Pitting corrosion of additive manufactured AISI 316L and Alloy 718

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Camilla SundsbakkPitting corrosion of additive manufactured AISI 316L and Alloy 718 NTNU Norwegian University of Science and Technology

Master ’s thesis

Pitting corrosion of additive

manufactured AISI 316L and Alloy 718

Master’s thesis in Chemical Engineering and Biotechnology Supervisor: Andreas Erbe

Co-supervisor: Roy Johnsen June 2020

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Pitting corrosion of additive

manufactured AISI 316L and Alloy 718

Master’s thesis in Chemical Engineering and Biotechnology Supervisor: Andreas Erbe

Co-supervisor: Roy Johnsen June 2020

Norwegian University of Science and Technology

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This Master’s Thesis was carried out at the Department of Material Science and Engi- neering at Norwegian University of Science and Technology (NTNU). The work has been in collaboration with DNV GL, and Equinor. The project was supervised by Professor Roy Johnsen and Andreas Erbe from NTNU.

The work has been performed by the author, with the exception of the CT scanning that was carried out by Ole Tore Buset at the Department of Physics. The experimental work was performed at the Department of Mechanical and Industrial Engineering at NTNU, with samples provided by DNV GL and Equinor.

Camilla Sundsbakk NTNU, Trondheim 28.06.2020

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I would like to show my greatest gratitude to my supervisors, Andreas Erbe and Roy Johnsen. I am particular gratefully to Roy Johnsen for his for all the guidance and support he has provided. His knowledge, enthusiasm, and availability have been highly appreciated.

I would also like to thank Cristian Torres Rodiguez and Iman Taji, who provided training and help throughout the work, and were always available and if needed. In addition, I would like to thank Di Wan for the training he provided and the effort he made to ensure enough work time in the labs after the time-limitations due to by Covid-19.

I am also thankful for all the friends I have made during my studies, the support they have given and the good times we have shared. Lastly, I would I like to express my gratitude towards my family for the encouragement they have showed me throughout my studies.

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Additive manufactured (AM) steel has favourable production aspects in both an efficient and economical view in comparison with traditional manufacturing (TM) methods. Be- fore replacing TM corrosion resistant alloys (CRA) with components made by AM, it is essential to know and understand the difference in properties between the different production methods.

This study examines the susceptibility of localized corrosion of selected CRAs produced by AM and TM. Cyclic potentiodynamic polarization according to ASTM G61 was used to measure initiation and growth of pitting corrosion on AISI 316L and Alloy 718 in water solutions containing 3.5wt% NaCl at room temperature (RT), 35^◦C, 60^◦C and 90^◦C. All tests showed more positive pitting potential for the AM samples compared to the respective TM samples. The repassivation potentials showed similar behaviour for both AM and TM samples at the different temperatures, but the TM metal showed higher difference between the open circuit potential (OCP) and repassivation potentials.

An as-built ”skin” sample from the AM 316L material was also tested in 3.5 wt% NaCl at 35^◦C, showing higher susceptibility for pitting corrosion than any other 316L sample, due to the high surface roughness.

Critical initiation temperature was measured by immersing a TM and AM sample in ferric chloride solution according to ASTM G48 and raising the temperature from room RT.

Both AM and TM 316L samples gave critical pitting temperature at RT. The TM 718 material showed a critical pitting temperature of 53^◦C, therefore being more susceptible to pitting corrosion as a effect of temperature, as the AM 718 showed critical pitting temperature at 60^◦C.

Surface analyses before the corrosion test were performed to determine the surface porosity%. The average porosity% were 0.124 for the AM 316L sample and 0.065 for the AM 718 sample. Both porosities were below 1%, which stayed in-line with the theory as the porosity did not affect the corrosion properties of the material in a negative matter. The chemical composition of the AISI 316L material was analysed with Energy-dispersive X- ray spectroscopy (EDS). The results suggested that MnS is a preferential pit initiation site on the TM samples, while small MnO and nano-sized Mn-Si oxides cause pitting corrosion on the AM samples. The EDS results for the AM 718 material suggested preferential pit initiation site at Al inclusions and NbC segregation due to high contents inside the pits.

Other surface analyses, optical microscopy (OM) and Scanning electron microscope (SEM), revealed the different amounts of corrosion attacks from each sample and the microstructure within the pits. Experiments with the Avesta cell suggested that pitting corrosion at the AM 718 preferentially occurs on the molten pool boundaries.

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Additivprodusert st˚al har gunstige produksjonsaspekter i b˚ade et effektivt og økonomisk syn sammenlignet med tradisjonelle produksjonsmetoder. Før tradisjonelt produserte korrosjonsbestandige legeringer blir skiftet ut med additivproduserte er det derimot vik- tig ˚a ha en solid forst˚aelse i forskjellene produksjonsmetodene utgjør p˚a egenskaper til metallet.

Denne studien undersøker motstandsdyktigheten for lokal korrosjon av utvalgte korrosjonsbestandige metall, produsert ved additiv og konvensjonelle metoder. Syklisk potensiodynamisk polarisering i henhold til ASTM G61 ble brukt til ˚a m˚ale initiering og vekst av gropkorrosjon p˚a AISI 316L og legering 718 i vannløsninger inneholdende 3,5 vekt% NaCl ved romtemperatur (RT), 35^◦C, 60^◦C og 90^◦C. Alle tester viste høyere positivt groppotensial for de additive prøvene sammenlignet med de korresponderende tradisjonelle prøvene. Repassiveringspotensialene viste samsvarende oppførsel for b˚ade additiv og tradisjonelle prøver ved de forskjellige test temperaturene, men høyere forskjell fra ˚apen-krets-potensial og repassiveringspotensialene for sistnevnte. En upreparert over- flateprøve fra AM 316L-materialet ble ogs˚a testet i 3,5 vekt% NaCl ved 35^◦C, som viste høyere følsomhet for gropkorrosjon enn noen annen 316L prøve p˚a grunn av høy over- flateruhet.

Initieringstemperatur ble m˚alt ved ˚a senke en tradisjonell og additiv prøve i jernkloridløsning i henhold til ASTM G48 og heve temperaturen fra RT. B˚ade additiv og tradisjonell 316L- prøver ga kritisk groptemperatur ved RT og ble derfor ikke videre økt. Det tradisjonelle 718 materialet viste en kritisk groptemperatur p˚a 53^◦C, og var derfor mer utsatt for grop korrosjon som en effekt av temperaturen, ettersom AM 718 viste kritisk groptemperatur ved 60^◦C.

Overflateanalyser ble gjennomført før korrosjonstestene for ˚a bestemme overflateporøsitet%.

Gjennomsnittlig porøsitet% var 0,124 for de additive 316L-prøvene og 0,065 for additiv 718 prøvene. Begge porøsitetene var under 1%, som i tr˚ad med teorien og forsøkene utført tilsier at korrosjonsbestandigheten av prøvene ikke ble p˚avirket i en negativ grad. Den kjemiske sammensetningen funnet ved EDS-analysen av AISI 316L-materialet antydet at MnS er et fortrinnsvis gropinitieringssted p˚a de tradisjonelle prøvene, mens sm˚a MnO og nano Mn-Si i de additive materialene er det som for˚arsaker gropkorrosjon hos dem.

EDS-resultatene for additiv 718-materialet antydet at gropinitieringen i dette materialet skjer ved Al-oppsamlinger og diffusjon av NbC.

Andre overflateanalyser, OM og SEM, avslørte den forskjellige mengden korrosjonsangrep fra hver prøve og mikrostrukturen i gropene. Eksperimenter med Avesta-cellen antydet at gropkorrosjon p˚a additiv 718 fortrinnsvis skjer p˚a smeltelinjene fra produksjonsprosessen.

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List of Figures

2.1.1 Different shapes of pitting corrosion . . . 6

2.1.2 Schematic representation of the environment inside a pit. . . 7

2.1.3 Illustration of CPP curve . . . 8

2.1.4 Crevice corrosion . . . 9

2.1.5 Maximum pit depth vs chloride concentration for AISI 316L. . . 10

2.1.6 Pitting potential vs the content of Mo in the composition at different temperatures. . . 11

2.1.7 The effect of pitting potential by increased Mo . . . 11

2.1.8 The effect of Cr on pitting potential . . . 12

2.1.9 The effect of nitrogen on pitting potential . . . 12

2.1.10 CPP plot showing the effect of temperature. . . 13

2.1.11 Potentiodynamic polarization curve of 316L in borate buffer solution at different exposure of oxygen. . . 14

2.2.1 Overview of the additive manufacturing technologies of metal production 16 2.2.2 DMLS/SLM . . . 17

2.2.3 Laser Metal Deposition . . . 18

2.2.4 Illustration of the machine used in Binder Jetting. . . 19

3.1.1 Overview of the austenitic SS microstructure by different processing methods. . . 24

3.1.2 Representation of orientation, grain boundaries, and inclusions in L- PBF 316L SS. . . 24

3.1.3 OM side views of L-PBF 625 microstructure produced by laser power of a) 150 W, b) 200 W, c) 250 W and d) 300 W . . . 25

3.1.4 Representation of porosity in down-skin . . . 26

3.1.5 Variations of porosity with varied laser power. Constant scanning speed of 0.8 m/s for L-PBF Ti6Al4V by EOS M290 system and process parameters according to their default values for 30 µm powder layer thickness 27 3.1.6 The residual stress as a function of height in a SLM component, determined by temperature . . . 29

3.2.1 Explanation of MPBs . . . 31

4.1.1 AISI 316L additive manufactured flange used to extract samples for electrochemical testing, both with and without skin. . . 36

4.1.2 Embedded, mirror polished sample . . . 38

4.1.3 Samples with copper rod . . . 39

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4.1.4 Samples tested in the beaker setup for ASTM G61 . . . 39

4.1.5 Representative image of AM 718 from ImageJ where all the pores are marked in red. . . 41

4.1.6 Schematic representation of the full Avesta cell . . . 42

4.1.7 Schematic representation of the cross section of the Avesta cell . . . . 42

4.1.8 Schematic representation of the beaker setup for CPP measurement . 44 4.1.9 Illustration of set up during CPT measurement. . . 45

4.2.1 10,000x magnification of a single pore at the AM 316L flange material. 47 4.2.2 Representation of different grade of porosity at the surface on the AM 316L flange test samples. . . 48

4.2.3 Representation of different grade of porosity at the surface on the AM 718 test samples. . . 48

4.2.4 Top view of 3D image of AM 316L flange. . . 49

4.2.5 Vertical 3D image of AM 316L flange sample. . . 49

4.2.6 Overview 3D image of AM 718 sample. . . 50

4.2.7 Vertical 3D image of the empty space inside and outside the CT scanned AM 718 sample. . . 50

4.2.8 3D image cross-section in x-y dimension of AM 625. . . 51

4.2.9 3D image cross-section in y-z dimension of AM 625. . . 51

4.2.10 CPP and OCP curves of 316L samples, TM and AM (AIDRO), measured in Avesta cell at 60^◦C in 3.5wt% NaCl. . . 52

4.2.11 Images from OM of 316L after CPP in Avesta cell, at 60^◦C and 3.5wt% NaCl. The vertical lines are from the imaging and not seen on the actual samples. . . 53

4.2.12 CPP and OCP curves of 316L embedded samples, TM and AM, measured at 25^◦C in 3.5wt% NaCl. . . 53

4.2.15 Representative OM images of 316L samples after ASTM G61 . . . 56

4.2.17 IFM of the pit depth at AM surface, polarized to +424 mV vs Ag/AgCl at 25^◦C in 3.5wt% NaCl. . . 58

4.2.18 IFM of the pit depth at TM surface, polarized to +177 mV vs Ag/AgCl at 25^◦C in 3.5wt% NaCl. . . 58

4.2.19 IFM of the pit depth at ATM surface, polarized to +240 mV vs Ag/AgCl at 60^◦C in 3.5wt% NaCl. . . 59

4.2.20 IFM of the pit depth at AM surface, polarized to -10.9 mV vs Ag/AgCl at 60^◦C in 3.5wt% NaCl. . . 59

4.2.21 SEM image of pit at AM 316L flange, 60^◦C . . . 60

4.2.22 Pit on the AM 316L flange sample after CPP measurement in 3.5wt% NaCl at 60^◦C, polarized to +366 mV vs Ag/AgCl. . . 61

4.2.23 Second pit on the AM 316L flange sample after CPP measurement in 3.5wt% NaCl at 60^◦C, polarized to +366 mV vs Ag/AgCl. . . 63

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4.2.24 2000x magnitude of SEM image inside of pit after CPP of TM 316L at

60^◦C. . . 64

4.2.25 EDS spots for 316L flange, 60^◦C. . . 65

4.2.26 A second pit on the TM 316L sample after CPP measurement in 3.5wt% NaCl at 60^◦C, polarized to -20 mV vs Ag/AgCl. . . 67

4.2.27 CPP and OCP curves of 718 samples, TM and AM, measured in Avesta cell at 60^◦C in 3.5wt% NaCl. . . 69

4.2.28 CPP and OCP curves of 718 samples, TM and AM, measured in Avesta cell at 90^◦C in 3.5wt% NaCl. . . 69

4.2.29 Images from optical microscope of 718 after cyclic potentiodynamic polarization in Avesta cell, at 60^◦C and 3.5wt% NaCl. . . 70

4.2.30 Images from optical microscope of 718 after CPP in Avesta cell, at 90^◦C and 3.5wt% NaCl. . . 71

4.2.31 CPP and OCP curves of 718 samples, TM and AM, measured at 25^◦C in 3.5wt% NaCl. . . 72

4.2.34 OM images of AM 718, (a)-(c), and TM 718, (d)-(f), after CPP measured at 60^◦C in 3.5wt% NaCl. The corrosion attacks are marked and counted using ImageJ. . . 74

4.2.35 OM overview of test samples after CPP in 3.5wt% NaCl at 90^◦C. . . . 74

4.2.36 Investigation of pit depth after ASTM G61 test of AM 718 in 3.5 wt% NaCl at 60^◦C, polarized to +790 mV vs Ag/AgCl. . . 75

4.2.37 Investigation of pit depth after ASTM G61 test of TM 718 in 3.5 wt% NaCl at 60^◦C, polarized to +397 mV vs Ag/AgCl. . . 76

4.2.38 Investigation of pit depth after ASTM G61 test of AM 718 in 3.5 wt% NaCl at 90^◦C, polarized to +693 mV vs Ag/AgCl. . . 76

4.2.39 Investigation of pit depth after ASTM G61 test of TM 718 in 3.5 wt% NaCl at 90^◦C, polarized to +198 mV vs Ag/AgCl. . . 77

4.2.40 SEM image of pit at TM 718, 60^◦C . . . 77

4.2.41 Image area of TM 718 for EDS, 60^◦C . . . 78

4.2.42 SEM image of a pit on TM 718, initiated at polarization to +198 mV vs Ag/AgCl in 3.5wt% at 90^◦C. . . 80

4.2.43 Image area on TM 718 for EDS, 90^◦C . . . 81

4.2.44 SEM images of AM 718 pit at 90^◦C . . . 83

4.2.45 Area of the first pit at AM 718 for EDS, 90^◦C . . . 84

4.2.46 SEM images of an AM 718 pit at 60^◦C . . . 86

4.2.47 Image area of AM 718 for EDS, 60^◦C . . . 87

4.2.48 Open circuit potential development as a function of time and temperature of 718 samples with epoxy. . . 88 4.2.49 OM images of 718 samples with epoxy after modified ASTM G48 test. 89

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4.2.50 Open circuit potential development as a function of time and temper-

ature on 718 samples . . . 90

4.2.51 Open circuit potential development as a function of time and temperature on 316L samples . . . 90

4.2.52 OM overview images of 718 samples after ASTM G48 . . . 91

4.2.53 AM 718 sample with green growth after ASTM G48 . . . 91

4.2.54 OM overview images of 316L samples after ASTM G48 test . . . 92

4.3.1 Average ∆E = E_rep - OCP of the AM 316L and TM 316L corrosion tests with standard deviation. . . 95

4.3.2 CPP curve of SLM 316L and wrought 316L by Sander et al. . . 96

4.3.3 Average ∆E = E_rep - OCP of the AM 718 and TM 718 corrosion tests with standard deviation. . . 97 G.1 Area of first pit at AM 718 for EDS, 90^◦C . . . XIII

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List of Tables

2.2.1 Advantages and disadvantages with the use of AM technology. . . 15 3.1.1 Effect on surface roughness at SS 316L due to different values of laser

energy density, tested with SLM . . . 22 3.2.1 Summary of the effect of AM variables that influences the corrosion

properties of the built material . . . 32 3.2.2 Summary of some corrosion tests performed on AISI 316L and Alloy 718 33 3.2.3 Summery of some corrosion tests performed on Alloy 718 and Alloy 625 34 4.1.1 Alloy composition and PREN of the different test materials, both AM

and TM . . . 35 4.1.2 The processing parameters of the AM 316L from AIDRO . . . 36 4.1.3 The processing parameters of the AM 316L flange given by the manu-

facturers. . . 37 4.1.4 Processing parameters of the AM 625 material. . . 37 4.1.5 The processing parameters of the AM 718 given by the manufacturers 38 4.1.6 Overview of test concentrations and temperatures . . . 43 4.2.1 The lowest and highest surface porosity% found with ImageJ by rep-

resentative OM images of the AM 316L flange and AM 718 polished samples. . . 47 4.2.2 Parameters obtained from the anodic curve of the CPP measurements

of TM and AM 316L in Avesta cell under different conditions. The potential is in reference to Ag/AgCl. . . 52 4.2.3 Parameters obtained from the anodic curve of the CPP measurements

of 316L in the beaker setup under different conditions. The potential is in reference to Ag/AgCl. . . 55 4.2.4 The amount of localized corrosion attacks in area% found in ImageJ.

The number of attacks in the form of both crevices and pits were counted by the software and is presented. . . 57 4.2.5 The chemical composition of the AM 316L at each spot in Fig 4.2.22,

found through EDS analysis. . . 62 4.2.6 The chemical composition of the AM 316L sample in each spot in Fig

4.2.23, found through EDS analysis. . . 64 4.2.7 The chemical composition of the TM 316L sample in each spot in Fig

4.2.25, found through EDS analysis. . . 66

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4.2.8 Chemical composition in and outside a pit at TM 316 after CPP in 3.5wt% NaCl at 60^◦C . . . 68 4.2.9 Parameters obtained from the anodic curve from CPP measurements

of 718 in the Avesta cell. The potential is in reference to Ag/AgCl. . . 70 4.2.10 Parameters obtained from the anodic curve from CPP measurements

of 718 under different conditions in the beaker setup. The potential is in reference to Ag/AgCl. . . 73 4.2.11 The amount of localized corrosion attacks in %area found in ImageJ.

The number of attacks in the form of both crevices and pits was counted by the software. . . 75 4.2.12 The chemical composition of TM 718 in each spot in Fig 4.2.41, found

through EDS analysis. . . 79 4.2.13 The chemical composition of TM 718 in each spot in Fig 4.2.43, found

through EDS analysis. . . 82 4.2.14 The table shows the chemical composition of AM 718 in each spot in

Fig 4.2.45, found through EDS analysis. . . 85 4.2.15 The table shows the chemical composition of AM 718 in each spot in

Fig 4.2.47, found through EDS analysis. . . 88 4.2.16 Overview of the amount of corrosion attacks on TM and AM 316L and

718 during the critical pitting temperature test, ASTM G48. . . 92 G.1 Chemical composition found on the extra TM 718 pit at 90^◦C. . . XIV

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Abbreviations

Ag Silver

AgCl Silver chloride

AISI American Iron and Steel Institute

AM Additive Manufacturing

ASTM American Society for Testing and Materials

CE Counter electrode

CNC Computer numerical control

CPP Cyclic potentiodynamic polarization CPT Critical pitting temperature

Cr Chromium

CRA Corrosion resistant alloys DLD Direct laser deposition

EBM Electron beam melting

EBSD Electron backscatter diffraction EDS Energy-dispersive X-ray spectroscopy

H Hydrogen

IFM Infinite Focus Microscope

LM Laser melting

L-PBF Laser powder bed fusion

LS Laser sintering

Mn Manganese

MnO Manganese oxide

MnS Manganese sulphide

Mo Molybdenum

MPB Molten pool boundaries

NaCl Sodium Chloride

O Oxygen

OCP Open circuit potential

OM Optical microscopy

PBF Powder bed fusion

SEM Scanning electron microscope

Si Silicon

SiC Silicon carbide

SLM Selective Laser Melting

SS Stainless steel

TM Traditional manufactured

WE Working electrode

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Symbol list

α Ferrite

γ Austenite

ω Energy density

∆E difference between repassivation potential and OCP E_pit Pitting potential

E_rep Repassivation potential i_pas Passivation current R_a Surface roughness

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Chapter 1 Introduction

1.1 Background and motivation

Additive manufacturing (AM), commonly known as 3D-printing, is an emerging technology with several benefits over traditional manufacturing (TM). The process is predicted to be a faster, easier, and more economic alternative for several industries in the future, especially for the airplane industry, but also within oil and gas. The manufacturing techniques are experiencing increased attention for their inherent advantages, resulting in growing application and development [1]. These advantages are utilized in, e.g., production of patient-specific implants, complex and structurally optimized parts that lead to performance-critical weight savings, or reparation of expensive metallic parts [2]. In comparison, TM, such as CNC machining or milling, can have up to 95% of the original material removed to create the product, due to billets, forgings or castings [3]. The digital model for AM can be used many times without any wear and re-used with additions or changes. The main material focus of the AM methods is usually bio-compatible, high temperature and lightweight material such as Ti-, Ni-, Al- and Mg-based alloys [2].

We are just starting to see the potential of AM, but have still much to learn about the different properties of the AM metal compared to the well-known TM metal. As the production methods are so different, it is natural that the microstructure, hence the properties, are different. At this stage, it is still a challenge to produce metal from AM that is defect-free. To achieve this, we need to know more about the feedstock materials, process parameters, structure, properties, and performance [4]. Post-treatments have been utilized to increase the mechanical properties, but there is still a scarcity of studies focusing on the corrosion properties of AM metal.

The main object for this work is, thus, to document the pitting corrosion properties for AM corrosion resistant alloys (CRA) compared to the respective TM CRA at different temperatures in chloride concentrated water solution. By repeating the experiment under the same conditions, the aim is to state corrosion property differences yielded by the

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production method. In order achieve this, available literature about other experiments of the same nature will be studied to compare with the following experimental work to be performed. The effect of process parameters will be discussed in order to find optimized parameters for pitting protection and repassivation. The critical pitting temperature (CTP) will also be found to see if AM metal can replace TM metal in environments with higher temperature. Furthermore, the surface and corrosion attacks will be characterized to understand the nature and possible causes of the corrosion attacks. As AM metals could be favourably replacing TM steel in the future, it is essential to know if it can be done without downgrading the function and properties of the installation in use. AISI 316L and Alloy 718 are chosen as test materials due to their strong mechanical and corrosion properties and versatile use in marine environments [5].

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Chapter 2 Theory

Relevant literature was surveyed during the specialization project work [6], on which this thesis is a further work. The following chapter therefore presents the same theory, although with some changes and additions.

2.1 Corrosion

Corrosion is a big economical problem and is estimated to cost approximately 5% of an industrialized nation’s income due to prevention, replacement or maintenance of products [7]. Therefore, both money, time and environment can be saved by developing metals that are corrosion resistance. The corrosion resistance improves by adding chromium, iron, nickel, titanium, and many of their alloys, which creates a phenomenon called passivity, making normally active metals and alloys lose their chemical reactivity [7]. These types of metals are often referred to as stainless steels (SS), but will still be susceptible to corrosion in some environments. The degree of corrosion is not only affected by the composition, but also by fluid compositing, velocity and temperature [7].

2.1.1 Corrosion resistant alloys

The surface treatment and corrosion resistance of metals are related to their respective oxide layers. An oxide layer is a chemical compound containing oxygen and one additional element from the metal, for example, iron. Dens and tightly bound oxides provide strong protection towards aggressive ions and corrosion will be prevented. Aluminium and titan are examples of metals that form oxide layers so dense they can be considered corrosion-resistant, while iron needs alloying elements to create this effect and become a corrosion-resistant alloy, later referred to as CRA.

Regular steel will create a complex hydroxide (Fe₃O₄·nH₂O) in humid environment. This layer is highly porous, hence can O²⁻easily diffuse through the surface, and the metal will keep corroding [8]. CRA, such as SS, forms a different kind of oxide layer due to chromium,

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which is above 11% of the composition wt%. Chromium will, in the presence of oxygen, create a diffusion-tight and self-repairing passive film on the surface. The oxide layer, hence passive film, contains chromium (Cr) and oxygen, forming an impervious stable oxide layer (Cr2O3, called chromia) along the grain boundaries and surface [8]. Corrosion attacks are usually the cause of defects on grain boundaries due to high energy sites, unless they are protected by passivation. This oxide is only similar to the oxide of ordinary steel at high temperatures. The layer will be more complex at lower temperature, e.g, as hydrated C₂O₃ or as an outer layer consisting of hydroxides and water on the surface of an inner layer, consisting of Fe-Cr-oxides [8]. The thickness of the film is stable around 1.5-3 nm. This passive film is not impenetrable for chlorides, hence are SS also exposed to corrosion under conditions with aggressive ions.

To prevent corrosion in aggressive environments further alloying elements are added, such as Mo, Mn, Ni, etc. The addition of Mo makes the creation of the passive film easier, meaning it can repassivate easier during attacks, and is therefore more stable. An increasing amount of Cr, Mo, and N gives increasing resistance to pitting corrosion. By adding Ni with Mo will assure corrosion resistance in sulfuric acid and saltwater. This is further explained in section 2.1.5.

A wide range of mechanical properties combined with excellent corrosion resistant make SS very versatile in their applicability [7]. SS are divided into three classes based on the predominant phase constituent of the microstructure: ferritic, composed of the α-ferrite phase (BCC), austenitic, where the austenite phase, γ, (FCC) field is extended to room temperature or martensitic [7]. Duplex can also be registered as a class, with both γ and α phase [9]. The austenitic SS is the most corrosion-resistant due to high chromium contents and nickel additions. Thus, austenitic SS is produced in the largest quantities.

Austenitic steel

Austenitic SS possesses austenite as its primary crystalline structure, which is achieved by the addition of nickel, manganese, and nitrogen. These stabilizing alloy elements also have FCC structure [10]. The chromium content is high, with 17-18wt% [2]. There are two under groups of austenitic steel, 200 group and 300 group materials, where 316 and 316L are the most common SS after 304 and are characterized by the addition of Mo. The ”L” stands for low carbon grade, which provides greater resistance to carbon precipitation at high temperatures [9].

Inconel

Inconel alloys fall under the category called superalloy, referring to their superlative combinations of properties. Most are used in environments that expose the metal to severe oxidation and high temperatures for reasonable time periods, such as in aircraft turbine components, nuclear reactors or petrochemical equipment [7]. The trade name Inconel defines nickel-based alloys which are produced to high-temperature corrosion resistance, toughness and strength [11]. The alloys vary widely in their compositions but

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are primarily nickel-based with chromium as the second element. Other alloying elements include, as the other superalloys, the refractory metals (Nb, Mo, W, Ta), and titanium [7].

Alloy 718 is age hardenable and the most widespread material in its class [12]. It has been the standard material for turbine discs of gas turbine engines manufacturing for many years and hold outstanding mechanical properties even at temperature ranges of 250-700^◦C. The extraordinary properties come to form the unique microstructure, consisting of an austenitic matrix, γ, precipitates γ”, γ’, δ and carbides. The γ matrix is a solid solution of Cr, Fe, Ni and Mo and has FCC crystal structure, while the γ” is a metastable phase (Ni₃Nb) with tetragonal space centred crystal structure and is the main strengthening phase, typically 15-20% of the volume. Exposure to high temperature over time transform this phase to δ phase, which precipitates at the grains boundaries and prevent grain growth [12, 13]. Therefore, this phase transformation is a positive effect on the mechanical properties, strengthening the structure and prohibiting brittle behaviour and fatigue cracks to induce [13].

Passive metals, such as SS and Ni-alloys, are susceptible to localized corrosion due to the breakdown of the passive film. This breakdown causes an accelerated dissolution of metal at the localized sites that are unprotected. Typical forms of localized corrosion are pitting and crevice corrosion, affected by aggressive anions and temperature of the electrolyte as well as the composition and nature of the metal [14].

2.1.2 Pitting corrosion

Pitting initiates on SS in the formation of small areas or points taking form as cavities due to slow repassivation of the passive film. They can take form as shallow and wide or deep and narrow, as illustrated in Fig 2.1.1. Aggressive anion species, e.g., chloride, damages this ultra-thin passive film causing exposure of the underlying metal to the electrolyte which allows corrosion. Shallow and short-lived pits are called metastable pits. These are pits that are stopped by repassivation of the passive film but are still important as they can re-initiate and grow to form further breakdown [15]. When stable pits are formed and grow, the propagation stage of pitting corrosion is reached, often referred to as the pitting potential.

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Fig. 2.1.1. Different shapes of pitting corrosion [16].

In a chloride solution, anodic dissolution of the metal is according to Equation 2.1.1 and happens at the bottom of the pit [14]. The reaction is balanced by the cathodic reaction of the water on the surface of the metal, shown in Equation 2.1.2 [17]. The rate of pit growth is therefore dictated by the pit chemistry, growth and potential at the pit bottom [18].

M →Mⁿ⁺+ne⁻ (2.1.1)

O2 + 2H2O+ 4e⁻ →4OH⁻ (2.1.2)

M⁺Cl⁻+H2O →M OH +H⁺Cl⁻ (2.1.3)

Pitting corrosion is aggressive and rapid, making repassivation inside a growing pit highly unlikely. The reason is due to metal ions dissolved in the surrounding environment by the corrosion. The pH decreases as the negative chloride ions migrate to the pit and the positive metal ions to maintain balance [17, 18]. As a result, the environment inside the pit becomes more and more aggressive, shown in Equation 2.1.3, causing a reaction of the metal chloride to form into metal hydroxide and free acid.The solution on the metal surface where the cathodic reaction occurs will at the same time experience increased pH [15]. This is illustrated in Figure 2.1.2.

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Fig. 2.1.2. Schematic representation of the environment inside a pit [14].

An existing pit can repassivate if the alloy contains a sufficient amount of the alloying elements Cr, Mo, Ti, N, V, etc. Mo will significantly enhance the enrichment of Cr in the oxide but these alloying elements will in general help to stabilize the passive film and induce rapid repassivation to heal a damaged area in case of breakdown. The com- positional effect in the alloy can be represented by PREN; Pitting resistance equivalent number

P REN =Cr% + 3.3(M o%) + 16(N%) + 1.65(W%) (2.1.4) This number can be used as an indication of pitting and crevice corrosion resistance, i.e., a PREN of 40 indicates that a metal can withstand localized corrosion in deoxygenated seawater at 20^◦C [19].

Investigation of localized corrosion usually starts with electrochemical measurements to study the passivity of the material. Cyclic potentiodynamic polarization (CPP) is one of the techniques used, measuring the cathodic polarization first and then the anodic polarization [20]. Open circuit potential (OCP) is often measured before CPP to find the potential where the cathodic and anodic reactions on the metal surface are stabilized and therefore knowing where to start the measurements of CPP. The shape of the anodic polarization curve gives information about passivity and breakdown of the passivity of metals under different conditions [18].

The anodic polarization starts from the corrosion potential after reaching the steady-

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Fig. 2.1.3. Schematic illustration of a CPP curve with the corrosion parameters [21].

state. The curve then goes from active corrosion to a stable increase in potential, which is the passive region, seen in Fig 2.1.3. The passive current density (i_pas) will remain stable as long as the metal is passive. A rapid increase in current density indicates that pitting initiates. This increase can repassivate, due to metastable pitting, or continue increasing, showing pit growth. The potential when pitting corrosion overcomes the passive layer, is called E_pit. Rapid increase at potential above 1.2 V vs SCE can also indicate the evolution of oxygen in the water, which then is called the transpassive region. The oxygen can therefore in favour be removed to ensure that the oxygen evolution does not disturb the curve for localized corrosion [18, 21]. Repassivation (E_rep) of the oxide is seen when the turning curve stabilizes in current density or intersect with the forward curve. No pitting corrosion is initiated under this point. Szklarska-Smialowska explains the character of the CPP curves in further details [18]. In order to determine an accurate potential, an electrode with a well-known and stable potential is used as reference, often SCE or Ag/AgCl.

There is no standard curve for each material as they change with different variables, such as temperature, chloride content, and pH. Szklarska-Smialowska [18] states that scatter of the results in the determination of the pitting potential is sometimes very high on several samples of the same materials, thus is independent of the kind of metal, alloy or test method. Most authors consider this dispersion due to the mechanism of localized corrosion, either to random probabilistic phenomena or to instability. Pitting corrosion is

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deterministic, but at the same time susceptible to many experimental parameters, making reproducibility difficult to achieve [18].

2.1.3 Crevice corrosion

Crevice corrosion is another form of localized corrosion, occurring in narrow gaps between two metals, one metal with another material, badly executed welds, etc. As for pitting, crevice often occurs in chloride solutions, but also in other corrosive liquids. Aggressive ions penetrate and create a higher pH level underneath the gap than outside, as seen in Fig 2.1.4. The intensity of the attack is dependent on the width and length of said gap. Differential to pitting corrosion, which initiates different vulnerable sites on the metallic surface, crevice corrosion only attacks a portion of the metallic surface to which the aggressive environment is limited to the geometric [18]. Pitting is often seen as a particular case of crevice corrosion, where they act in a similar matter between small gaps on the surface, while other authors state that crevice propagates from pits inside the gap. However, crevice corrosion has lower crevice potential than pitting corrosion and, in most cases, initiates more rapidly [18].

Fig. 2.1.4. Schematic illustration of crevice corrosion in chloride solution [22].

2.1.4 Effect of aggressive ions and pH

There are several parameters influencing localized corrosion, both affecting critical corrosion parameters but also affecting the interpretation of results. An electrolyte containing ions, such as Cl⁻, will be an aggressive environment for the metal in direct contact because of the attrition of the oxide film on the material surface. As previously stated, it is the chloride that causes the suffering of localized corrosion as pits generate on the surface.

Metal in a diluted electrolyte with a small concentration of chloride ions will have a high

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critical pitting potential. This potential decreases with increasing chloride concentration, meaning pitting corrosion will more easily initiate in electrolytes with higher chloride contents [21]. Other ions, such as F⁻ or Br⁻, have the same aggressive behaviour on the passive film as chloride. The chloride content does not always cause pitting corrosion, but decrease the stability of the passive film by lowering the breakdown potential.

Because of the hydrolysis occurring during the anodic dissolution of the metal or alloy, acidification of the solution is contained in the pits or crevice as they develops [18]. Low pH in the solution will more naturally cause lower and more aggressive environments inside propagating pits, causing more bottomless pits, as illustrated in Fig 2.1.5, in addition to easier initiation and lower pitting potential.

Fig. 2.1.5. Maximum pit depth vs chloride concentration for AISI 316L in chloride contained solution in various pH values, static (S) and dynamic

(D) conditions [5].

It is possible to add corrosion inhibitors to the electrolyte to reduce the influence of aggressive ions, such as Na₂SO₄. Inhibitors will increase pitting potential.

2.1.5 The effect of metallurgical variables

As previously stated, the metallurgical variables affect the passive film, its stability, and where the corrosion attacks the material. Precipitation at grain boundaries, impurities by atoms, or inclusions are some factors [21]. The pit’s morphology and the composition of the metal surface, thickness, and the composition and structure of the passive film can all influence the pit nucleation. Nickel-based alloys can have the corrosion resistance increased by selectively alloying. The major elements for enhancing resistance for pitting corrosion are chromium, molybdenum, tungsten and nitrogen [23]. The strong influence of Mo is shown in Fig 2.1.6 and 2.1.7, giving increased corrosion resistance and higher

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repassivation with increasing content [18]. Increasing Mo concentration of the bulk material also increases the concentration in the passive film, which has shown to increase film thickness. By decreasing the susceptibility of pitting, Mo also increases the metal’s repassivation abilities [23].

Fig. 2.1.6. Pitting potential vs the content of Mo in the composition at different temperatures [18].

Fig. 2.1.7. The effect of pitting potential for 15 wt% Cr, 13 wt% Ni SS by increasing Mo, exposed in 0.1 M NaCl at 25^◦C [23].

Increasing the Cr content gives further Mo contribution to pitting resistance as the Cr

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compliments Mo. Cr in general also contributes to increased pitting resistance, seen in Fig 2.1.8. Another contribution in pitting potential is also given by nitrogen, which makes the microstructure denser and prohibit segregation at the grain boundaries [8]. The effect of N can be seen in Fig 2.1.9. However, a large concentration of nitrogen cause intergranular Cr₂N which can cause corrosion. Nickel has often been substituted for N in recent years to prohibit this affect [24]. Also, increasing the Cr content can promote precipitates of σ in CRA which reduces the corrosion resistance to intergranular, crevice and pitting corrosion [24]. Pitting will happen at the points with low Cr content, which often is at the austenite grain boundaries. The ferrite and austenite boundaries are also preferential sites to form σ, inducing galvanic effect that corrodes the austenite [24]. Small amounts of titanium and niobium are frequently present in CRA but are of a greater significance in enhancing resistance to intergranular corrosion rather than pitting [23].

Fig. 2.1.8. Critical pitting potential for Ni-Cr alloys in 0.1 M NaCl at 25^◦C [23].

Fig. 2.1.9. Pitting potential dependence of nitrogen for austenitic SS containing 22 wt% Cr, 20 wt% Ni, 4 wt% Mn and 0.1 or 2 wt% Mo [23].

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On the other hand, CRA often has some sulfur content. S is detrimental for pitting due to formation of sulfur-rich inclusions that dominate as pitting nucleation sites and, therefore, the lifetime of the unstable pits is related to the particle size of the sulfur [18].

2.1.6 Effect of temperature

Temperature is a critical factor in pitting corrosion as many metals will not initiate or propagate pits below a certainty temperature [15]. This temperature is commonly known as the critical pitting temperature (CPT). A high critical temperature corresponds to high corrosion resistance as increasing temperature weakens the passive layer. Thus, high temperature also affect the repassivation of the film. The effect of pitting corrosion is illustrated in Fig 2.1.10, showing additional increase in i_pas. Critical crevice temperature (CCT) can also be found and is usually lower than the corresponding CPT [15].

Fig. 2.1.10. Typical cyclic potentiodynamic polarization curves for CrMn stainless steel in buffer solution of pH 8 with 0.5 Cl⁻ showing the effect of

temperature [25].

2.1.7 Effect of dissolved gases in the electrolyte

The atmosphere the specimen is tested in profoundly influences its behaviour during the CPP test. Dissolution of different gases, such as O₂, H₂, and CO₂, influence the corrosion behaviour, formation (thus also reformation), of the passive film, and pitting corrosion [21]. Hydrogen has destructive effects on high alloyed steel and will cause a decrease in both E_pit and E_rep, a result of a disruption in the stability of the passive film [21]. The concentration of dissolved oxygen can have two different effects on the corrosion process.

The first case being the limitation of the corrosion rate as oxygen is consumed in the

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cathodic reaction. An increase in oxygen will, therefore, increase the corrosion rate. The second effect is on the oxide film, where the presence of oxygen reduces the corrosion rate as it raises the passivation [26]. However, there is a lack of studies on the effect of oxygen on CRA regarding localized corrosion. The effect of deaerated and open-air conditions was tested for SS, which showed that oxygen in a system with materials that are susceptible for localized corrosion may evolve concentrated cells that can cause pitting or crevice corrosion [27]. Feng et al. [28] also studied the oxygen and nitrogen effect on the corrosion resistance of 316L, as some authors claim that deaerated electrolyte with nitrogen can lower the corrosion potential. However, Khobragade et al. [27] found that deaerated electrolyte at different concentration of Cl⁻ all showed higher Epit that in open-air. The results are shown in Fig 2.1.11, revealing a low difference in the localized corrosion potential.

Fig. 2.1.11. Potentiodynamic polarization curve of 316L in borate buffer solution at different exposure of oxygen [28].

Oxygen evolution may replace the corrosion reaction at high potential, interfering with the results. As already mentioned, this causes transpassive behaviour of the material, which will give inaccurate results in the CPP curves [21].

2.1.8 Effect of surface roughness

The condition of the surface can have a significant influence on the behaviour of pitting corrosion on the metal [15]. Pits initiate at specific sites on the surface, such as roughness and inclusions [21]. The stabilization criteria explain the relation of surface roughness and pitting. As roughness gives more occluded sites, these can sustain the conditions required for active dissolution at lower current densities and lower potentials [15]. Studies on 316 SS have shown that pitting potential can be changed by as much as 500 mV using different surface preparation methods [29, 28]. Changing the grit size from 360 to 2200

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to produce finer surface finish can result in 200 mV increase in pitting potential in 0.1 M NaCl at RT [28]. The same effect is shown for other SS, removing groves, crevices and related defects that act as a form of preexisting defects where acidification can occur which exhibit metastable pitting as well. Higher surface roughness also exposes inclusions which increases the number of sites for pit initiation [28].

2.2 Additive Manufacturing

Parts made from additive manufacturing (AM) are designed as a three-dimensional digital model that is virtually sliced down to thin layers, usually around 20µm-1mm, dependent on the AM method used [30]. The physical three-dimensional component is then printed based on this model by single layers of material that are locally melted by a heating source. Therefore, in contrast to traditional manufacturing (TM), AM is carried out by thin layers of material in an additive process, hence the AM term [31, 32]. Special tools used for TM, such as cutting tools or molds, are not needed to make AM parts as the layers are built directly on a surface or built platform. Advantages and disadvantages are summarised in Table 2.2.1.

Tab. 2.2.1. Advantages and disadvantages with the use of AM technology.

Advantages Disadvantages

Complex geometries at no or small extra cost Unknown properties of new materials

No material waste Rough surface finish

Lighter parts with no excess material Cost of new machines and knowledge No tool cost Higher cost if the quantity of parts is high Faster from design till finished product

Strong material properties

Reparation possibilities of other products

The AM processes for metal can use two different types of feedstock: wire or powder.

The feed is selectively and continuously melted by a heat source and then cooled to form a surface. The processes can be divided into two categories: Powder Bed Fusion (PBF), including Laser Beam Melting (LBM) and Electron Beam Melting (EBM), and Direct Laser Deposition (DLD)/Direct Laser Deposition (DLD) [33]. Other methods also exist, as seen in Fig 2.2.1, but PBF and DED are the most versatile for producing AM components today. However, the demand for Binder Jetting (section 2.2.4) processed metal is expected to increase. The names of the processes are usually trademarked from the different manufacturers; LBM is, for example, also known as Selective Laser Melting (SLM), Direct Metal Laser Sintering (DMLS), LaserCUSING, Laser Metal Fusion or industrial 3D printing [30]. For simplicity, this process will be referred to as SLM further on.

As powder is mostly used as the feedstock and laser as the power source in AM processes, laser sintering, laser melt, and direct laser melting are the three methods of physically

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Fig. 2.2.1. Overview of the additive manufacturing technologies used to produce metal. The information is retrieved from [32].

forming the powder to a material. To produce an AM component with favourable microstructure and properties, design strategy, process parameters, and powder material are essential. It is usually difficult to tailor the microstructure and mechanical properties as these are strongly dependent on process and material, hence powder characteristics.

This includes particle size, particle shape, distribution, flowability, packing, etc. The process variables set apart the different melting or sintering methods and involve the laser type, spot size, scan speed, laser power, scan line spacing, and powder layer thickness [34]. Therefore, a distinct difference between AM processes have been apparent for the produced metallic components.

2.2.1 Laser sintering

Laser sintering (LS) uses layer by layer powder spreading that is sintered by a laser. It also contains an automatic powder layering apparatus, a computer system for process control, and inert gas or preheating of the powder bed. The mechanism for the powder densification depends on the wavelength from the laser, so the different types of lasers will, therefore, affect the powder in different ways [34]. The process begins with a lowered building platform. An inert gas, nitrogen or argon, is fed to the sealed building chamber to reduce oxygen, which contaminates the metal in production because of oxidation.

Then, usually, a 100µmeter thick layer of loose powder is deposited on the surface of the build platform by a recoater arm, and the laser starts to scan the powder bed surface according to the CAD data. The recoater arm then lays a new layer of powder as the process is repeated [33, 34, 35], as seen in Fig 2.2.2.

The process is based on liquid phase sintering, meaning partial melting of the powder.

The mechanism has short and rapid thermal cycles where the temperature of the laser is determined between the highest melting point of the metallic component, acting as the structural metal, and the lowest melting point of the metallic component, acting as the binder. Gu et al. [34] refer to this temperature as the mushy zone. The metal working as the binder will melt completely while the structural metal stays solid in the liquid.

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Fig. 2.2.2. Schematic representation of DMLS and SLM [36].

The wetting liquid causes capillary forces on the solid particles, which determines the rearrangement of the particles, hence the result of the metal.

Optimal incongruent melting during the sintering process requires strict control of the laser processing parameters. The localized rapid nature of the thermal cycle during the process creates a significant difficulty in controlling the temperature between the solid and liquid phases. Typical processing problems with pre-alloyed powder are insufficient densification, heterogeneous microstructures and properties, etc. However, it is possible to obtain sufficient mechanical properties by post-treatment such as furnace post-sintering, hot isostatic pressing (HIP), or secondary infiltration with a material with a lower melting point [34].

DMLS is a typical processing method used for PBF-LS, which is, as mentioned, the same process as SLM but with different laser parameters. DMLS needs less energy and lower temperature as the particles are only heated, while SLM needs a higher temperature to completely melts the particles. These thermal parameters make DMLS a better option for metal alloys, while SLM works best in producing pure metals. DMLS and SLM can produce SS (AISI 316L), aluminum, titanium, nickel alloys (Inconel), cobalt chrome, and precious metals [35].

2.2.2 Laser Melting

Laser melting (LM) has been developed as a process to produce fully dense components with mechanical properties that, in theory, could be comparable to those of bulk material without post-treatments. As mentioned, LS and LM have the same process and apparatus, differential by a cycle of complete melt and solidification of the metallic powder in the LM process. The idea of a full melt has required improved laser processing condi-

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tions such as higher laser power, smaller focused spot size, and thinner layer thickness, causing a more time-consuming process. However, results show that microstructural and mechanical properties indeed are significantly improved compared to early time LS processed metal. The density of the LM processed metal is highly controllable and can be produced up to 99.9% in a direct way. Nevertheless, the LM process suffers a high risk of instability of molten pools due to the complete melting. As the powder solidifies it also shrinks, causing a build-up of residual stress during cooling that gives distortion and even delamination of the final metal. Also, the instability of the melt may cause spheroidizing of the particles, which leads to internal porosity. Therefore, both processing and powder parameters are fundamental to create a desirable result.

2.2.3 Laser Metal Deposition

The laser metal deposition does not involve a powder bed and can be seen as the AM way of welding. The material is added to the melt pool by a wire or powder feed system while a heat source, typically laser, creates the melt pools by heating the surface [37]. Gas delivers the powder more efficient and limits oxidation and contamination; the process is seen in Fig 2.2.3.

Fig. 2.2.3. The figure shows the apparatus and process of LMD [37].

By adding the material directly to the surface it is possible to repair parts, scratches, wear, or to prepare corrosion or wear-resistant coating, saving the cost of buying new parts that previously were unrepairable. Also, rebuilding a worn part may result in the repaired component having a longer life than the new part would have [34]. The LMD also gives an unique possibility of depositing multiple materials at different parts on one component with high precision. This can be utilised to the development of new materials with properties and characteristics that do not exist in natural environments. To be able to do this, or make a favorable design with LMD, knowledge of temperature, velocity, and

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composition distribution is important as these parameters are essential for the process, microstructure, and properties.

2.2.4 Binder Jetting

Binder Jetting may become a strong competitor to the popular SLM production. Binder Jetting can have several inkjet whilst SLM only has one laser, providing the possibility of a faster process than for SLM. It is in addition a cheaper process method. The process uses an inkjet print head to quickly deposits a bonding agent onto a thin layer of metal powder, binding the particles together. This process is repeated, layer-by-layer, to form the component. After this stage, the component is called a ”green” part which is further cured or dried in an oven so the redundant material is removed. The final stage is sintering at high temperature to fuse the particles together. The process is shown in Fig 2.2.4.

Further information about Binder Jetting is reviewed by Ziaee et al. [38] or can be found at ExOne’s homepage [39].

Fig. 2.2.4. A full image of the machine used in Binder Jetting [40].

Binder Jetting today is mostly applied in the production of powdery materials for casting, medicine and electronics [38]. The new attention for metal Binder Jetting production come from the interest of companies and long-time users of other AM methods that want a faster and cheaper process. Binder Jetting compared to SLM and DMLS, operate at a consistent low temperature, which prohibited thermal stress in the component and therefore also stress related defects [39].

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Chapter 3 State of the art

3.1 Mechanical properties of AM

The layer-by-layer process is practically a repeating process of casting new layers of metal on top of each other. The alloy experiences repeated solid and liquid state phase transformation that differ from traditional wrought or cast metal production processes, which thus give different microstructure. A combination of rapid solidification, direc- tional cooling, and phase transformation induces the formation of unique alloy structures that typically have a refined grain structure, dislocation cell sub-structures, and residual stresses [33]. This structure gives higher hardness and yield strength from the finer grains and the enhanced solid solubility, coming from the rapid quenching, while the residual stress built up in the structure makes it susceptible for fracture, reducing the ductility [2].

The residual stress builds up as a result of the uneven heat distribution and shrinkage as the powder solidifies. The rapid solidification also reduces elemental diffusion and extend solid solubility, which can result in a metastable phase [41]. Microstructural banding is another possible defect caused by the repeated thermal cycles, giving microstructural differences between deposition layers [41]. The process can also cause metallurgical defects, entrapped gas porosity, dendritic growth withing the build, and lack of fusion porosity [33]. The presence of these defects or flaws can have a negative impact on the resulting properties of the produced component, regardless of the production methods. There- fore, it is of high importance that the processing parameters are constantly improved to minimize defects with influence.

3.1.1 Roughness

The surface roughness (R_a) of AM components produced by SLM is usually high. The R_a for components from milling is 1-2µm, while the roughness in SLM components can be in the range of 10-30 µm [33]. Wang, et al. [42] did a study on the effect on the roughness of SLM 316L SS due to laser energy density (ω), meaning the energy imparted

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by the laser per volume of the melt pool. The ω is defined by the Equation 3.1.1, where P_{ef f} is the effective laser power (W), υ_s is the hatch distance (mm) and d is the layer thickness (mm) [43]. The hatch distance represent the area affected by the laser, which together with the thickness representing the volume of the melt pool.

ω = P_{ef f}

υ_s·h·d⁰ (3.1.1)

The reports show high Ra below 75 J/mm³, decreasing in the range of 100-170 J/mm³ and again increasing above 180 [33, 42]. Low ω gave sporadic lack of melting, resulting in given surface roughness, while at high ω, the roughness possibly increases as a result of excessive melting. At ω between 75 J/mm³ and 120 J/mm³ the phenomena balling - the formation of metallic droplets instead of the desired spred of liquid to form uniform molten surface - occurred predominently [33]. Balling has been registered to be attributed to low laser power, large laser focus diameter and hence a low laser energy per unit area, in addition to contribution from poor oxygen control and coarse powder (diameter of 75µm) [2]. A low degree of melting, hence low ω or high scanning speed, can also cause balling on the surface of the melted powder [44].

Table 3.1.1 summarize the divided fabrication into not completed melting zone, successful fabrication zone, excessive melting zone and balling zone. Reportedly the upper surface roughness was affected by the width of the track, scan space, and thickness of powder layer [42]. These results show only the indication for process parameters for SS 316L using SLM, as other AM processes have other factors affecting the resulting component.

In addition, other alloys produced by SLM have different powder characteristics and may, therefore, give different R_a-ω relationship.

Tab. 3.1.1. Effect on surface roughness at SS 316L due to different values of laser energy density, tested with SLM

Laser energy density [J/mm³] Surface roughness [µm]

<75 15

not completed melting zone

100< ω <170 <10

successful fabrication zone

>180 >10

excessive melting zone 75< ω < 120 Balling effect

The energy density was suggested used as an identity parameter by Kamath et al. after Tolosa et al. in 2010 produced an L-PBF 316L SS with 99% relative density implementing 200 W laser with an 80µm laser focus diameter and scan velocity up to 1000 mm/s [2].

Furthermore, Liu, et al. [45] observed that residual stress has relation with ω and will decrease with lower ω. The study they used x-ray to determine residual stress in SLM

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316L, observing that it is mainly compressing at the lower build and tensile at the top of the build. It was also found that the residual stress is much larger in the parallel direction to the laser scanning direction. Furthermore, other studies have found a negative effect of surface roughness of as-built samples, which result in poor fatigue properties irrespective of other defects, highlighting the important role surface roughness for fatigue [46].

3.1.2 Microstructure

The SS produced by L-PBF is reported to be fully austenitic with columnar grain structure that has fine solidification cells with diameters of 1 µm or smaller [30, 2]. One austenite grain is formed by tens or hundreds of these cells, which are very similar in crystallographic orientation, i.e. high grain boundary binds the material volume. As previously stated, SS produced in L-PBF has smaller grains than TM components. The microstructure differences depends on processing methods, clearly viewed by comparing 316L SS from DED and L-PBF process. The component from DED shows microseg- regation at the regions along the borders of the solidification cells during solidification, which leads to an enrichment in the ferrite stabilising elements Cr and Mo, hence giving fine ferritic films. This is not observed in L-PBF as all studies observe fully austenite microstructure with no indication of any solid-state phase transformation [2]. However, the intercellular regions show enrichment with Cr and Mo, but not sufficient enough to stabilise ferrite. This is illustrated in Fig 3.1.1, showing the stated difference in conventional, L-PBF and DED microstructure. Furthermore, no Cr carbides have been observed at the grain boundaries due to the rapid cooling.

The microstructure of austenitic SS produced by L-PBF gives a combination of higher strength without a reduction in ductility. Wang et al. [47] attributed this property combination to the order of the microstructure, regarding solidification cells, high and low angle grain boundaries, dislocations, and oxide inclusions. Fig 3.1.2(a) shows the grain orientation of a L-PBF 316L SS where the colour and thus orientation change continuously within one grain. These Electron backscatter diffraction (EBSD) measurements suggest a broad grain size distribution in the SS microstructure, which together with the different grain boundaries and inclusions (Fig 3.1.2 c-e) show both structural and chemical heterogeneous microstructure [47]. This order is unique for AM and has also been reported in E-PBF of SS 316L [2]. Furthermore, conventional austenitic SS typically show strain-induced martensite formation upon plastic deformation, but the fine microstructure in L-PBF produced 316L hinder formation of martensite. Nevertheless, the metal possesses high yield strength but has low work hardening [48].

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Fig. 3.1.1. Overview of the austenitic SS microstructure by different processing methods. Depending on the processing conditions, different observations may occur. γ: austenite phase,δ: retained primary ferritic

phase [2].

Fig. 3.1.2. a, Various length scales uncovered in L-PBF 316L SS. b, EBSD showing grain orientationc, SEM image revealing fusion boundaries, high-angle grain boundaries (HAGBs), and solidification cellular structures.

The inset shows the cellular structure at a higher magnification. d, TEM image of solidification cells. e, A high-angle annular dark-field (HAADF) scanning TEM (STEM) image of the solidification cells shown in d [47]

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AM opens the possibility to control the cooling rates during and after solidification in each layer of the metal by varying process parameters. It thereby presents the possibility to local and digital control over the microstructure of the built component, a possibility that is not possible for TM [2]. To achieve a microstructure with strong mechanical and corrosion properties, the process parameters need to be optimized. As seen in Fig 3.1.3, the effect of just laser power is strong. Low laser power, as in Fig 3.1.3(a), creates small molten pools with low melt, resulting from poor track overlap. Large keyhole pores are therefore created between the layers. Fig 3.1.3(c) shows that high laser power also causes pores along with the molten pools, but more spherical. Even higher laser power (300 W in Fig 3.1.3(d)) can cause entrapped gas pores in addition to long and irregular tracks in the structure. Pores and molten pools are therefore seemingly correlated [49].

Fig. 3.1.3. OM side views of L-PBF 625 microstructure produced by laser power of a) 150 W, b) 200 W, c) 250 W and d) 300 W [49].

3.1.3 Porosity

Pores and their distribution on the metal parts are complex. Various process parameters, scanning and building strategies, feedstock material, deformation during manufacturing, etc, lead to a wide range of potential flaws, including different pore shapes, sizes and total volumetric porosity [46]. The effect of these defects has a similar range. Even with optimal parameters, errors or imperfections may still occur. X-ray tomography/Computed tomography (CT) scans allow a quick, non-destructive investigation of larger material volumes to inspect final parts based on defect sizes, such as pores and cracks, location of defects or the geometrical accuracy for the complex and internal features of the compo-

Pitting corrosion of additive manufactured AISI 316L and Alloy 718

Master ’s thesis

Pitting corrosion of additive

manufactured AISI 316L and Alloy 718

Pitting corrosion of additive

manufactured AISI 316L and Alloy 718

Contents

List of Figures

List of Tables

Abbreviations

Symbol list

Chapter 1

Introduction

1.1 Background and motivation

Chapter 2

Theory

2.1 Corrosion

2.1.1 Corrosion resistant alloys

2.1.2 Pitting corrosion

2.1.3 Crevice corrosion

2.1.4 Effect of aggressive ions and pH

2.1.5 The effect of metallurgical variables

2.1.6 Effect of temperature

2.1.7 Effect of dissolved gases in the electrolyte

2.1.8 Effect of surface roughness

2.2 Additive Manufacturing

2.2.1 Laser sintering

2.2.2 Laser Melting

2.2.3 Laser Metal Deposition

2.2.4 Binder Jetting

Chapter 3

State of the art

3.1 Mechanical properties of AM

3.1.1 Roughness

3.1.2 Microstructure

3.1.3 Porosity