual samples showing signs of crevice corrosion

after 200 mV potentiostatic hold for 24 hours at 45°C in deaerated 3.5 wt% NaCl

solution a) 625+ b)718

Figure 9 shows the potentiostatic polarisation curves for 718 coupled with 625+

during a hold of +200 mV vs Ag/AgCl for 24 hours, immediately followed by a hold of

+250 mV vs ag/AgCl for 24 hours. The sharp increase in current density observed

shortly after experiment initiation indicates that a localised corrosion event has taken

place was confirmed by visual inspection that both 718 and 625+ had undergone

crevice corrosion. The crevice corrosion had taken place both underneath the crevice

former, and at the metal-metal interface between 718 and 625+.


154

Figure 9 Potentiostatic Polarisation of 718 coupled with 625+ in 3.5 wt% NaCl

solution at 45°C; 24 hr hold at +200 mV –followed by 24hr hold +250 mV

Figure 10 a) and b) shows the crevice damage of 718 that occurred underneath the

crevice formers and that occurred at the metal-metal crevice interface with 625+

respectively. Figure 10 c) and d) show the same damage pattern for 625+.

2D Graph 3

Time / s

0 25000 50000 75000 100000 125000 150000 175000 200000

i /

mA

cm

-2

0.0

0.2

0.4

0.6

0.8

1.0


155

Figure 10 a) 718 crevice damage underneath the crevice formers; b) 718 crevice

damage at the metal-metal crevice interface with 625+; c) 625+ crevice damage

underneath the crevice formers and b) 625+ crevice damage at the metal-metal

crevice interface with 718.

Figure 11 shows SEM images the crevices formed underneath the ceramic crevice

formers of post-immersion 718 and 625+ specimens at 45°C. Figure 11 a/b) show an

apparent intergranular attack pathway for 718. Figure 11 c/d) shows a similar attack

pathway for 625+. Figure 12 shows the crevice damage on 718 from the metal

interface with 625+ which too follows an intergranular attack pathway.


156

Figure 11 SEM images of crevices formed underneath the ceramic crevice former

a/b) 718; c/d) 625+


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Figure 12 Grain boundary exposure of 718 after crevice attack at a metal interface

with 625+

8.5 Discussion The galvanic and crevice corrosion behaviour of alloys 718 and 625+ in deaerated 3.5

wt% NaCl solution was investigated through potentiostatic electrochemical testing at

and 45°C. Alloy 625+ showed a greater resistance to crevice corrosion than 718; both

in the post-immersion sample inspections and through the electrochemical

behaviour. It is suspected that 625+ has greater corrosion resistance due to its

composition. The main difference in composition between the two alloys is an

increased Mo content for alloy 625+ when compared to 718. Molybdenum additions

have been known to restrict the initiation of crevices and reduce propagation rates14

. Although one mechanism has not been agreed for how the molybdenum can affect


158

the crevice initiation and propagation, it is thought to be due to the production of a

MoO42- ion which can provide a protective oxide layer14.

The SEM images for both alloys show that the crevice corrosion show the attack

pathway may be intergranular, due to the exposed grain boundaries, beneath the

accumulated corrosion product, which indicates the precipitates, or matric elements,

found at the grain boundaries may have a role in facilitating the corrosion pathway.

In conditions where the y’ and y” are the most abundant an intergranular corrosion

pathway is observed. Hwang et al found that the gamma-phases can be more

electrochemically active than the matrix thus acting as a localised anode15. As these

precipitates can be found at the grain boundaries, it may rationalise the intergranular

pathway seen on SEM images during post-immersion analysis. The intergranular

corrosion (ICG) pathway was also seen by Mansfield9 in galvanic corrosion tests for

Cu-Ni alloys.

The corrosion product build-up was more visible on the 718 samples than 625+ which

may be due to the 718 being more susceptible to corrosion. The confocal data

showed that all 718 crevices were of a greater volume than 625+. The greater volume

indicates a greater amount of matrix loss; hence the increased presence of corrosion

product. There could also be a combined effect of the corrosion product acting as a

stabiliser for the crevice, with the corrosion continuing to propagate underneath.

Confocal images and analysis show that crevice corrosion may propagate quicker

through 718 than 625+ as there is a greater volume within the crevice, supporting

findings by Francis10 where localised corrosion propagates quicker through 316 when


159

coupled to higher-alloyed stainless steels, or the corrosion product is stabilising the

crevice.

These tests have shown that 625+ can undergo crevice corrosion lower than the

previously reported CCT. This may be due to the severity of the test, as this is also

seen in the modified THE (chapter reference). Hornus et al16 also reported nickel

alloys undergoing crevice corrosion below previously believed CCTs after undergoing

the modified THE method. To investigate whether the severe testing is the cause of

the premature crevice corrosion a potentiostatic temperature ramping test could be

employed.

A crevice was only observed at the metal-metal interface of the test specimen which

was subjected to the double polarisation, indicating that galvanic corrosion is unlikely

to occur between these two alloys under normal service conditions. It is worth noting,

however, that the corrosion suffered during the doubled testing was extremely

severe both at the metal-metal interface and underneath the crevice formers. There

may be an element of galvanic coupling, as both samples show a greater number of

crevices in the coupled tests than the individual tests.

8.6 Conclusions From the results obtained regarding the effect of coupling 718 and 625+ on the

galvanic and crevice corrosion, and the effects of changing the potential during

potentiostatic scans of the Ni-base, the following conclusions can be drawn:

Alloy 718 is more susceptible to crevice corrosion than 625+ as it experienced

more severe crevice corrosion at all potentials.


160

• 625+ underwent crevice corrosion below its previously recorded CCT (50°C-

60°C). This is thought to be due to the severity of the potentiostatic testing.

• Both alloys suffered an intergranular corrosion pathway, within both the

crevice former artificial crevice and the metal-to-metal crevice.

• Only during the double potentiostatic experiment where the test specimens

were subjected to 200 mV, immediately followed by 250 mV did a crevice

occur at the metal-metal interface.

8.7 References 1. Wolf, S. D., Brillson, L. J. & Dimos, D. B. Advanced Materials for Our Energy

Future. Mater. Challenges Altern. Renew. Energy Conf. (2010).

2. Wilhelm, S. M. Galvanic corrosion in oil and gas production: Part 1 - laboratory

studies. Corrosion 48, 691–703 (1992).

3. Olajire, A. A. Recent advances on organic coating system technologies for

corrosion protection of offshore metallic structures. J. Mol. Liq. 269, 572–606

(2018).

4. Bahadori, A. & Bahadori, A. Principle of Electrochemical Corrosion and

Cathodic Protection. Cathodic Corros. Prot. Syst. 1–34 (2014).

5. Wang, W., Shen, K., Yi, J. & Wang, Q. A mathematical model of crevice

corrosion for buried pipeline with disbonded coatings under cathodic

protection. J. Loss Prev. Process Ind. 41, 270–281 (2016).

6. Tsujino, B. & Miyase, S. Galvanic Corrosion of Steel in Sodium Chloride


161

Solution. Corrosion 38, 226–230 (1982).

7. Du, X. Q., Yang, Q. S., Chen, Y., Yang, Y. & Zhang, Z. Galvanic corrosion behavior

of copper/titanium galvanic couple in artificial seawater. Trans. Nonferrous

Met. Soc. China (English Ed. 24, 570–581 (2014).

8. Barik, R. C., Wharton, J. a., Wood, R. J. K. & Stokes, K. R. Galvanic corrosion

performance of high strength copper-nickel alloys in seawater. 65, 359–367

(2009).

9. Mansfeld, F., Liu, G., Xiao, H., Tsai, C. H. & Littler, B. J. of Copper Alloys , in

Seawater. Science (80-. ). 36, 2063–2095 (1994).

10. Francis, R. Galvanic Corrosion of High Alloy Stainless Steels in Sea Water. Br.

Corros. J. 29, 53–57 (1994).

11. Gallagher, P., Malpas, R. E. & Shone, E. B. Corrosion of stainless steels in

natural, transported, and artificial seawaters. Br. Corros. J. 23, 229–233 (1988).

12. Rhodes, P. R. Environment-assisted cracking of corrosion-resistant alloys in oil

and gas production environments: A review. Corrosion 57, 923–966 (2001).



Solution. 1–13 (2011).

14. Lillard, R. S., Jurinski, M. P. & Scully, J. R. Crevice Corrosion of Alloy-625 in

Chlorinated Astm Artificial Ocean Water. Corrosion 50, 251–265 (1994).



162


16. Hornus, E. C., Rodríguez, M. A., Carranza, R. M., Giordano, C. M. & Rebak, R. B.

Effect of Environmental Variables on Crevice Corrosion Susceptibility of Ni–Cr–

Mo Alloys for Nuclear Repositories. Procedia Mater. Sci. 8, 11–20 (2015).

9; Discussion and Conclusions

163

9; DISCUSSION AND CONCLUSIONS

This chapter sets out to build upon the discussions in the previous chapters, by

encompassing decisions about experimental preferences, and bringing the chapters

together into a coherent body of work; culminating in presenting the conclusions of

the thesis.

9.1 Corrosion Behaviour

This project set out to investigate the localised corrosion behaviour of two commonly

used nicker-superalloys; 718 and 625+ through electrochemical testing. A variety of

electrochemical methods have been utilised, each with their own advantages; which

are discussed in more details in the following sections and in Chapter 10; Future

Work.

Throughout this project, it has been consistently shown that alloy 718 has a greater

susceptibility to crevice corrosion when compared to alloy 625+. Not only in the lower

recorded Critical Crevice Temperatures, but also in the corresponding crevice

potentials (ECREV). Differences in the corrosion susceptibility have been attributed to

the differences in elemental compositions. The main difference between the alloy

compositions is Alloy 718 having a lower Mo content when compared to Alloy 625+.

Molybdenum is understood to generate protective molybdenum oxides during the

passive film breakdown which can slow the crevice propagation rate.1,2


164

Despite the difference in susceptibility, the attack pathways have been similar for

both alloys, and the possible causes are discussed in the proceeding section of this

chapter. Both alloys are also unlikely to undergo crevice corrosion in service during

standard operating conditions, as neither alloy undergoes crevice corrosion at room

temperature; although temperatures below this have not been investigated.

9.2 Microstructure

Through changing the microstructure of the alloys through heat treatments, the

effects of the precipitates have been able to be investigated. An unexpected result

came from the electrochemical tests in that for both alloys the solution anneal

condition had the least corrosion resistance when compared to the other heat

treatments. Initial assumptions assumed the solution anneal would have increased

corrosion resistance as the precipitates would not be present to act as initiation

points for the corrosion to occur.

It is however, the combination of precipitates which provide the greatest corrosion

resistance, with δ phase seemingly the most contributing factor, despite it being

incoherent with the matrix. Post immersion analysis of samples were delta phase was

the primary precipitate showed that the delta phase remained, and assumption was

made that the matrix was being attacked preferentially over the delta phase.

In conditions where the y’ and y” are the most abundant show an intergranular attack

pathway. Previous studies Hwang et el3 have concluded that the gamma-phases are

more electrochemically active than the matrix, thus they corrode preferentially. As

these precipitates like to congregate at the grain boundaries, this would explain the


165

intergranular pathway seen on SEM images during post-immersion analysis. The

galvanic coupling of the gamma phases to the matrix could have further contributions

to crack pathways in stress corrosion cracking and hydrogen embrittlement3.

9.2 Environment

When the change in environment was taken into account, the same pattern of

corrosion resistance was seen throughout the alloys and heat treatments, indicating

at the amount of chloride present does not change the preferred attack pathway.

Intergranular attack was seen in the solution anneal, and for the alloys containing y’

and y”, and matrix attack occurred for the alloys with delta phase as the primary

precipitate.

The chloride concentration could, however, change the minimum temperature at

which crevice corrosion could occur. So although the chloride does not change how

the corrosion occurs, it can change when. There was no pattern to the effect of

chloride on the CCT, with some alloys and heat treatments maintaining a steady CCT

for 0.1 M and 0.6 M; but for others it remained constant between 0.6 M and 1 M.

In the literature it has been common to see the CCT reducing linearly with log [Cl-]4,

but this has not been observed during this work. It could be possible to obtain this

result if experiments were to be completed at smaller temperature ranges than 10°C.

Literature investigations of localised corrosion often utilise solutions containing FeCl3

and HCl5, which is a severe environment due to the high chloride concentration, and

the Fe3+ ions promote electron production which increases the corrosion rates.

Through operating in a 3.5 wt% NaCl envrionment it is more representative of the


166

conditions that these alloys will experience during service, and thus a more accurate

depiction of the localised corrosion behaviour of these alloys can be obtained. The

crevice environments during this work have, however, been artificially created. It is

discussed in Chapter 10; Future Work that using samples with a threaded geometry

may be more representative of in-service conditions. Despite this, the results

presented here provide a good baseline understanding of how these alloys are

susceptible to crevice corrosion. The information provided during this thesis could be

of more use for alternative marine applications of the alloys or for the nuclear

industry where operating temperatures can be higher.

9.4 Galvanic Crevice Corrosion

The galvanic crevice corrosion tests during this project have shown that there can be

an additional galvanic effect when similar metals are coupled together in a geometry

which can cause crevice corrosion. As with the potentiodynamic tests that took place

throughout the work, Alloy 718 was more susceptible to crevice corrosion during

potentiostatic testing. The difference in corrosion behaviour between the two alloys

has been attributed to differences in composition; with alloy 625+ having less iron

and more molybdenum. Molybdenum has been knows to increase the corrosion

resistance when added as an alloying element.6

The coupled testing method, however, is something that has not been widely seen in

the literature; and so there is little to compare the results with; ensuring what has

been observed here is a true representation of the corrosion behaviour of the alloys

under the coupled conditions. The results obtained, in particular the CCT, contradict


167

previous results. Despite the CCT for 625+ being recorded as being in the range of 50-

60°C during the potentiodynamic testing in 3.5 wt% NaCl, during the galvanic

coupling tests, it underwent crevice corrosion at 45°C. This is thought to the due to

the severity of the test methods. A similar concept was observed during this work

when the THE method was used, and it has been seen in the literature by7 when also

using the THE method. The two electrochemical methods are similar in that they

utilise a potentiostatic set up, rather than a potentiodynamic one.

The environment between the two samples is good enough to propagate crevice

growth, as seen in the experiment where there were two potentiostatic holds (200

mV followed by 250 mV), but all other potentiostatic holds only yielded crevices

underneath the ceramic crevice formers, indicting this is the preferential crevice

initiation site. The preference could be due to the distance between the coupled

samples being too great to stabilise a crevice, except when aided by a high potential.

A future investigation would be to possibly utilise threaded samples with the bolt

acting as the coupling alloy. It is known that threaded components are susceptible to

crevice corrosion due to their geometry.8

The galvanic couple, however, is a more realistic representation of the conditions in

service, when compared to single samples with crevice formers. These alloys are

unlikely to be used in marine situations where they isolated from other alloys,

therefore, investing the potential galvanic effect on the crevice corrosion behaviour

is important to understand when wanting to understand how these alloys may

behave in service; especially if there is a breakdown in the cathodic protection.


168

9.4 Conclusions

Alloy 625+ has a greater resistance to crevice corrosion than alloy 718. Alloy 625+ has

a higher CCT (50-60) compared to 718 (30), a higher ECREV at CCT, and when it does

suffer crevice corrosion it is to a lesser extent with fewer shallower crevices. The

resistance is thought to be due to the compositional differences between the alloys.

As the average in-service temperature of the areas surrounding the alloys is

approximately 4°C, it is unlikely either of these alloys would undergo crevice

corrosion in service, unless other external factors can into consideration.

Microstructure does have an effect on the corrosion behaviour of alloys 718 and

6 5 , with the combined effects of the γ’, γ”, and δ-phase precipitates providing the

most localised corrosion resistance. The alloys are currently utilised in their as-

received states which do contain all precipitates and so no change to the heat

treatment procedures currently in place for these alloys should be changed; as doing

so would not increase the localised corrosion resistance.

Changes in the chloride concentration also have an effect on the localised corrosion

behaviour of the two alloys. Although no discernible pattern was distinguished,

reducing the chloride concentration could raise the reported CCT; and similarly

increasing the chloride concentration could give way to a decreased CCT. Changes in

chloride concentration are unlikely to be so varied throughout the oceans where

these alloys are in service, so the CCT for in service alloys can be taken at as that

recorded for 3.5 wt% NaCl.


169

9.5 References

1. Henderson, J. D., Li, X., Shoesmith, D. W., Noël, J. J. & Ogle, K. Molybdenum

surface enrichment and release during transpassive dissolution of Ni-based

alloys. Corros. Sci. 147, 32–40 (2019).

2. Lloyd, A. C., Noël, J. J., McIntyre, S. & Shoesmith, D. W. Cr, Mo and W alloying

additions in Ni and their effect on passivity. Electrochim. Acta 49, 3015–3027

(2004).



4. Han, D., Jiang, Y. M., Shi, C., Deng, B. & Li, J. Effect of temperature, chloride ion





296–305 (2017).

6. Hayes, J. R., Gray, J. J., Szmodis, A. W. & Orme, C. A. Influence of Chromium

and Molybdenum on the Corrosion of Nickel-Based Alloys. CORROSION 62,

491–500 (2006).



8. Larché, N. & Thierry, D. Crevice Corrosion performance of High Alloy Stainless


170

Steels and Ni-based Alloys in Seawater Applications, Join Industry Program

(confidential). (2012).

Chapter 10; Future Work:

171

CHAPTER 10; FUTURE WORK:

10.1 Microstructure The microstructure has been resolved for the alloys 718 and 625+ in this thesis in so

far as micrographs were established for each heat treatment, and general trends of

corrosion behaviour have been established. Further analysis of the effects of

particular precipitates and carbides have on the corrosion resistance could be

investigated with other alloys. As previously mentioned, 625+ is a descendant of the

alloy 625 which is solution hardened alloy. The solution annealed condition for 625+

would theoretically be similar to that of 625. As this work has compared the corrosion

resistance of 625+ to 718; it could also compare the corrosion behaviour of 625 and

625+ to determine if the precipitation hardening bares an effect on the corrosion

behaviour.

10.1.1 THE for other microstructures

Through using the THE technique for the as-received microstructures of the alloys

718 and 625+ the protection potential was established. Although the other

microstructures have been deemed less corrosion resistant, it would be pertinent to

understand the role of the microstructure of the passivation of the alloys. This work

would have the potential to be extended to cover repassivation of an established

crevice which could simulate the breakdown, and re-employment of cathodic


172

protection which is known phenomenon within the oil industry. It could contain the

above mentioned comparisons with 625, other alloy alloys similar to 718.

10.1.2 Confirm CCT through potentiostatic temperature ramping

To further define the CCT, a potentiostatic approach to electrochemical crevice

corrosion could be taken. A similar method to that employed by Han et al1 when

studying the effect so temperature on the crevice corrosion behaviour of super

duplex stainless steel could be utilised in the study for alloys 718 and 625+. A holding

potential of +750 mV v SCE was set, and then the temperature increased by 2 °C over

a period of 2 minutes and held for 30 minutes. When there was a significant current

response registered this was taken as the crevice initiation point, and thus the CCT

could be recorded.

Figure 4 Potentiostatic Crevice Corrosion test with temperature ramping to

determine CCT1


173

Through having two independents routes to determine the CCT, a more accurate

temperature range of where the alloys could undergo damage via crevice corrosion

could be established.

10.2 Environment

10.2.1 Other Chloride concentrations

Only three chloride concentrations have been investigated in this thesis. If these

alloys are to continue to have applications in other environments, other than marine,

move severe chloride concentrations could be investigated. New nuclear reactors are

considering the use of molten chloride salts heat transport fluids to transfer high-

temperature process heat from nuclear reactors to power chemical plants.2 The

current findings of this work indicate that increasing the chloride concentration can

decrease the CCT.

10.2.2 Galvanic Corrosion of other heat treatments

The galvanic crevice corrosion behaviour commenced to be investigated in the work

presented in this thesis. As previously mentioned these alloys, or alloys with

dissimilar galvanic properties, may be electrical contact with one another during

service. Further work into the galvanic crevice corrosion behaviour of these alloys

need to take place for it to be fully understood. Changes in the temperature and the

potential range would be the first parameters to be changed and investigated.


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10.2.3 Combined effects of crevice and stress

When these alloys are in service, they are not only susceptible to corrosion

environments, but also stress. Work by Rezig et al has looked at the combined effects

of crevice corrosion and stress, finding that a crevice may act as an initiation point for

a fatigue crack.4 Other work has similarly been conducted to study the pit to crack

transition in X65 steel.5 In order for this work to be completed, first a baseline tensile

and fatigue strength would have to be established. Once established, crevices can be

introduced and the effects on the strengths monitored. Post-experiment analysis

could then reveal if the crevice acts as an initiation point for cracking. This work may

also tie in to the microstructure studies and evaluating the effects of microstructure

as well.

10.3 References 1. Han, D., Jiang, Y. M., Shi, C., Deng, B. & Li, J. Effect of temperature, chloride ion



2. Sridharan, K. & Allen, T. R. Corrosion in Molten Salts. Molten Salts Chemistry

(Elsevier Inc., 2013).

3. Rebak, R. B. Mechanisms of Inhibition of Crevice Corrosion in Alloy 22. MRS

Proc. 985, 0985-NN08-04 (2006).

4. Rezig, E., Irving, P. E. & Robinson, M. J. Development and early growth of

fatigue cracks from corrosion damage in high strength stainless steel. Procedia

Eng. 2, 387–396 (2010).


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5. Evans, C. Understanding the Pit-Crack Transition. (The University of

Manchester, 2016).

Localised Corrosion of Ni-base Superalloys in Seawater

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