Licentiate ThesisProduction Technology2020 No. 30

Laser-Directed Energy Deposition: Influence of Process Parameters and Heat-TreatmentsSuhas Sreekanth

Tryck: Stema Specialtryck AB, August 2020 Trycksak3041 0234

SVANENMÄRKET

Trycksak3041 0234

SVANENMÄRKET

Licentiate ThesisProduction Technology2020 No. 30

Laser-Directed Energy Deposition: Influence of Process Parameters and Heat-TreatmentsSuhas Sreekanth

University West SE-46186 Trollhättan Sweden +46 52022 30 00 www.hv.se © Suhas Sreekanth 2020

ISBN 978-91-88847-67-6 (Printed version) ISBN 978-91-88847-66-9 (Electronic version)

iii

Acknowledgements

I would like to express my sincere gratitude to my supervisors Joel Andersson and Shrikant Joshi for being a constant source of guidance. Their continual support and motivation have kept me challenging myself in this pursuit for learning.

I would like to thank all the research engineers at (Production Technology West) PTW for sharing their expertise. I have deeply benefited from their knowledge bank with a combined experience of about 100 years. In particular, working with Kjell Hurtig has been immensely beneficial in this learning process. I would also like to send huge appreciation to all my colleagues at PTW for their support and discussions while creating a safe learning environment.

I would like to acknowledge the research partners involved through SUMAN-Next and CAM2 project for their continual insights. Particularly, the inputs of Peter Harlin proved invaluable in detailing the scientific aspects of this work.

Finally, I would like to thank my friends and family for being there for me through the course of every sojourn.

Suhas Sreekanth 1st of August 2020

University West SE-46186 Trollhättan Sweden +46 52022 30 00 www.hv.se © Suhas Sreekanth 2020

ISBN 978-91-88847-67-6 (Printed version) ISBN 978-91-88847-66-9 (Electronic version)

iii

Acknowledgements

I would like to express my sincere gratitude to my supervisors Joel Andersson and Shrikant Joshi for being a constant source of guidance. Their continual support and motivation have kept me challenging myself in this pursuit for learning.

I would like to thank all the research engineers at (Production Technology West) PTW for sharing their expertise. I have deeply benefited from their knowledge bank with a combined experience of about 100 years. In particular, working with Kjell Hurtig has been immensely beneficial in this learning process. I would also like to send huge appreciation to all my colleagues at PTW for their support and discussions while creating a safe learning environment.

I would like to acknowledge the research partners involved through SUMAN-Next and CAM2 project for their continual insights. Particularly, the inputs of Peter Harlin proved invaluable in detailing the scientific aspects of this work.

Finally, I would like to thank my friends and family for being there for me through the course of every sojourn.

Suhas Sreekanth 1st of August 2020

iv

v

Populärvetenskaplig Sammanfattning

Nyckelord: Lasermetalldeponering; Legering 718; Processparametrar; Värmebehandlingar; Nb-rika faser

Lasermetalldeponering med pulver som tillsatsmaterial (L-DED) är en additiv tillverkningsmetod där man genom lager-på-lager-principen bygger upp en specifik geometri. Metoden har visat en enorm potential under det senaste decenniet. Flyg- och rymdindustrin står som de främsta mottagarna på grund av processens förmåga att bygga komponenter utan något större behov av omfattande skärande bearbetning vilket bidrar till mindre materialspill. Användningen av superlegering 718 inom flygmotorapplikationer är mycket omfattande vilket har lett till ett stort forskningsintresse gällande utveckling av L-DED-processen för specifikt denna superlegering. AM-processer pliktas ofta av låga bygghastigheter och långa ledtider vilket direkt påverkar produktionskostnaderna. För att övervinna låga bygghastigheter så har detta forskningsarbete fokuserat på att uppnå höga deponeringshastigheter genom högre materialtillförsel.

Inverkan av processparametrar vid L-DED processen är av största vikt eftersom de påverkar mikrostrukturen för det specifika materialet och hur materialet slutligen beter sig vid användning. I det aktuella arbetet varierades processparametrar såsom lasereffekt, framföringshastighet, pulvermatningshastighet och pulverfokusförskjutning där deras inflytande på geometri och mikrostruktur av ensträngsväggar gällande legering 718 har analyserats. Ensträngsväggarnas geometri har undersökts i form av höjd, bredd, inträngningsdjup och effektivitet gällande pulvermatning vilket bestämts genom att mäta tvärsnittsarean för deponering samt uppblandning. Deponeringens mikrostruktur visade kolumnär dendritisk struktur i mitten varav likaxlig dendritisk struktur i bottenregionen samt i det övre området. Segringar i form av Nb-rika faser som Laves och NbC kunde identifieras i de interdendritiska områdena samt i korngränserna. Segringen ökade längs deponeringshöjden, där bottenregionen hade den minsta och det övre området den högsta koncentrationen av Nb-rika faser. Den ökande mängden segring i höjdled anses bero på variationen i kylningshastighet. En hög lasereffekt (1600 W - 2000 W) och en hög framföringshastighet (1100 mm/min) visade sig vara lämplig för att minimera mängden segring.

Förutom användandet av specifika processparametrar för att minska mängden segring kan man även använda olika typer av värmebehandlingar efter tillverkningsprocessen. Upplösningsbehandling (954 ° C/1 timme) följt av en två-stegs åldring (718 °C/8 timmar + 621 °C/ 8 timmar) är standard för legering 718 och applicerades därför på byggda geometrier. Efter genomförd värmebehandling minskade segringen på grund av upplösningen av de Nb-rika faserna. Vid upplösningsbehandlingen ägde även utskiljning utav Delta-fas rum, vilket huvudsakligen kunde associeras till topp- och bottenregionerna varav mittenregionen hade betydligt mer sparsamt gällande utskiljning av Delta-fas.

iv

v

Populärvetenskaplig Sammanfattning

Nyckelord: Lasermetalldeponering; Legering 718; Processparametrar; Värmebehandlingar; Nb-rika faser

Lasermetalldeponering med pulver som tillsatsmaterial (L-DED) är en additiv tillverkningsmetod där man genom lager-på-lager-principen bygger upp en specifik geometri. Metoden har visat en enorm potential under det senaste decenniet. Flyg- och rymdindustrin står som de främsta mottagarna på grund av processens förmåga att bygga komponenter utan något större behov av omfattande skärande bearbetning vilket bidrar till mindre materialspill. Användningen av superlegering 718 inom flygmotorapplikationer är mycket omfattande vilket har lett till ett stort forskningsintresse gällande utveckling av L-DED-processen för specifikt denna superlegering. AM-processer pliktas ofta av låga bygghastigheter och långa ledtider vilket direkt påverkar produktionskostnaderna. För att övervinna låga bygghastigheter så har detta forskningsarbete fokuserat på att uppnå höga deponeringshastigheter genom högre materialtillförsel.

Inverkan av processparametrar vid L-DED processen är av största vikt eftersom de påverkar mikrostrukturen för det specifika materialet och hur materialet slutligen beter sig vid användning. I det aktuella arbetet varierades processparametrar såsom lasereffekt, framföringshastighet, pulvermatningshastighet och pulverfokusförskjutning där deras inflytande på geometri och mikrostruktur av ensträngsväggar gällande legering 718 har analyserats. Ensträngsväggarnas geometri har undersökts i form av höjd, bredd, inträngningsdjup och effektivitet gällande pulvermatning vilket bestämts genom att mäta tvärsnittsarean för deponering samt uppblandning. Deponeringens mikrostruktur visade kolumnär dendritisk struktur i mitten varav likaxlig dendritisk struktur i bottenregionen samt i det övre området. Segringar i form av Nb-rika faser som Laves och NbC kunde identifieras i de interdendritiska områdena samt i korngränserna. Segringen ökade längs deponeringshöjden, där bottenregionen hade den minsta och det övre området den högsta koncentrationen av Nb-rika faser. Den ökande mängden segring i höjdled anses bero på variationen i kylningshastighet. En hög lasereffekt (1600 W - 2000 W) och en hög framföringshastighet (1100 mm/min) visade sig vara lämplig för att minimera mängden segring.

Förutom användandet av specifika processparametrar för att minska mängden segring kan man även använda olika typer av värmebehandlingar efter tillverkningsprocessen. Upplösningsbehandling (954 ° C/1 timme) följt av en två-stegs åldring (718 °C/8 timmar + 621 °C/ 8 timmar) är standard för legering 718 och applicerades därför på byggda geometrier. Efter genomförd värmebehandling minskade segringen på grund av upplösningen av de Nb-rika faserna. Vid upplösningsbehandlingen ägde även utskiljning utav Delta-fas rum, vilket huvudsakligen kunde associeras till topp- och bottenregionerna varav mittenregionen hade betydligt mer sparsamt gällande utskiljning av Delta-fas.

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vii

Abstract

Title: Laser-Directed Energy Deposition: Influence of Process Parameters and Heat-Treatments

Keywords: Directed Energy Deposition; Alloy 718; Process Parameters; Heat Treatments; Nb- rich Phases

ISBN: Printed: 978-91-88847-67-6; Electronic: 978-91-88847-66-9

Laser-Directed Energy Deposition (L-DED), an Additive Manufacturing (AM) process used for the fabrication of parts in a layer-wise approach has displayed an immense potential over the last decade. The aerospace industry stands as the primary beneficiary due to the L-DED process capability to build near-net-shape components with minimal tooling and thereby producing minimum wastage because of reduced machining. The widespread use of Alloy 718 in the aero-engine application has prompted huge research interest in the development of L-DED processing of this superalloy. AM processes are hindered by low build rates and high cycle times which directly affects the process costs. To overcome these issues, the present work focusses on obtaining high deposition rates through a high material feed. Studying the influence of process parameters during the L-DED process is of prime importance as they determine the performance of in-service structures. In the present work, process parameters such as laser power, scanning speed, feed rate and stand-off distances are varied and their influence on geometry and microstructure of Alloy 718 single-track deposits are analyzed. The geometry of deposits is measured in terms of height, width and depth; and the powder capture efficiency is determined by measuring areas of deposition and dilution. The microstructure of the deposits shows a columnar dendritic structure in the middle and bottom region of the deposits and equiaxed grains in the top region. Nb-rich segregation involving laves and NbC phases, typical of Alloy 718 is found in the interdendritic regions and grain boundaries. The segregation increases along the height of the deposit with the bottom region having the least and the top region showing the highest concentration of Nb-rich phases due to the variation in cooling rates. A high laser power (1600 W – 2000 W) and a high scanning speed (1100 mm/min) are found to be the preferable processing conditions for minimizing segregation. Another approach to minimize segregation is by performing post-build heat treatments. The solution treatment (954 °C/1 hr) and double aging (718 °C/8 hr + 621 °C/ 8 hr) standardized for the wrought form of Alloy 718 is applied to as-built deposits which showed a reduction in segregation due to the dissolution of Nb-rich phases. Upon solution treatment, this reduction is accompanied by precipitation of the delta phase, found predominantly in top and bottom regions and sparsely in the middle region of the deposit.

vi

vii

Abstract

Title: Laser-Directed Energy Deposition: Influence of Process Parameters and Heat-Treatments

Keywords: Directed Energy Deposition; Alloy 718; Process Parameters; Heat Treatments; Nb- rich Phases

ISBN: Printed: 978-91-88847-67-6; Electronic: 978-91-88847-66-9

Laser-Directed Energy Deposition (L-DED), an Additive Manufacturing (AM) process used for the fabrication of parts in a layer-wise approach has displayed an immense potential over the last decade. The aerospace industry stands as the primary beneficiary due to the L-DED process capability to build near-net-shape components with minimal tooling and thereby producing minimum wastage because of reduced machining. The widespread use of Alloy 718 in the aero-engine application has prompted huge research interest in the development of L-DED processing of this superalloy. AM processes are hindered by low build rates and high cycle times which directly affects the process costs. To overcome these issues, the present work focusses on obtaining high deposition rates through a high material feed. Studying the influence of process parameters during the L-DED process is of prime importance as they determine the performance of in-service structures. In the present work, process parameters such as laser power, scanning speed, feed rate and stand-off distances are varied and their influence on geometry and microstructure of Alloy 718 single-track deposits are analyzed. The geometry of deposits is measured in terms of height, width and depth; and the powder capture efficiency is determined by measuring areas of deposition and dilution. The microstructure of the deposits shows a columnar dendritic structure in the middle and bottom region of the deposits and equiaxed grains in the top region. Nb-rich segregation involving laves and NbC phases, typical of Alloy 718 is found in the interdendritic regions and grain boundaries. The segregation increases along the height of the deposit with the bottom region having the least and the top region showing the highest concentration of Nb-rich phases due to the variation in cooling rates. A high laser power (1600 W – 2000 W) and a high scanning speed (1100 mm/min) are found to be the preferable processing conditions for minimizing segregation. Another approach to minimize segregation is by performing post-build heat treatments. The solution treatment (954 °C/1 hr) and double aging (718 °C/8 hr + 621 °C/ 8 hr) standardized for the wrought form of Alloy 718 is applied to as-built deposits which showed a reduction in segregation due to the dissolution of Nb-rich phases. Upon solution treatment, this reduction is accompanied by precipitation of the delta phase, found predominantly in top and bottom regions and sparsely in the middle region of the deposit.

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ix

Appended Publications

Paper A. Influence of Laser-Directed Energy Deposition Process Parameters and Thermal Post-Treatments on Nb-rich Secondary Phases of Alloy 718 Specimens, Submitted to Metals (Under Review)

Suhas Sreekanth, Kjell Hurtig, Shrikant Joshi, Joel Andersson The lead author was responsible for preliminary planning, characterization, analysis of results and preliminary writing. Co-authors contributed towards planning, experiments, interpretation of results and reviewing the manuscript.

Paper B. Effect of Direct Energy Deposition Process Parameters on Single-Track Deposits of Alloy 718, Metals 2020, 10, 96; DOI: https://doi.org/10.3390/met10010096

Suhas Sreekanth, Ehsan Ghassemali, Kjell Hurtig, Shrikant Joshi, Joel Andersson

The lead author was responsible for preliminary planning, characterization, analysis of results and preliminary writing. Co-authors contributed towards conducting experiments, interpretation of results and reviewing the manuscript.

viii

ix

Appended Publications

Paper A. Influence of Laser-Directed Energy Deposition Process Parameters and Thermal Post-Treatments on Nb-rich Secondary Phases of Alloy 718 Specimens, Submitted to Metals (Under Review)

Suhas Sreekanth, Kjell Hurtig, Shrikant Joshi, Joel Andersson The lead author was responsible for preliminary planning, characterization, analysis of results and preliminary writing. Co-authors contributed towards planning, experiments, interpretation of results and reviewing the manuscript.

Paper B. Effect of Direct Energy Deposition Process Parameters on Single-Track Deposits of Alloy 718, Metals 2020, 10, 96; DOI: https://doi.org/10.3390/met10010096

Suhas Sreekanth, Ehsan Ghassemali, Kjell Hurtig, Shrikant Joshi, Joel Andersson

The lead author was responsible for preliminary planning, characterization, analysis of results and preliminary writing. Co-authors contributed towards conducting experiments, interpretation of results and reviewing the manuscript.

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Abbreviations

AB As-Built

AM Additive Manufacturing

BSE Back Scattered Electron

BTF Buy-to-Fly

CET Columnar-Equiaxed transition

DA Direct Aging

EBSD Electron Back Scatter Diffraction

EDS Electron Dispersive Spectroscopy

FCC Face Centred Cubic

GA Gas Atomized

GB Grain Boundary

HIP Hot Isostatic Pressing

HT Heat Treated

L-DED Laser-Directed Energy Deposition

PA Plasma Atomized

PDAS Primary Dendritic Arm Spacing

PREP Plasma Rotary Electrode Process

PSD powder size distribution

SEM Scanning Electron Microscopy

ST Solution Treatment

STA Solution Treatment + Aging

TTT Time Temperature Transformation

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xi

Abbreviations

AB As-Built

AM Additive Manufacturing

BSE Back Scattered Electron

BTF Buy-to-Fly

CET Columnar-Equiaxed transition

DA Direct Aging

EBSD Electron Back Scatter Diffraction

EDS Electron Dispersive Spectroscopy

FCC Face Centred Cubic

GA Gas Atomized

GB Grain Boundary

HIP Hot Isostatic Pressing

HT Heat Treated

L-DED Laser-Directed Energy Deposition

PA Plasma Atomized

PDAS Primary Dendritic Arm Spacing

PREP Plasma Rotary Electrode Process

PSD powder size distribution

SEM Scanning Electron Microscopy

ST Solution Treatment

STA Solution Treatment + Aging

TTT Time Temperature Transformation

xii

Symbols and Notations

γʺ Gamma Double prime phase

γʹ Gamma Prime phase

γ Gamma Matrix

δ Delta Phase

G Thermal Gradient

R Solidification Velocity

H Height of Deposit

W Width of Deposit

D Depth of Deposit

Ad Area of Deposition

As Area of Dilution

Lm Line Mass

M Feed Rate

V Deposition Speed

HI Heat Input

P Laser Power

SE Specific Energy

Ds Laser Spot Diameter

η Powder Capture Efficiency

ρ Density

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Table of Contents

Acknowledgments........................................................................... iii Populärvetenskaplig Sammanfattning .............................................. v

Abstract ......................................................................................... vii Appended Publications ................................................................... ix

Abbreviations .................................................................................. xi Table of Contents ......................................................................... xiii

1 Introduction .................................................................... 1

1.1 Research Gap and Objectives ............................................. 2

1.2 Research Questions ............................................................ 4

2 Alloy 718 ......................................................................... 7

2.1 Chemistry and Role of Alloying Elements............................. 7

2.2 Heat Treatments ................................................................ 10

3 Laser-Directed Energy Deposition .............................. 13

3.1 Lasers Properties ............................................................... 13

3.2 Powder Feedstock ............................................................. 14

3.3 Effect of Process Parameters ............................................. 16

3.4 Solidification and Microstructure ........................................ 17

3.5 Mechanical Properties........................................................ 17

4 Experiments .................................................................. 21

4.1 Laser-Directed Energy Deposition Equipment .................... 21

4.2 Material .............................................................................. 22

4.3 Design of Experiments ....................................................... 22

4.4 Material Preparation and Metallography ............................. 24

4.5 Single-Track Characterization ............................................ 26

5 Summary of Results ..................................................... 29

5.1 Geometry of Single-Track Deposits.................................... 29

5.2 Microstructure and Phase Constituents .............................. 30

xii

Symbols and Notations

γʺ Gamma Double prime phase

γʹ Gamma Prime phase

γ Gamma Matrix

δ Delta Phase

G Thermal Gradient

R Solidification Velocity

H Height of Deposit

W Width of Deposit

D Depth of Deposit

Ad Area of Deposition

As Area of Dilution

Lm Line Mass

M Feed Rate

V Deposition Speed

HI Heat Input

P Laser Power

SE Specific Energy

Ds Laser Spot Diameter

η Powder Capture Efficiency

ρ Density

xiii

Table of Contents

Acknowledgments........................................................................... iii Populärvetenskaplig Sammanfattning .............................................. v

Abstract ......................................................................................... vii Appended Publications ................................................................... ix

Abbreviations .................................................................................. xi Table of Contents ......................................................................... xiii

1 Introduction .................................................................... 1

1.1 Research Gap and Objectives ............................................. 2

1.2 Research Questions ............................................................ 4

2 Alloy 718 ......................................................................... 7

2.1 Chemistry and Role of Alloying Elements............................. 7

2.2 Heat Treatments ................................................................ 10

3 Laser-Directed Energy Deposition .............................. 13

3.1 Lasers Properties ............................................................... 13

3.2 Powder Feedstock ............................................................. 14

3.3 Effect of Process Parameters ............................................. 16

3.4 Solidification and Microstructure ........................................ 17

3.5 Mechanical Properties........................................................ 17

4 Experiments .................................................................. 21

4.1 Laser-Directed Energy Deposition Equipment .................... 21

4.2 Material .............................................................................. 22

4.3 Design of Experiments ....................................................... 22

4.4 Material Preparation and Metallography ............................. 24

4.5 Single-Track Characterization ............................................ 26

5 Summary of Results ..................................................... 29

5.1 Geometry of Single-Track Deposits.................................... 29

5.2 Microstructure and Phase Constituents .............................. 30

xiv

5.3 Hardness Testing ............................................................... 34

6 Conclusions and Future Work .................................... 37

6.1 Conclusions ....................................................................... 37

6.2 Future Work ....................................................................... 38

References .............................................................................. 39

Appended Publications

Paper A Influence of Laser-Directed Energy Deposition Process Parameters and Thermal Post-Treatments on Nb-rich Secondary Phases of Alloy 718 Specimens

Paper B Effect of Direct Energy Deposition Process Parameters on Single-Track Deposits of Alloy 718

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xiv

5.3 Hardness Testing ............................................................... 34

6 Conclusions and Future Work .................................... 37

6.1 Conclusions ....................................................................... 37

6.2 Future Work ....................................................................... 38

References .............................................................................. 39

Appended Publications

Paper A Influence of Laser-Directed Energy Deposition Process Parameters and Thermal Post-Treatments on Nb-rich Secondary Phases of Alloy 718 Specimens

Paper B Effect of Direct Energy Deposition Process Parameters on Single-Track Deposits of Alloy 718

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1

1 Introduction

The modern-day turbine works on the principle of conversion of thermal energy into mechanical work. Operation capability at higher temperatures translates into higher work output. This is the very basis of jet-engines and gas-turbines, in which the combustion products of a compressed fuel-air mixture are allowed to expand inside a turbine and consequently the energy is converted into shaft power [1]. The continual improvements in alloy development, coating material, fabricability of intricate cooling channels and close composition control by the processing of material have increased the high-temperature capability of jet-engines by two-fold in the past 70 years [2]. In hot-sections of jet-engine, the gas temperatures vary between 1500 °C in the combustion chamber to about 600 °C in the exhaust casing [3]. The portfolio of alloy systems that can offer mechanical stability and inertness in such harsh working environments is quite limited. Nickel-based superalloys are one such group of materials used for high-temperature applications and Alloy 718 is one of the predominantly used superalloys especially in aerospace and oil and gas applications.

The operation temperature of the Alloy 718 is in the vicinity of 650 °C. It has quite a complex alloy chemistry that involves about 10 major alloying elements. The alloy is routinely used in fabrication of stator components of a turbine such as fasteners, diffusers and exhaust casings, intermediate and low-pressure turbine vanes [4,5]. Alloy 718 has relatively good weldability and fabricability compared to other classes of Ni-based superalloys. These properties coupled with the relatively low cost, make Alloy 718 a viable choice for fabrication of aero-engine components.

A relatively new process of fabricating components called Additive Manufacturing (AM) has attracted major interest across the manufacturing sector. Principally, AM refers to a group of bottom-up processes wherein the components are manufactured in a layer-wise sequence without the aid of dies and fixtures [6]. The actual definition refers to a CAD model utilized for the building process. Some derivatives of traditional welding processes use a CNC program or similar tools for defining the build path of 3D deposition. The idea of manufacturing with minimal tooling is an attractive feature for aerospace industries as it saves both time and cost of production. One such AM process to draw major attention is the processing of powder form of feedstock by a laser energy source called Laser Directed Energy Deposition (L-DED), which is the focus of this present research.

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1

1 Introduction

The modern-day turbine works on the principle of conversion of thermal energy into mechanical work. Operation capability at higher temperatures translates into higher work output. This is the very basis of jet-engines and gas-turbines, in which the combustion products of a compressed fuel-air mixture are allowed to expand inside a turbine and consequently the energy is converted into shaft power [1]. The continual improvements in alloy development, coating material, fabricability of intricate cooling channels and close composition control by the processing of material have increased the high-temperature capability of jet-engines by two-fold in the past 70 years [2]. In hot-sections of jet-engine, the gas temperatures vary between 1500 °C in the combustion chamber to about 600 °C in the exhaust casing [3]. The portfolio of alloy systems that can offer mechanical stability and inertness in such harsh working environments is quite limited. Nickel-based superalloys are one such group of materials used for high-temperature applications and Alloy 718 is one of the predominantly used superalloys especially in aerospace and oil and gas applications.

The operation temperature of the Alloy 718 is in the vicinity of 650 °C. It has quite a complex alloy chemistry that involves about 10 major alloying elements. The alloy is routinely used in fabrication of stator components of a turbine such as fasteners, diffusers and exhaust casings, intermediate and low-pressure turbine vanes [4,5]. Alloy 718 has relatively good weldability and fabricability compared to other classes of Ni-based superalloys. These properties coupled with the relatively low cost, make Alloy 718 a viable choice for fabrication of aero-engine components.

A relatively new process of fabricating components called Additive Manufacturing (AM) has attracted major interest across the manufacturing sector. Principally, AM refers to a group of bottom-up processes wherein the components are manufactured in a layer-wise sequence without the aid of dies and fixtures [6]. The actual definition refers to a CAD model utilized for the building process. Some derivatives of traditional welding processes use a CNC program or similar tools for defining the build path of 3D deposition. The idea of manufacturing with minimal tooling is an attractive feature for aerospace industries as it saves both time and cost of production. One such AM process to draw major attention is the processing of powder form of feedstock by a laser energy source called Laser Directed Energy Deposition (L-DED), which is the focus of this present research.

2

L-DED technology is used in processing aerospace alloys for almost two decades. The earliest applications conformed to cladding, coating and surface alloying applications. In the 1990s, efficient laser technologies with high power output enabled laser processing of metallic materials through laser sintering and net-shaping technologies [7]. In the present-day scenario, L-DED is utilized for building near-net-shape components and part features, thereby imparting a buy-to-fly (BTF) ratio of 2:1. The BTF of components processed by conventional operations ranges between 6:1 to 30:1 as a consequence of extensive machining operations [8], which results in direct economic and environmental implications.

1.1 Research Gap and Objectives

In general, process-related challenges stem from the fact that the scope of AM is multidisciplinary. It broadly covers many fields including process automation, materials and metallurgy, supply-chain management, design and production. Knowledge and expertise related to multiple fields are necessary to tackle issues related to the AM processes. The build rates and costs are other hindrances, specifically in a mass-production setting which necessitates volume manufacturing.

The L-DED research on Alloy 718 mainly concerns with the influence of process parameters on geometry of deposit [9-12]. The effect of three parameters such as laser power, deposition speed and feed rate is the preliminary focus of such research. Various researchers concluded that the height of the deposits increases with feed rates but decreases with speed [10,13,14], width mainly depends on the power and speed and dilution or depth is a function of power. The effect of processing conditions on the solidification dynamics and microstructures is briefly mentioned in early L-DED works that require further attention [14]. The laser beam delivery in each processing unit is unique which leads to empirical studies that are process and machine specific [15], necessitating a need for process parameter evaluation done in the present study.

Most of the work performed involved determination of parameter window suitable for producing a certain geometry of deposit. Steen et al. [11] developed a combined parameter approach to accommodate parameter sets such as specific energy and energy density with aspect ratio and feed rates considered for Nimonic 75. The aspect ratio (width/height) of the deposit increased with an increase in power and feed rate. Low material feed presented higher dilution and high feed rates presented porosity issues marking the non-feasible regions for good deposition characteristics. Zhong et al. [16] followed a similar procedure in identifying the regions of vapourization and porosity for Alloy 718 using specific energy density. More focus on determining process-microstructural maps needs

INTRODUCTION

3

to be done. The present research focusses on finding relationships between process parameters and resulting microstructures.

Typical L-DED experiments consider feed rates less than 0.5 kg/hr which is termed as low deposition rates. Most of the earlier cited works consider such low feed rates resulting in longer deposition times to complete a build. Increasing the feed rate is advantageous in situations involving mass production, thereby resulting in higher productivity and cost savings. Apart from these benefits, a higher feed material results in better capture of laser energy [17]. Zhong et al. [16,18,19] performed a series of experiments in high deposition rate conditions (~2.1 kg/hr) which involve geometric analysis, defect identification and tensile properties. The geometric relations are similar to the aforementioned low deposition rate experiments, however, no relationships either empirical or theoretical were established between the process parameters and the bead geometry. The deposits showed nominal levels of porosities which decreased with increasing power conditions and tensile strength was comparable to a wrought form of Alloy 718. Basic microstructural characterization involving the identification of phases and solidification behavior is performed in the above research, whereas extensive research needs to be performed in order to understand the relationship between process parameters, microstructure and tensile properties of deposits.

Figure 1 shows the processing window of Alloy 718 where a linear heat input of 300 J/mm is utilized for accommodating a feed rate of 18 g/min (1.08 kg/hr) for deposition [20]. The dotted rectangle shows the parameter window considered in the present research work, where the linear heat input is determined based on a screening set of experiments for the deposition of material pertaining to a feed rate of 1.2 kg/hr. The uniqueness comes from the linear energy considered in the processing which is relatively lower compared to the other experiments. Also, the processing equipment capabilities that arise due to laser optics and powder feeder limits the feed rate to 1.2 kg/hr (max.) considered in this research.

The properties of as-deposited L-DED specimens can be enhanced by post-build heat treatments. The segregation of Nb often reported in as-built (AB) L-DED builds [21,22] is reduced by heat treatments. Many studies report the influence of varied solution treatment (ST) temperatures and heat treatment schemes on microstructure, hardness and tensile properties of deposits. Reduction in the laves phase is reported for a ST temperature of 980 °C [21]. Double solutionizing treatment with the first step performed at 1100 °C and variable second step done at temperatures between 960 °C – 1020 °C showed recrystallization of columnar grains found in AB deposits to equiaxed grains [23]. For aerospace applications, temperatures between 954 °C – 982 °C is used for solutionizing in order to

2

L-DED technology is used in processing aerospace alloys for almost two decades. The earliest applications conformed to cladding, coating and surface alloying applications. In the 1990s, efficient laser technologies with high power output enabled laser processing of metallic materials through laser sintering and net-shaping technologies [7]. In the present-day scenario, L-DED is utilized for building near-net-shape components and part features, thereby imparting a buy-to-fly (BTF) ratio of 2:1. The BTF of components processed by conventional operations ranges between 6:1 to 30:1 as a consequence of extensive machining operations [8], which results in direct economic and environmental implications.

1.1 Research Gap and Objectives

In general, process-related challenges stem from the fact that the scope of AM is multidisciplinary. It broadly covers many fields including process automation, materials and metallurgy, supply-chain management, design and production. Knowledge and expertise related to multiple fields are necessary to tackle issues related to the AM processes. The build rates and costs are other hindrances, specifically in a mass-production setting which necessitates volume manufacturing.

The L-DED research on Alloy 718 mainly concerns with the influence of process parameters on geometry of deposit [9-12]. The effect of three parameters such as laser power, deposition speed and feed rate is the preliminary focus of such research. Various researchers concluded that the height of the deposits increases with feed rates but decreases with speed [10,13,14], width mainly depends on the power and speed and dilution or depth is a function of power. The effect of processing conditions on the solidification dynamics and microstructures is briefly mentioned in early L-DED works that require further attention [14]. The laser beam delivery in each processing unit is unique which leads to empirical studies that are process and machine specific [15], necessitating a need for process parameter evaluation done in the present study.

Most of the work performed involved determination of parameter window suitable for producing a certain geometry of deposit. Steen et al. [11] developed a combined parameter approach to accommodate parameter sets such as specific energy and energy density with aspect ratio and feed rates considered for Nimonic 75. The aspect ratio (width/height) of the deposit increased with an increase in power and feed rate. Low material feed presented higher dilution and high feed rates presented porosity issues marking the non-feasible regions for good deposition characteristics. Zhong et al. [16] followed a similar procedure in identifying the regions of vapourization and porosity for Alloy 718 using specific energy density. More focus on determining process-microstructural maps needs

INTRODUCTION

3

to be done. The present research focusses on finding relationships between process parameters and resulting microstructures.

Typical L-DED experiments consider feed rates less than 0.5 kg/hr which is termed as low deposition rates. Most of the earlier cited works consider such low feed rates resulting in longer deposition times to complete a build. Increasing the feed rate is advantageous in situations involving mass production, thereby resulting in higher productivity and cost savings. Apart from these benefits, a higher feed material results in better capture of laser energy [17]. Zhong et al. [16,18,19] performed a series of experiments in high deposition rate conditions (~2.1 kg/hr) which involve geometric analysis, defect identification and tensile properties. The geometric relations are similar to the aforementioned low deposition rate experiments, however, no relationships either empirical or theoretical were established between the process parameters and the bead geometry. The deposits showed nominal levels of porosities which decreased with increasing power conditions and tensile strength was comparable to a wrought form of Alloy 718. Basic microstructural characterization involving the identification of phases and solidification behavior is performed in the above research, whereas extensive research needs to be performed in order to understand the relationship between process parameters, microstructure and tensile properties of deposits.

Figure 1 shows the processing window of Alloy 718 where a linear heat input of 300 J/mm is utilized for accommodating a feed rate of 18 g/min (1.08 kg/hr) for deposition [20]. The dotted rectangle shows the parameter window considered in the present research work, where the linear heat input is determined based on a screening set of experiments for the deposition of material pertaining to a feed rate of 1.2 kg/hr. The uniqueness comes from the linear energy considered in the processing which is relatively lower compared to the other experiments. Also, the processing equipment capabilities that arise due to laser optics and powder feeder limits the feed rate to 1.2 kg/hr (max.) considered in this research.

The properties of as-deposited L-DED specimens can be enhanced by post-build heat treatments. The segregation of Nb often reported in as-built (AB) L-DED builds [21,22] is reduced by heat treatments. Many studies report the influence of varied solution treatment (ST) temperatures and heat treatment schemes on microstructure, hardness and tensile properties of deposits. Reduction in the laves phase is reported for a ST temperature of 980 °C [21]. Double solutionizing treatment with the first step performed at 1100 °C and variable second step done at temperatures between 960 °C – 1020 °C showed recrystallization of columnar grains found in AB deposits to equiaxed grains [23]. For aerospace applications, temperatures between 954 °C – 982 °C is used for solutionizing in order to

4

minimize grain growth [24,25]. Therefore, the lowest temperature recommendation of ST is investigated in the present work.

To summarize, the research gap that exists in processing high feed rates of Alloy 718 and process parameter evaluation with respect to the deposit geometry and microstructure forms the preliminary objective of this study. Further, the effects of heat treatments on deposits are performed to study the phase transformation that happens in L-DED specimens as the precipitates are morphologically different in L-DED processed specimens compared to conventionally processed alloys [26,27]. The study aims at investigating the suitability of standard heat treatments recommended for wrought alloys on L-DED deposits.

Figure 1: Optimal processing window of L-DED for various material systems; the dotted area indicates the processing window of Alloy 718 considered in the current work (Reproduced from [20])

1.2 Research Questions

The following questions framed as per the objectives of this research, denoted by RQ1 and RQ2:

INTRODUCTION

5

RQ1. What is the effect of laser power, scan speed, feed rate, laser stand-off distance and laser spot-diameter (primary process parameters) on geometry, microstructure and hardness of deposits? RQ2. What is the influence of solution and aging treatments on microstructural characteristics such as segregation, grain size and texture of L-DED deposits?

The above research questions are answered in the subsequent results and discussion section. Furthermore, RQ1 is answered in both Paper A and Paper B and the RQ2 is reported in Paper A.

4

minimize grain growth [24,25]. Therefore, the lowest temperature recommendation of ST is investigated in the present work.

To summarize, the research gap that exists in processing high feed rates of Alloy 718 and process parameter evaluation with respect to the deposit geometry and microstructure forms the preliminary objective of this study. Further, the effects of heat treatments on deposits are performed to study the phase transformation that happens in L-DED specimens as the precipitates are morphologically different in L-DED processed specimens compared to conventionally processed alloys [26,27]. The study aims at investigating the suitability of standard heat treatments recommended for wrought alloys on L-DED deposits.

Figure 1: Optimal processing window of L-DED for various material systems; the dotted area indicates the processing window of Alloy 718 considered in the current work (Reproduced from [20])

1.2 Research Questions

The following questions framed as per the objectives of this research, denoted by RQ1 and RQ2:

INTRODUCTION

5

RQ1. What is the effect of laser power, scan speed, feed rate, laser stand-off distance and laser spot-diameter (primary process parameters) on geometry, microstructure and hardness of deposits? RQ2. What is the influence of solution and aging treatments on microstructural characteristics such as segregation, grain size and texture of L-DED deposits?

The above research questions are answered in the subsequent results and discussion section. Furthermore, RQ1 is answered in both Paper A and Paper B and the RQ2 is reported in Paper A.

6

7

2 Alloy 718

Superalloy refers to a group of materials that possess the capability to perform in extreme conditions of temperature, pressure, and chemical activity. Superalloys can withstand high homologous temperature (operation temperature/melting temperature) and sustain high tensile, creep and fatigue loads, commonly encountered in an aero-engine [28]. In the 1950s, superalloy research gathered pace to fulfill the need for better aerospace materials that could withstand higher temperatures [2]. The advent of industrial era catalyzed superalloys research by supplementing innovations like Vacuum Induction Melting for alloy processing and development of fabrication technologies like investment casting that could implement internal cooling channels in a turbine blade [3]. New classes of superalloys such as directionally solidified alloys and single-crystal systems were developed during the 1970s, which further extended the operating temperatures of aero-engines.

During its formative years, superalloys underwent many changes with respect to the alloy chemistry [29]. Earliest alloys used Fe as the base material, which was replaced by Ni having increased additions of refractory elements and Al+Ti to improve high-temperature stability and precipitation strengthening. The fast precipitation kinetics in these alloys led to difficulties involving fabricability, welding and applications in certain areas of thermal cyclic loading in aero-engine. In the late 1950s, Eiselstein [30] substituted Al+Ti with Nb which led to the invention of Alloy 718. The patent for Alloy 718 was awarded in 1962 and it accounted for about 50% of total superalloy production by 2006 [4].

2.1 Chemistry and Role of Alloying Elements

The Alloy 718 consists of nearly 10 major elements and about 8 trace and tramp elements affecting the alloy properties. It consists of a Ni-matrix with high additions of Fe that partitions into the matrix. Nb is the most vital element in the formation of gamma double prime (γʺ) precipitates, delta phase, carbide and laves phases further elaborated in Section 2.1.1. Minor additions of Al+Ti results in low amounts of gamma prime (γʹ) precipitation. The role of major alloying elements is as shown in Table 1.

6

7

2 Alloy 718

Superalloy refers to a group of materials that possess the capability to perform in extreme conditions of temperature, pressure, and chemical activity. Superalloys can withstand high homologous temperature (operation temperature/melting temperature) and sustain high tensile, creep and fatigue loads, commonly encountered in an aero-engine [28]. In the 1950s, superalloy research gathered pace to fulfill the need for better aerospace materials that could withstand higher temperatures [2]. The advent of industrial era catalyzed superalloys research by supplementing innovations like Vacuum Induction Melting for alloy processing and development of fabrication technologies like investment casting that could implement internal cooling channels in a turbine blade [3]. New classes of superalloys such as directionally solidified alloys and single-crystal systems were developed during the 1970s, which further extended the operating temperatures of aero-engines.

During its formative years, superalloys underwent many changes with respect to the alloy chemistry [29]. Earliest alloys used Fe as the base material, which was replaced by Ni having increased additions of refractory elements and Al+Ti to improve high-temperature stability and precipitation strengthening. The fast precipitation kinetics in these alloys led to difficulties involving fabricability, welding and applications in certain areas of thermal cyclic loading in aero-engine. In the late 1950s, Eiselstein [30] substituted Al+Ti with Nb which led to the invention of Alloy 718. The patent for Alloy 718 was awarded in 1962 and it accounted for about 50% of total superalloy production by 2006 [4].

2.1 Chemistry and Role of Alloying Elements

The Alloy 718 consists of nearly 10 major elements and about 8 trace and tramp elements affecting the alloy properties. It consists of a Ni-matrix with high additions of Fe that partitions into the matrix. Nb is the most vital element in the formation of gamma double prime (γʺ) precipitates, delta phase, carbide and laves phases further elaborated in Section 2.1.1. Minor additions of Al+Ti results in low amounts of gamma prime (γʹ) precipitation. The role of major alloying elements is as shown in Table 1.

8

Table 1: Role of critical alloying elements in Alloy 718 [31]

Alloying elements

Amount (wt.%)

Effect on alloy properties

Nickel 50-55 Matrix stabilizer, forms strengthening precipitates Aluminium Titanium

0.2-0.8 0.65-1.15

Provides γʹ precipitate strengthening while Al aids oxidation resistance and Ti form carbides

Chromium 17-21 Provides solution strengthening, form carbides, aids in hot corrosion and oxidation protection

Iron 15-25 Partitions into the matrix, stabilizes laves phase Niobium Tantalum

4.75-5.5 0.05 max.

Provides both solution and γʺ precipitate strengthening and form carbides.

Molybdenum Tungsten

2.8-3.3 Provides solution strengthening, oxidation resistance and form carbides

Carbon 0.08 max. Provides grain boundary (GB) stability and participates in carbide formation

2.1.1 Phase Constituents of Alloy 718

Alloy 718 is a Ni-Fe based superalloy that consists of about 50-55% Ni and about 15-25% Fe additions with the capability of operation up to 650 °C. It is relatively inexpensive compared to other Ni-based superalloys as it contains high amounts of Fe and reduced amounts of Co. The phase constituents of Alloy 718 are discussed here below.

Austenitic Matrix (γ) predominantly made of Ni and has the Face Centered Cubic (FCC) structure. Ni provides high solubility for many elements comparable to Fe and provides higher coherency strains compared to Fe due to limited compressibility of Ni, which is the basis of Solid solution strengthening [2]. Mo+W, Nb, Ta and Al provide solid solution strengthening due to the relative difference in atomic size compared to Ni.

Gamma double prime (γʺ) is the primary strengthening precipitate in Alloy 718 that has a Ni3Nb stoichiometry and a Body-Centered Tetragonal crystal structure. γʺ precipitate has a disc-shaped morphology with a lattice parameter c:a of approximately 2:1. The coherency strain between the matrix γ-γʺ is around 2.86% which produces hardening [32]. Since the coherency strains of γʺ precipitates are higher compared to γʹ, they coarsen at much lower temperatures of around 650 °C limiting the service temperature of the alloy.

Gamma Prime (γʹ) is an intermetallic compound of A3B type with Ni3(Al,Ti) stoichiometry, where electronegative elements (Co, Fe) in the matrix substitute

9

for A and electropositive elements (Al, Ti, Ta, Nb) substitute for B [2]. It has the FCC structure and is nearly coherent with the matrix. The γ-γʹ mismatch is generally around 0.1% and imparts better temperature stability of up to 800 °C [33].

Delta (δ) phase is an orthorhombic precipitate which forms due to the dissolution of metastable γʺ [34], laves and NbCs [35]. The temperature of formation of δ-phase is between 650 °C to 980 °C depending on the Nb-content in the alloy. Between 700 °C - 885 °C, γʺ coarsens rapidly and above 885 °C dissolves into the matrix due to instability [2,36]. In the temperature range of 900 °C - 950 °C, δ precipitation is accelerated [36]. Minor amounts of δ phase are precipitated which aids in hindering grain growth. δ precipitates are found to be influential in increasing strength and rupture ductility by inhibiting long-range GB sliding [37].

Carbides are found to precipitate at GBs which are the sites of concentrated vacancies. They are detrimental to ductility but increases creep resistance and rupture strengths. The most commonly found carbide in Alloy 718 is MCs and M23C6s [38] having a cubic crystal structure [39], although M6Cs are reported to be present based on alloy modifications [40]. The MCs precipitate at high temperatures in the range of 1080 °C – 1250 °C [41] and 650 °C in lower volumes as reported by Sundararaman et al. [38]. On the increase in temperatures to 750 °C, MC and M23C6 co-precipitation is reported in the above research.

Laves phase appears as islands of globular morphology having a hexagonal structure containing up to 25% wt. of Nb [41] and forms preferentially at inter-dendritic areas and GBs as a consequence of segregation. Increased amounts of Nb, Ta, Mo and Si content in the alloy facilitates the formation of laves phase [27]. The dissolution of laves upon heat treatments can benefit the ductility properties of the alloy and increases the strength due to the increased amounts of Nb in the matrix. Laves phase offers a site for crack initiation and propagation resulting in lower strength, ductility and brittleness.

The L-DED deposits show a varied amount of Nb-rich phase (laves + MC) precipitation based on the process parameter conditions [14,42,43]. Also, a compositional gradient exists along with the height of the deposits as depicted by Tabernero et al. [44]. The successive thermal cycling can aid in precipitation of minor amounts of strengthening precipitates in the bottom and middle layers of the deposits [45].

The microstructural characteristics of high deposition rate L-DED builds are studied by Li et al. [45]. The laves phase having an Nb concentration of ~25 wt%

8

Table 1: Role of critical alloying elements in Alloy 718 [31]

Alloying elements

Amount (wt.%)

Effect on alloy properties

Nickel 50-55 Matrix stabilizer, forms strengthening precipitates Aluminium Titanium

0.2-0.8 0.65-1.15

Provides γʹ precipitate strengthening while Al aids oxidation resistance and Ti form carbides

Chromium 17-21 Provides solution strengthening, form carbides, aids in hot corrosion and oxidation protection

Iron 15-25 Partitions into the matrix, stabilizes laves phase Niobium Tantalum

4.75-5.5 0.05 max.

Provides both solution and γʺ precipitate strengthening and form carbides.

Molybdenum Tungsten

2.8-3.3 Provides solution strengthening, oxidation resistance and form carbides

Carbon 0.08 max. Provides grain boundary (GB) stability and participates in carbide formation

2.1.1 Phase Constituents of Alloy 718

Alloy 718 is a Ni-Fe based superalloy that consists of about 50-55% Ni and about 15-25% Fe additions with the capability of operation up to 650 °C. It is relatively inexpensive compared to other Ni-based superalloys as it contains high amounts of Fe and reduced amounts of Co. The phase constituents of Alloy 718 are discussed here below.

Austenitic Matrix (γ) predominantly made of Ni and has the Face Centered Cubic (FCC) structure. Ni provides high solubility for many elements comparable to Fe and provides higher coherency strains compared to Fe due to limited compressibility of Ni, which is the basis of Solid solution strengthening [2]. Mo+W, Nb, Ta and Al provide solid solution strengthening due to the relative difference in atomic size compared to Ni.

Gamma double prime (γʺ) is the primary strengthening precipitate in Alloy 718 that has a Ni3Nb stoichiometry and a Body-Centered Tetragonal crystal structure. γʺ precipitate has a disc-shaped morphology with a lattice parameter c:a of approximately 2:1. The coherency strain between the matrix γ-γʺ is around 2.86% which produces hardening [32]. Since the coherency strains of γʺ precipitates are higher compared to γʹ, they coarsen at much lower temperatures of around 650 °C limiting the service temperature of the alloy.

Gamma Prime (γʹ) is an intermetallic compound of A3B type with Ni3(Al,Ti) stoichiometry, where electronegative elements (Co, Fe) in the matrix substitute

9

for A and electropositive elements (Al, Ti, Ta, Nb) substitute for B [2]. It has the FCC structure and is nearly coherent with the matrix. The γ-γʹ mismatch is generally around 0.1% and imparts better temperature stability of up to 800 °C [33].

Delta (δ) phase is an orthorhombic precipitate which forms due to the dissolution of metastable γʺ [34], laves and NbCs [35]. The temperature of formation of δ-phase is between 650 °C to 980 °C depending on the Nb-content in the alloy. Between 700 °C - 885 °C, γʺ coarsens rapidly and above 885 °C dissolves into the matrix due to instability [2,36]. In the temperature range of 900 °C - 950 °C, δ precipitation is accelerated [36]. Minor amounts of δ phase are precipitated which aids in hindering grain growth. δ precipitates are found to be influential in increasing strength and rupture ductility by inhibiting long-range GB sliding [37].

Carbides are found to precipitate at GBs which are the sites of concentrated vacancies. They are detrimental to ductility but increases creep resistance and rupture strengths. The most commonly found carbide in Alloy 718 is MCs and M23C6s [38] having a cubic crystal structure [39], although M6Cs are reported to be present based on alloy modifications [40]. The MCs precipitate at high temperatures in the range of 1080 °C – 1250 °C [41] and 650 °C in lower volumes as reported by Sundararaman et al. [38]. On the increase in temperatures to 750 °C, MC and M23C6 co-precipitation is reported in the above research.

Laves phase appears as islands of globular morphology having a hexagonal structure containing up to 25% wt. of Nb [41] and forms preferentially at inter-dendritic areas and GBs as a consequence of segregation. Increased amounts of Nb, Ta, Mo and Si content in the alloy facilitates the formation of laves phase [27]. The dissolution of laves upon heat treatments can benefit the ductility properties of the alloy and increases the strength due to the increased amounts of Nb in the matrix. Laves phase offers a site for crack initiation and propagation resulting in lower strength, ductility and brittleness.

The L-DED deposits show a varied amount of Nb-rich phase (laves + MC) precipitation based on the process parameter conditions [14,42,43]. Also, a compositional gradient exists along with the height of the deposits as depicted by Tabernero et al. [44]. The successive thermal cycling can aid in precipitation of minor amounts of strengthening precipitates in the bottom and middle layers of the deposits [45].

The microstructural characteristics of high deposition rate L-DED builds are studied by Li et al. [45]. The laves phase having an Nb concentration of ~25 wt%

10

in the interdendritic regions of the top, middle and bottom regions of the build is envisaged. γʺ and γʹ precipitates are found to precipitate inhomogeneously close to the laves phase. Inter- and intra-granular δ-precipitates are reported in this study. Precipitate Free Zones are reported in the vicinity of GBs due to the presence of Nb-rich phases, which is proved to facilitate the propagation of cracks and thereby affecting the mechanical properties of deposits [46]. The influence of heat treatments on the phase fraction of precipitates is discussed below.

2.2 Heat Treatments

The diffusion processes in Alloy 718 are similar to those in high alloyed ferrous materials. The composition of the alloy involves many d-block transition group elements that enable the substitutional diffusion reaction. The presence of carbides and borides facilitates interstitial diffusion reactions.

The Time-Temperature-Transformation (TTT)- diagram of wrought Alloy 718 is as shown in Figure 2 which depicts the temperature and time for the evolution of 0.1% Volume δ, γʹ, γʺ and α-Cr phases. Solidification study of Alloy 718 by Knorovsky, Cieslak, Headley, Romig and Hammetter [41] having Nb around 5.2 wt.% showed that the eutectic transformation temperatures for γ/NbC and γ/laves in the range of 1250 °C and 1195 °C, respectively.

Figure 2: TTT diagram for wrought Alloy 718 (©1991 by The Minerals, Metals and the Materials Society. Used with permission [47]).

11

The δ-solvus temperature is denoted isothermally at 1010 °C (1850 °F) for an alloy containing 5.38 %wt. Nb [47]. Since the alloy considered for TTT-diagram is in the wrought form, no laves phase transformation is considered. In the temperature range of 600 °C (1110 °F) - 900 °C (1650 °F), γʹand γʺ precipitates are formed. The time transformation of γʹand γʺ to δ precipitates is denoted by the curve γʹ+γʺ+δ. Above 950 °C (1740 °F), only δ-phase persists with the continual dissolution of γʹ+γʺ- precipitates.

Depending on the material forms and applications, varied heat treatment schemes are recommended for Alloy 718. ST and aging are the most common heat treatment procedures used in wrought form processing and cast materials are frequently subjected to homogenization treatment.

Homogenization: The cast microstructure of Alloy 718 consists of large amounts of secondary Nb-rich precipitates such as NbCs and laves, owing to the slow cooling rates. The size of the precipitates are relatively bigger and hence necessitates high-temperature ST [24]. These treatments are performed in the temperature range close to 1093 °C (2010 °F) as per the AMS 5383D standard specification. In some cases that require close control of grain size, extensive δ-phase is precipitated prior to homogenization in the temperature ranging between 875 °C (1607 °F)- 950 °C (1740 °F) called δ-dumping treatment.

Solution treatment: ST is performed to saturate Nb in the γ-matrix by creating a near homogeneous solid-solution. The slow dissolution of higher-order Nb-rich phases such as laves is often reported for solutionizing temperatures. Based on the application and refractory element content in the alloy, ST temperature varies between 950 °C (1740 °F) – 1020 °C (1870 °F). A quasi-treatment to homogenize, stress-relieve and ST specimens are often incorporated in the temperature field close to 1065 °C (1950 °F). In aerospace applications, a lower ST temperature range of 925 °C (1700 °F) – 1010 °C (1850 °F) is used to hinder grain growth that influences fatigue strength and rupture ductility properties. In oil and gas turbines, resistance to corrosion and toughness takes precedence over the strength of the alloy and hence a higher temperature in the range of 1010 °C – 1065 °C is employed for solutionizing. The times for ST can vary between 1hr. – 4 hrs., based on the treatment temperature and component size.

Aging: A double step aging treatment is the most common procedure followed in aerospace applications. The higher temperature aging step is performed at a temperature range of 718 °C (1325 °F) – 780 °C (1435 °F) to precipitate γʹ- nuclei and to assist the growth of δ-precipitates. The lower temperature aging is performed at temperatures just above the operating temperature of in-service components which varies between 600 °C (1110 °F) - 650 °C (1200 °F). γʺ-

10

in the interdendritic regions of the top, middle and bottom regions of the build is envisaged. γʺ and γʹ precipitates are found to precipitate inhomogeneously close to the laves phase. Inter- and intra-granular δ-precipitates are reported in this study. Precipitate Free Zones are reported in the vicinity of GBs due to the presence of Nb-rich phases, which is proved to facilitate the propagation of cracks and thereby affecting the mechanical properties of deposits [46]. The influence of heat treatments on the phase fraction of precipitates is discussed below.

2.2 Heat Treatments

The diffusion processes in Alloy 718 are similar to those in high alloyed ferrous materials. The composition of the alloy involves many d-block transition group elements that enable the substitutional diffusion reaction. The presence of carbides and borides facilitates interstitial diffusion reactions.

The Time-Temperature-Transformation (TTT)- diagram of wrought Alloy 718 is as shown in Figure 2 which depicts the temperature and time for the evolution of 0.1% Volume δ, γʹ, γʺ and α-Cr phases. Solidification study of Alloy 718 by Knorovsky, Cieslak, Headley, Romig and Hammetter [41] having Nb around 5.2 wt.% showed that the eutectic transformation temperatures for γ/NbC and γ/laves in the range of 1250 °C and 1195 °C, respectively.

Figure 2: TTT diagram for wrought Alloy 718 (©1991 by The Minerals, Metals and the Materials Society. Used with permission [47]).

11

The δ-solvus temperature is denoted isothermally at 1010 °C (1850 °F) for an alloy containing 5.38 %wt. Nb [47]. Since the alloy considered for TTT-diagram is in the wrought form, no laves phase transformation is considered. In the temperature range of 600 °C (1110 °F) - 900 °C (1650 °F), γʹand γʺ precipitates are formed. The time transformation of γʹand γʺ to δ precipitates is denoted by the curve γʹ+γʺ+δ. Above 950 °C (1740 °F), only δ-phase persists with the continual dissolution of γʹ+γʺ- precipitates.

Depending on the material forms and applications, varied heat treatment schemes are recommended for Alloy 718. ST and aging are the most common heat treatment procedures used in wrought form processing and cast materials are frequently subjected to homogenization treatment.

Homogenization: The cast microstructure of Alloy 718 consists of large amounts of secondary Nb-rich precipitates such as NbCs and laves, owing to the slow cooling rates. The size of the precipitates are relatively bigger and hence necessitates high-temperature ST [24]. These treatments are performed in the temperature range close to 1093 °C (2010 °F) as per the AMS 5383D standard specification. In some cases that require close control of grain size, extensive δ-phase is precipitated prior to homogenization in the temperature ranging between 875 °C (1607 °F)- 950 °C (1740 °F) called δ-dumping treatment.

Solution treatment: ST is performed to saturate Nb in the γ-matrix by creating a near homogeneous solid-solution. The slow dissolution of higher-order Nb-rich phases such as laves is often reported for solutionizing temperatures. Based on the application and refractory element content in the alloy, ST temperature varies between 950 °C (1740 °F) – 1020 °C (1870 °F). A quasi-treatment to homogenize, stress-relieve and ST specimens are often incorporated in the temperature field close to 1065 °C (1950 °F). In aerospace applications, a lower ST temperature range of 925 °C (1700 °F) – 1010 °C (1850 °F) is used to hinder grain growth that influences fatigue strength and rupture ductility properties. In oil and gas turbines, resistance to corrosion and toughness takes precedence over the strength of the alloy and hence a higher temperature in the range of 1010 °C – 1065 °C is employed for solutionizing. The times for ST can vary between 1hr. – 4 hrs., based on the treatment temperature and component size.

Aging: A double step aging treatment is the most common procedure followed in aerospace applications. The higher temperature aging step is performed at a temperature range of 718 °C (1325 °F) – 780 °C (1435 °F) to precipitate γʹ- nuclei and to assist the growth of δ-precipitates. The lower temperature aging is performed at temperatures just above the operating temperature of in-service components which varies between 600 °C (1110 °F) - 650 °C (1200 °F). γʺ-

12

precipitation and growth of γʹ-phase take effect at lower aging temperatures. In case of oil and gas applications, a single step aging treatment is commonly performed in the temperature range of 650 °C – 815 °C (1500 °F).

The homogenization of L-DED deposits is rarely performed as high-temperature treatments induce grain growth. But Liu et al. [23] reported the recrystallization of columnar dendrites into equiaxed grains upon subjecting the AB specimen to 1100 °C/ 1.5 hr treatment. The thermal stresses present in the AB deposits were reported to be the reason for recrystallization. However, the recrystallized grain sizes were too large (100 µm) compared to wrought materials. Further, both homogenization and solution treatments are reported to influence the dissolution of Nb-rich phases [43,46]. Also, ST is found to aid the precipitation of δ-phase at the expense of the dissolved Nb-rich phase. Table 2 shows the influence of post-build heat treatments on tensile properties of deposits which is further elaborated in Section 3.5.

13

3 Laser-Directed Energy Deposition

Directed Energy Deposition processes by principle use thermal energy to melt material as they are being deposited on the substrate in a layer-wise sequence as per the definition by ISO/ASTM 52900 [6]. The characteristics of the DED process depend on two main factors viz., energy source and material feedstock. A laser beam, electron beam, electric arc or joule heating are the common energy sources utilized in DED processes [48]. The wire form of feedstock can be used for all variants of DED energy sources whereas, powder feedstock is generally processed with the L-DED system. Wire strips are used for coating and cladding applications in laser variant of the DED process. Since the energy source and material forms are similar to those used in welding applications, the maturity level of DED technology is quite advanced compared to other AM processes [49]. Hybrid processes such as Controlled Metal Build-up by Fraunhofer IPT implements DED and machining to achieve high geometrical accuracy. Induction coils and plasma arc are used in tandem with laser power in the variants of L-DED process in order to increase deposition and process efficiency, and also to minimize part distortion.

L-DED, also known as Laser Metal Deposition, is a common AM process used for fabrication of 3D components, part features, repair and refurbishing of metallic components which is attractive for aerospace applications. The application portfolio also extends to prototyping, cladding and surface alloying applications. The above application features combined with the feasibility of the process in utilizing many forms of feedstock for fabrication of varied structural components make L-DED a versatile process.

Laser properties and feedstock material and their delivery are the important controllable factors that affect the deposition properties. Substrate conditions and protective atmospheres constitute other factors affecting the quality of the deposits. A brief account of these factors that influence the deposition properties is presented in this chapter.

3.1 Lasers Properties

Laser is a term used to describe a highly coherent beam of light with high intensity compared to visible light. The lasing action is a result of stimulated emission where the electron in the higher energy state makes a transition to a lower energy state under the influence of a photon. The common types of lasers used in

12

precipitation and growth of γʹ-phase take effect at lower aging temperatures. In case of oil and gas applications, a single step aging treatment is commonly performed in the temperature range of 650 °C – 815 °C (1500 °F).

The homogenization of L-DED deposits is rarely performed as high-temperature treatments induce grain growth. But Liu et al. [23] reported the recrystallization of columnar dendrites into equiaxed grains upon subjecting the AB specimen to 1100 °C/ 1.5 hr treatment. The thermal stresses present in the AB deposits were reported to be the reason for recrystallization. However, the recrystallized grain sizes were too large (100 µm) compared to wrought materials. Further, both homogenization and solution treatments are reported to influence the dissolution of Nb-rich phases [43,46]. Also, ST is found to aid the precipitation of δ-phase at the expense of the dissolved Nb-rich phase. Table 2 shows the influence of post-build heat treatments on tensile properties of deposits which is further elaborated in Section 3.5.

13

3 Laser-Directed Energy Deposition

Directed Energy Deposition processes by principle use thermal energy to melt material as they are being deposited on the substrate in a layer-wise sequence as per the definition by ISO/ASTM 52900 [6]. The characteristics of the DED process depend on two main factors viz., energy source and material feedstock. A laser beam, electron beam, electric arc or joule heating are the common energy sources utilized in DED processes [48]. The wire form of feedstock can be used for all variants of DED energy sources whereas, powder feedstock is generally processed with the L-DED system. Wire strips are used for coating and cladding applications in laser variant of the DED process. Since the energy source and material forms are similar to those used in welding applications, the maturity level of DED technology is quite advanced compared to other AM processes [49]. Hybrid processes such as Controlled Metal Build-up by Fraunhofer IPT implements DED and machining to achieve high geometrical accuracy. Induction coils and plasma arc are used in tandem with laser power in the variants of L-DED process in order to increase deposition and process efficiency, and also to minimize part distortion.

L-DED, also known as Laser Metal Deposition, is a common AM process used for fabrication of 3D components, part features, repair and refurbishing of metallic components which is attractive for aerospace applications. The application portfolio also extends to prototyping, cladding and surface alloying applications. The above application features combined with the feasibility of the process in utilizing many forms of feedstock for fabrication of varied structural components make L-DED a versatile process.

Laser properties and feedstock material and their delivery are the important controllable factors that affect the deposition properties. Substrate conditions and protective atmospheres constitute other factors affecting the quality of the deposits. A brief account of these factors that influence the deposition properties is presented in this chapter.

3.1 Lasers Properties

Laser is a term used to describe a highly coherent beam of light with high intensity compared to visible light. The lasing action is a result of stimulated emission where the electron in the higher energy state makes a transition to a lower energy state under the influence of a photon. The common types of lasers used in

14

additive applications are CO2 lasers, Yb-doped fiber lasers, Nd:YAG solid state lasers and diode lasers. In the present work, Yb-doped lasers are pumped with the help of diode lasers that produces a focused beam in the range of 1030 nm - 1080 nm [50] which is preferable for processing of metallic materials [51]. Different modes of power distribution can be obtained in laser operation by having suitable optical configurations. A continuous-wave gaussian form of power distribution is the most common mode used for L-DED applications (see Figure 1 in Paper B). Top-hat power distribution can be obtained by having a working distance equal to the focal length of the lens.

Laser pulsing mode of operation is commonly applied in order to minimize the energy density. A pulsing operation can be accomplished by modulation of power source, Q-switching or mode-locking functions [52]. Xiao et.al [53] reported the results of processing Alloy 718 using a pulsed mode and continuous mode laser, which shows that discrete laves phase form in pulsed-wave as the energy density per unit volume is comparatively less compared to the continuous mode of operation. Chen et.al [54] performed deposition by utilizing flat-top and gaussian modes which showed less segregation in the gaussian mode of power distribution. The related mechanical properties from these studies are further discussed in Section 3.5.

3.2 Powder Feedstock

Laser DED process is the preferred option for processing powder feedstock as EB and plasma-arc DED processes utilize wire form of material. Powder feedstock offers higher precision and can be used in sensitive applications as the heat input for melting the powder is relatively lower compared to wire. This minimizes part distortions in applications necessitating the addition of part features and repair for in-service components [55].

Two primary forms of powders are used in metal-AM applications, viz., Gas Atomized (GA) and Plasma Atomized (PA) powders, although GA powders are the primary source of metal-AM feedstock as they are inexpensive compared to PA powders. Zhong et al. [56] compared GA and Plasma Rotary Electrode Process (PREP) powder L-DED builds of Alloy 718 and reported higher porosity levels and discrete Nb segregation in case of GA deposits compared to PREP specimens. The Nb-rich phases were extensive in case of PREP powder deposits compared to GA builds, and this was attributed to the partitioning of laser energy in determining the area of deposition (Ad) and the area of dilution (As) of the deposit. Ad of PREP deposit was higher (2.47 mm2) compared to GA deposit (2.20 mm2), and hence it was assumed that the PREP deposits encounter lower

LASER-DIRECTED ENERGY DEPOSITION

15

cooling rates. Also, the nominal Nb content in PREP powders was slightly higher (0.35 %wt.) compared to GA powders.

Qi et al. [57] studied the deposition properties of course GA (74 µm – 125 µm), fine GA (44 µm – 74 µm) and PREP (44 µm – 149 µm) powders which showed that the deposits produced from coarse GA powder had about 2 % - 2.25 % volume fraction porosity, whereas the fine GA powder which had 1.5 %. The PREP powder deposits showed the lowest volume fraction of porosity of less than 0.5% for a heat input of ~53 J/mm. Upon increasing the heat input to 110 J/mm, the porosities related to the coarse and fine GA powders dropped below 1%, indicating dense build at high energy conditions.

The powder size distribution (PSD) is a vital factor that influences the powder focus diameter. Smaller particle diameter and a narrow PSD led to a concentrated powder focus and better capture efficiency [58]. Mean particle diameter (D50) of <20 µm and >70 µm are prone to clogging effects as per the research of Kong et al. [59]. The research shows an upper and lower limit that needs to be adhered for obtaining a sound deposition process, as the values depicted in the research is a function of annular nozzle opening and powder morphology.

3.2.1 Powder Delivery

The process of powder delivery starts at the powder hopper in the powder feeder arrangement. The mode of powder delivery can be batch-wise or continuous operation as per the process requirements. In batch-wise feeding, the amount of powder is delivered in bulk based on threshold weight or per unit time, whereas the continuous feeding arrangement ensures a seamless supply of powder into the melt-pool. A lead screw or a graded disk is used for controlling the powder supply in case of volumetric feeders and load cells are accommodated in case of gravimetric feeders [60]. Thus, the controlled amount of powder is discharged into the interaction area of laser blown through a coaxial or an off-axial (lateral) nozzle arrangement with the help of carrier gas. The coaxial nozzle is proven to have higher capture efficiency compared to off-axis nozzles that has a skewed powder-laser energy distribution [61]. A discrete coaxial 3-jet nozzle arrangement is found to have better accessibility for building features in 3D-applications but has a lower capture efficiency compared to the coaxial nozzle arrangement due to particle rebounding effect [62,63]. Also, the powder stream density in discrete nozzle varies in a different direction of deposition which yields varied deposit geometry [63].

14

additive applications are CO2 lasers, Yb-doped fiber lasers, Nd:YAG solid state lasers and diode lasers. In the present work, Yb-doped lasers are pumped with the help of diode lasers that produces a focused beam in the range of 1030 nm - 1080 nm [50] which is preferable for processing of metallic materials [51]. Different modes of power distribution can be obtained in laser operation by having suitable optical configurations. A continuous-wave gaussian form of power distribution is the most common mode used for L-DED applications (see Figure 1 in Paper B). Top-hat power distribution can be obtained by having a working distance equal to the focal length of the lens.

Laser pulsing mode of operation is commonly applied in order to minimize the energy density. A pulsing operation can be accomplished by modulation of power source, Q-switching or mode-locking functions [52]. Xiao et.al [53] reported the results of processing Alloy 718 using a pulsed mode and continuous mode laser, which shows that discrete laves phase form in pulsed-wave as the energy density per unit volume is comparatively less compared to the continuous mode of operation. Chen et.al [54] performed deposition by utilizing flat-top and gaussian modes which showed less segregation in the gaussian mode of power distribution. The related mechanical properties from these studies are further discussed in Section 3.5.

3.2 Powder Feedstock

Laser DED process is the preferred option for processing powder feedstock as EB and plasma-arc DED processes utilize wire form of material. Powder feedstock offers higher precision and can be used in sensitive applications as the heat input for melting the powder is relatively lower compared to wire. This minimizes part distortions in applications necessitating the addition of part features and repair for in-service components [55].

Two primary forms of powders are used in metal-AM applications, viz., Gas Atomized (GA) and Plasma Atomized (PA) powders, although GA powders are the primary source of metal-AM feedstock as they are inexpensive compared to PA powders. Zhong et al. [56] compared GA and Plasma Rotary Electrode Process (PREP) powder L-DED builds of Alloy 718 and reported higher porosity levels and discrete Nb segregation in case of GA deposits compared to PREP specimens. The Nb-rich phases were extensive in case of PREP powder deposits compared to GA builds, and this was attributed to the partitioning of laser energy in determining the area of deposition (Ad) and the area of dilution (As) of the deposit. Ad of PREP deposit was higher (2.47 mm2) compared to GA deposit (2.20 mm2), and hence it was assumed that the PREP deposits encounter lower

LASER-DIRECTED ENERGY DEPOSITION

15

cooling rates. Also, the nominal Nb content in PREP powders was slightly higher (0.35 %wt.) compared to GA powders.

Qi et al. [57] studied the deposition properties of course GA (74 µm – 125 µm), fine GA (44 µm – 74 µm) and PREP (44 µm – 149 µm) powders which showed that the deposits produced from coarse GA powder had about 2 % - 2.25 % volume fraction porosity, whereas the fine GA powder which had 1.5 %. The PREP powder deposits showed the lowest volume fraction of porosity of less than 0.5% for a heat input of ~53 J/mm. Upon increasing the heat input to 110 J/mm, the porosities related to the coarse and fine GA powders dropped below 1%, indicating dense build at high energy conditions.

The powder size distribution (PSD) is a vital factor that influences the powder focus diameter. Smaller particle diameter and a narrow PSD led to a concentrated powder focus and better capture efficiency [58]. Mean particle diameter (D50) of <20 µm and >70 µm are prone to clogging effects as per the research of Kong et al. [59]. The research shows an upper and lower limit that needs to be adhered for obtaining a sound deposition process, as the values depicted in the research is a function of annular nozzle opening and powder morphology.

3.2.1 Powder Delivery

The process of powder delivery starts at the powder hopper in the powder feeder arrangement. The mode of powder delivery can be batch-wise or continuous operation as per the process requirements. In batch-wise feeding, the amount of powder is delivered in bulk based on threshold weight or per unit time, whereas the continuous feeding arrangement ensures a seamless supply of powder into the melt-pool. A lead screw or a graded disk is used for controlling the powder supply in case of volumetric feeders and load cells are accommodated in case of gravimetric feeders [60]. Thus, the controlled amount of powder is discharged into the interaction area of laser blown through a coaxial or an off-axial (lateral) nozzle arrangement with the help of carrier gas. The coaxial nozzle is proven to have higher capture efficiency compared to off-axis nozzles that has a skewed powder-laser energy distribution [61]. A discrete coaxial 3-jet nozzle arrangement is found to have better accessibility for building features in 3D-applications but has a lower capture efficiency compared to the coaxial nozzle arrangement due to particle rebounding effect [62,63]. Also, the powder stream density in discrete nozzle varies in a different direction of deposition which yields varied deposit geometry [63].

16

3.3 Effect of Process Parameters

Laser power, laser spot diameter, deposition speed and feed rates form the important set of parameters that are sensitive to the quality of deposition [15]. The extent of overlap and deposition sequence are important factors in 3D builds as these control the powder efficiency, cooling rate, residual stresses and density of builds [64]. Carrier and shielding gas flow rate and powder stand-off distance further improve the quality of builds.

Porosities, lack of fusion, cracks, residual stresses and poor surface finish are some of the defects that pose challenges in L-DED research. These defects are a function of processing parameters and to minimize defects process optimization needs to be performed. Time and again it is shown that the process related porosities can be eliminated by implementing high power density [65,66]. But very high-power density tends to vaporize material, re-melt pre-deposit layer or substrate, yielding high residual stresses [67]. Segregation and precipitation of unwanted phases are inherent properties of the material and can only be eliminated by suitable process parameters or suitable post-treatments. As deposition speed increases, the melt pool dynamics and the solidification mode change [68]. Too high deposition speed can result in insufficient dwell times and hence the lack of fusion and porosity and the other extreme results in vaporization.

Segerstark et al. [13] performed L-DED using GA Alloy 718 powder on a stainless-steel substrate in order to study the microstructure of builds and dilution effects on build quality and geometry. It shows higher Nb-rich phases precipitate at the bottom of the deposits due to the dilution of Fe from the substrate. It is also reported that low energy input has a higher tendency to form equiaxed grains [69,70]. Zhang et al. [71] has performed experiments by varying one parameter at a time while the other processing conditions were kept a constant and studied the influence it has on the geometry of the build and microstructure. An empirical relation to find overlap distance on the surface finish is verified and suggested for L-DED in the same work.

Ma et al. [69] studied the effects of energy input, the extent of overlap between successive deposits and solidification conditions that yielded different grain morphologies. It is showed that a high extent of overlap of above 30% bead width with low power and speed can produce equiaxed crystals. If the alloy is not meant to be directionally solidified, it is desirable to have an equiaxed grain structure that imparts isotopic properties to materials. Kistler et.al [72] and Corbin et al. [9] introduced a normalized enthalpy method which accounts for a greater number of factors compared to energy relationships with the process to describe and

LASER-DIRECTED ENERGY DEPOSITION

17

model the geometry of builds. The effects of pre-heat were also studied on the geometry of the deposition. Güpner et al. [73] introduced a factor called process efficiency wherein the partial energies for melting the powder, melting the substrate and energy losses are compared to the power input. A lower energy condition obtained by increased speed resulted in good deposition properties as lower dilution was obtained. At high energy conditions, most of the energy was utilized for melting the substrate.

3.4 Solidification and Microstructure

Solidification and texturing of materials in the L-DED process is often controlled by variation of process parameter or feedstock modification. Thermal Gradient (G in K/mm) and solidification velocity (R in mm/s) relationship is important in determining the nature and scale of microstructural features. The product of these two terms G x R is the cooling rate (K/s), which determines the amount of segregation and size of microstructural features in the materials. The ratio G/R determines the type of microstructures. At high G and low R, corresponding to a high G/R ratio cellular or columnar microstructures prevail depending on the alloying elements present in the process [74]. In Alloy 718, due to the presence of refractory elements, dendritic solidification persists in conditions of high G [75].

Almost all researchers claim that grain growth starts epitaxially from the substrate. Low Heat Input achieved by either low power or high scanning speed is termed as the best condition to achieve equiaxed grains [76]. Segerstark et.al [77] depicted that heat input can be used for alloy modification as per the application. A high heat input showed a columnar microstructure with a directional solidification texture of grains, whereas the low heat input condition shows a bimodal equiaxed and columnar microstructure having smaller grain size. The overlap distances affect the recrystallization of grains in Alloy 718 processed by L-DED techniques [64]. An overlap extent of 20-50% of bead width was considered in the research which showed that 50% iteration has the lowest grain size after heat treatment of material at 1100 °C. Dinda et.al [78] demonstrated that the texture is a strong function of deposition pattern by incorporating unidirectional and bidirectional deposition strategies. The change in the direction of heat flow affects the grain growth and thereby breaking the epitaxial growth of grains between many layers.

3.5 Mechanical Properties

The mechanical properties of Alloy 718 depend on the alloying content, process routes and processing condition. The AB form of Alloy 718 has relatively lower strength but better ductility properties compared to the heat-treated conditions as seen in Table 2. Zhao et al. [79] showed that for low deposition rate experiments,

16

3.3 Effect of Process Parameters

Laser power, laser spot diameter, deposition speed and feed rates form the important set of parameters that are sensitive to the quality of deposition [15]. The extent of overlap and deposition sequence are important factors in 3D builds as these control the powder efficiency, cooling rate, residual stresses and density of builds [64]. Carrier and shielding gas flow rate and powder stand-off distance further improve the quality of builds.

Porosities, lack of fusion, cracks, residual stresses and poor surface finish are some of the defects that pose challenges in L-DED research. These defects are a function of processing parameters and to minimize defects process optimization needs to be performed. Time and again it is shown that the process related porosities can be eliminated by implementing high power density [65,66]. But very high-power density tends to vaporize material, re-melt pre-deposit layer or substrate, yielding high residual stresses [67]. Segregation and precipitation of unwanted phases are inherent properties of the material and can only be eliminated by suitable process parameters or suitable post-treatments. As deposition speed increases, the melt pool dynamics and the solidification mode change [68]. Too high deposition speed can result in insufficient dwell times and hence the lack of fusion and porosity and the other extreme results in vaporization.

Segerstark et al. [13] performed L-DED using GA Alloy 718 powder on a stainless-steel substrate in order to study the microstructure of builds and dilution effects on build quality and geometry. It shows higher Nb-rich phases precipitate at the bottom of the deposits due to the dilution of Fe from the substrate. It is also reported that low energy input has a higher tendency to form equiaxed grains [69,70]. Zhang et al. [71] has performed experiments by varying one parameter at a time while the other processing conditions were kept a constant and studied the influence it has on the geometry of the build and microstructure. An empirical relation to find overlap distance on the surface finish is verified and suggested for L-DED in the same work.

Ma et al. [69] studied the effects of energy input, the extent of overlap between successive deposits and solidification conditions that yielded different grain morphologies. It is showed that a high extent of overlap of above 30% bead width with low power and speed can produce equiaxed crystals. If the alloy is not meant to be directionally solidified, it is desirable to have an equiaxed grain structure that imparts isotopic properties to materials. Kistler et.al [72] and Corbin et al. [9] introduced a normalized enthalpy method which accounts for a greater number of factors compared to energy relationships with the process to describe and

LASER-DIRECTED ENERGY DEPOSITION

17

model the geometry of builds. The effects of pre-heat were also studied on the geometry of the deposition. Güpner et al. [73] introduced a factor called process efficiency wherein the partial energies for melting the powder, melting the substrate and energy losses are compared to the power input. A lower energy condition obtained by increased speed resulted in good deposition properties as lower dilution was obtained. At high energy conditions, most of the energy was utilized for melting the substrate.

3.4 Solidification and Microstructure

Solidification and texturing of materials in the L-DED process is often controlled by variation of process parameter or feedstock modification. Thermal Gradient (G in K/mm) and solidification velocity (R in mm/s) relationship is important in determining the nature and scale of microstructural features. The product of these two terms G x R is the cooling rate (K/s), which determines the amount of segregation and size of microstructural features in the materials. The ratio G/R determines the type of microstructures. At high G and low R, corresponding to a high G/R ratio cellular or columnar microstructures prevail depending on the alloying elements present in the process [74]. In Alloy 718, due to the presence of refractory elements, dendritic solidification persists in conditions of high G [75].

Almost all researchers claim that grain growth starts epitaxially from the substrate. Low Heat Input achieved by either low power or high scanning speed is termed as the best condition to achieve equiaxed grains [76]. Segerstark et.al [77] depicted that heat input can be used for alloy modification as per the application. A high heat input showed a columnar microstructure with a directional solidification texture of grains, whereas the low heat input condition shows a bimodal equiaxed and columnar microstructure having smaller grain size. The overlap distances affect the recrystallization of grains in Alloy 718 processed by L-DED techniques [64]. An overlap extent of 20-50% of bead width was considered in the research which showed that 50% iteration has the lowest grain size after heat treatment of material at 1100 °C. Dinda et.al [78] demonstrated that the texture is a strong function of deposition pattern by incorporating unidirectional and bidirectional deposition strategies. The change in the direction of heat flow affects the grain growth and thereby breaking the epitaxial growth of grains between many layers.

3.5 Mechanical Properties

The mechanical properties of Alloy 718 depend on the alloying content, process routes and processing condition. The AB form of Alloy 718 has relatively lower strength but better ductility properties compared to the heat-treated conditions as seen in Table 2. Zhao et al. [79] showed that for low deposition rate experiments,

18

heat treatment involving homogenization and standard solution treatment + aging (STA) enhanced tensile properties of deposits built using GA powder. A further enhancement of tensile properties was reported, especially ductility is improved on using PREP powder as the feedstock. The high deposition rate L-DED experiment by Zhong et al. [19] showed tensile properties similar to the GA powder deposits of Zhao et al. Further, the inclusion of 1150 °C/4 hr Hot Isostatic Pressing (HIP) treatment did not effect changes in tensile properties of specimen [80]. An increase in ST temperature to 1100 °C and consequent double aging showed a decline in tensile properties [81].

Chen et al. [54] showed that varying the laser beam shape has effects on the final properties of L-DED specimens. Flat-top and gaussian mode of power distribution was used in the experiments, wherein the gaussian form yielded a continuous band of laves phase. The flat-top power distribution yielded a higher volume fraction of discrete laves phase which displayed better mechanical properties. Xiao et al. [53] used pulsed and continuous wave (CW) mode of laser operation, that resulted in varied microstructural features that were subjected to direct aging (DA) treatments. Pulsed mode generated equiaxed grains and CW mode yielded a columnar microstructure that affected the mechanical properties. The failure mode is determined to be along the interdendritic region in case of CW specimens which indicated the presence of laves phase on contrary to the pulsed mode specimens that showed a transgranular ductile fracture. The research by Qi et al. [57] shows the anisotropy of L-DED components as samples built vertically perpendicular to the build plate showed better strengths compared to the horizontally built specimens that had better ductility properties.

The research mentioned in Section 3.3 - 3.5 provides an overview of the influence of processing conditions on geometry, microstructure and mechanical properties of L-DED builds. The post-treatments have proved to be effective in the reduction of segregation in Alloy 718 and thereby enhancing the tensile properties. Recrystallization is achievable by the treatment of L-DED builds at homogenization temperatures. Further, 954 °C – 980 °C ST are influential in increasing the tensile properties of deposits as indicated by the aforementioned research.

19

Tabl

e 2:

Ten

sile

pro

perti

es o

f L-D

ED

Allo

y 71

8

Res

earc

hers

M

ater

ial

Con

ditio

n

0.2%

YS

UT

S D

uctil

ity

Proc

ess

or H

T D

escr

iptio

n M

Pa

MPa

%

EL.

Zha

o et

al.

[79]

AB

590

845

26

1080

°C h

omog

eniz

atio

n +

980

°C S

TA

STA

-GA

11

33

1240

9

Built

usin

g G

A p

owde

r

STA

-PRE

P 13

60

1170

18

Bu

ilt u

sing

PRE

P po

wde

r

Zho

ng e

t al.

[80]

A

B 52

5 84

7 29

10

93 °C

hom

ogen

izat

ion

+ S

TA

STA

/ H

IP+

STA

11

27

1303

16

N

o ch

ange

s on

HIP

at 1

150

°C /

4 hr

Sui e

t al.

[81]

ST

A

992

1245

15

11

00 °C

ST+

stan

dard

agi

ng

Che

n et

al.

[54]

A

B- F

lat to

p 54

3 95

3 23

Fl

at to

p las

er m

ode

AB-

Gau

ssia

n 42

8 77

6 24

G

auss

ian m

ode

pow

er d

istrib

utio

n

Qi e

t al.

[57]

AB*

55

2 90

4 16

V

ertic

al bu

ild o

n an

age

d su

bstra

te

DA

* 10

84

1333

8

STA

* 10

07

1221

16

A

MS

5662

M st

anda

rd sp

ec.

Hom

ogen

ize+

ST

A*

949

1194

20

10

93 °C

hom

ogen

izat

ion+

980

°C S

TA

AM

S 53

83D

stan

dard

spec

.

Xia

o et

al.

[82]

D

A-P

ulse

d 11

21

1404

12

Pu

lsed

mod

e las

er

DA

-CW

10

62

1270

7

cont

inuo

us-w

ave

mod

e las

er

Not

e: A

ll sa

mpl

es a

re h

oriz

onta

lly b

uilt

on th

e su

bstra

te u

nles

s ind

icat

ed b

y th

e A

ster

ix (*

) tha

t rep

rese

nts

a ve

rtica

lly b

uilt

sam

ple.

18

heat treatment involving homogenization and standard solution treatment + aging (STA) enhanced tensile properties of deposits built using GA powder. A further enhancement of tensile properties was reported, especially ductility is improved on using PREP powder as the feedstock. The high deposition rate L-DED experiment by Zhong et al. [19] showed tensile properties similar to the GA powder deposits of Zhao et al. Further, the inclusion of 1150 °C/4 hr Hot Isostatic Pressing (HIP) treatment did not effect changes in tensile properties of specimen [80]. An increase in ST temperature to 1100 °C and consequent double aging showed a decline in tensile properties [81].

Chen et al. [54] showed that varying the laser beam shape has effects on the final properties of L-DED specimens. Flat-top and gaussian mode of power distribution was used in the experiments, wherein the gaussian form yielded a continuous band of laves phase. The flat-top power distribution yielded a higher volume fraction of discrete laves phase which displayed better mechanical properties. Xiao et al. [53] used pulsed and continuous wave (CW) mode of laser operation, that resulted in varied microstructural features that were subjected to direct aging (DA) treatments. Pulsed mode generated equiaxed grains and CW mode yielded a columnar microstructure that affected the mechanical properties. The failure mode is determined to be along the interdendritic region in case of CW specimens which indicated the presence of laves phase on contrary to the pulsed mode specimens that showed a transgranular ductile fracture. The research by Qi et al. [57] shows the anisotropy of L-DED components as samples built vertically perpendicular to the build plate showed better strengths compared to the horizontally built specimens that had better ductility properties.

The research mentioned in Section 3.3 - 3.5 provides an overview of the influence of processing conditions on geometry, microstructure and mechanical properties of L-DED builds. The post-treatments have proved to be effective in the reduction of segregation in Alloy 718 and thereby enhancing the tensile properties. Recrystallization is achievable by the treatment of L-DED builds at homogenization temperatures. Further, 954 °C – 980 °C ST are influential in increasing the tensile properties of deposits as indicated by the aforementioned research.

19

Tabl

e 2:

Ten

sile

pro

perti

es o

f L-D

ED

Allo

y 71

8

Res

earc

hers

M

ater

ial

Con

ditio

n

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21

4 Experiments

4.1 Laser-Directed Energy Deposition Equipment

The L-DED equipment can be broadly divided into laser and laser delivery systems, powder handling system, gas supply system and workbench. A 6kW Yb-doped fiber laser (delivery fiber Ø800 µm) which in-turn uses diode laser modules as a pumping medium is used as the power source in the present set of experiments. The focusing optics consists of a collimator lens and a focus lens having focal lengths of 160 mm and 200 mm. This arrangement of lenses yields a laser spot diameter of 1 mm in focus.

A volumetric powder feeder arrangement coupled with a Fraunhofer ILT- D50 or COAX50- coaxial nozzle is used powder delivery. Argon gas is used for the purposes of shielding the melt-pool and also as a carrier gas for delivering feedstock. The deposition head consisting of the laser optics and the nozzle is housed on a 4-axis CNC machine. Figure 3 shows the schematic arrangement of the L-DED system utilized in the present work.

Figure 3: L-DED system used for building deposits showing the laser head, workbench and powder feeder system. The inset picture (a) shows the coaxial D50 nozzle and (b) shows the laser head assembly including powder delivery and gas supply tubes.

20

21

4 Experiments

4.1 Laser-Directed Energy Deposition Equipment

The L-DED equipment can be broadly divided into laser and laser delivery systems, powder handling system, gas supply system and workbench. A 6kW Yb-doped fiber laser (delivery fiber Ø800 µm) which in-turn uses diode laser modules as a pumping medium is used as the power source in the present set of experiments. The focusing optics consists of a collimator lens and a focus lens having focal lengths of 160 mm and 200 mm. This arrangement of lenses yields a laser spot diameter of 1 mm in focus.

A volumetric powder feeder arrangement coupled with a Fraunhofer ILT- D50 or COAX50- coaxial nozzle is used powder delivery. Argon gas is used for the purposes of shielding the melt-pool and also as a carrier gas for delivering feedstock. The deposition head consisting of the laser optics and the nozzle is housed on a 4-axis CNC machine. Figure 3 shows the schematic arrangement of the L-DED system utilized in the present work.

Figure 3: L-DED system used for building deposits showing the laser head, workbench and powder feeder system. The inset picture (a) shows the coaxial D50 nozzle and (b) shows the laser head assembly including powder delivery and gas supply tubes.

22

4.2 Material

PA Alloy 718 powder is used as feedstock having a chemical composition as shown in Table 3. The powder has a particle size distribution of 45–105 µm with a D50 of 79µm. A wrought solution treated Alloy 718 plate is used as a substrate. Both the powder and substrate plate conform to the specifications of the UNS N07718 standard. Powders have a spherical morphology with the presence of some satellite particles as shown in Figure 4.

Figure 4: PA powder morphology used as feedstock

Table 3:Chemical composition (Wt. %) of Alloy 718 feedstock and substrate plate

Elements Ni Cr Fe Nb+Ta Mo C Ti Al Powder 52.89 18.7 18.52 4.9 2.94 0.05 0.92 0.61

Substrate 53.57 18.7 17.58 4.97 2.89 0.04 0.91 0.59

Elements Co Ta B Cu Mn Si P S Powder 0.11 <0.01 <0.001 <0.1 0.05 0.19 0.005 0.004

Substrate 0.25 0.004 0.002 0.23 0.09 0.06 0.008 0.001

4.3 Design of Experiments

4.3.1 As-built Deposits

Two different experimental design approaches are followed in two articles, Paper A [43] and Paper B [42] appended to this thesis. L9-orthogonal array design and partial factorial method are used in Paper A and Paper B, respectively, to

EXPERIMENTS

23

minimize the number of experimental trials and to obtain qualitative answers with respect to the influence of deposition parameter on geometry, phase content and defects in the single-track deposits.

In Paper A, three process parameters, viz., laser power, deposition speed and powder feed rates are varied in three different levels which results in nine different conditions as shown in Table 4. Other parameters such as laser stand-off distance and laser spot diameter are kept constant at 7.5 mm and 1.3 mm, respectively.

Table 4: Experimental trials designed for work in Paper A as per L9 orthogonal array

Designation Laser Power, W Deposition Speed, mm/min

Feed Rate, g/min

A1 1200 700 13.33 A2 1200 900 16.67 A3 1200 1100 20 A4 1600 700 16.67 A5 1600 900 20 A6 1600 1100 13.33 A7 2000 700 20 A8 2000 900 13.33 A9 2000 1100 16.67

In Paper B, laser power is varied in three different levels; deposition speed and laser stand-off distance are varied in two different levels. An additional parameter set having a laser power of 1800 W is considered in order to validate a few experimental results from Paper A. The resultant tracks from experimental set in Table 4 showed high dilution. Therefore, a suitable consideration for minimizing dilution is considered by increasing stand-off distance and selecting two higher-level scanning speeds [42] from the experimental design of Paper A. The resultant experimental conditions having nine different trials are as recorded in Table 5. The mass flow rate of 20 g/min, is kept constant in experimental conditions pertaining to Paper B. The shielding gas flow rate (12 l/min) and carrier gas flow rate (5 l/min) are kept constant in experimental conditions of both Paper A and Paper B. The single-track deposits were 55 mm long built on an Alloy 718 plate having dimensions of 200 mm x 60 mm x 3 mm.

22

4.2 Material

PA Alloy 718 powder is used as feedstock having a chemical composition as shown in Table 3. The powder has a particle size distribution of 45–105 µm with a D50 of 79µm. A wrought solution treated Alloy 718 plate is used as a substrate. Both the powder and substrate plate conform to the specifications of the UNS N07718 standard. Powders have a spherical morphology with the presence of some satellite particles as shown in Figure 4.

Figure 4: PA powder morphology used as feedstock

Table 3:Chemical composition (Wt. %) of Alloy 718 feedstock and substrate plate

Elements Ni Cr Fe Nb+Ta Mo C Ti Al Powder 52.89 18.7 18.52 4.9 2.94 0.05 0.92 0.61

Substrate 53.57 18.7 17.58 4.97 2.89 0.04 0.91 0.59

Elements Co Ta B Cu Mn Si P S Powder 0.11 <0.01 <0.001 <0.1 0.05 0.19 0.005 0.004

Substrate 0.25 0.004 0.002 0.23 0.09 0.06 0.008 0.001

4.3 Design of Experiments

4.3.1 As-built Deposits

Two different experimental design approaches are followed in two articles, Paper A [43] and Paper B [42] appended to this thesis. L9-orthogonal array design and partial factorial method are used in Paper A and Paper B, respectively, to

EXPERIMENTS

23

minimize the number of experimental trials and to obtain qualitative answers with respect to the influence of deposition parameter on geometry, phase content and defects in the single-track deposits.

In Paper A, three process parameters, viz., laser power, deposition speed and powder feed rates are varied in three different levels which results in nine different conditions as shown in Table 4. Other parameters such as laser stand-off distance and laser spot diameter are kept constant at 7.5 mm and 1.3 mm, respectively.

Table 4: Experimental trials designed for work in Paper A as per L9 orthogonal array

Designation Laser Power, W Deposition Speed, mm/min

Feed Rate, g/min

A1 1200 700 13.33 A2 1200 900 16.67 A3 1200 1100 20 A4 1600 700 16.67 A5 1600 900 20 A6 1600 1100 13.33 A7 2000 700 20 A8 2000 900 13.33 A9 2000 1100 16.67

In Paper B, laser power is varied in three different levels; deposition speed and laser stand-off distance are varied in two different levels. An additional parameter set having a laser power of 1800 W is considered in order to validate a few experimental results from Paper A. The resultant tracks from experimental set in Table 4 showed high dilution. Therefore, a suitable consideration for minimizing dilution is considered by increasing stand-off distance and selecting two higher-level scanning speeds [42] from the experimental design of Paper A. The resultant experimental conditions having nine different trials are as recorded in Table 5. The mass flow rate of 20 g/min, is kept constant in experimental conditions pertaining to Paper B. The shielding gas flow rate (12 l/min) and carrier gas flow rate (5 l/min) are kept constant in experimental conditions of both Paper A and Paper B. The single-track deposits were 55 mm long built on an Alloy 718 plate having dimensions of 200 mm x 60 mm x 3 mm.

24

Table 5: Experimental runs designed for work in Paper B as per the partial factorial approach.

Designation Laser Power, W Deposition Speed, mm/min

Offset Distance, mm

B1 1200 1100 9.5 B2 1600 900 9.5 B3 2000 1100 9.5 B4 1200 1100 13 B5 1600 900 13 B6 2000 1100 13

B7* 1800 1100 13 B8 1600 1100 9.5 B9 1600 1100 13

*B7 trial run is used to determine the optimal conditions for phase fraction analysis

4.3.2 Heat Treatments

In order to study the influence of heat treatment on phase constituents and microstructure, ST and aging steps are performed in accordance with AMS 5663N standard specification [25]. The ST was performed at 954 °C/1 hr followed by air cooling. A double aging treatment typical for Alloy 718 is performed at 718 °C /8 hr followed by furnace cooling at a rate of 56 °C/hr to 621 °C and 621 °C/8 hr, air cooling to room temperature. Three different schedules of heat treatment involving ST, ST+ double aging (STA) and DA involving just the double aging treatments are performed on AB specimens. The treatments were performed in a tube furnace with an argon atmosphere to avoid interference of active gases. The difference between different heat treatments and AB condition are analyzed through volume fraction analysis of Nb-rich precipitates, elemental analysis and hardness testing.

4.4 Material Preparation and Metallography

4.4.1 Sample Preparation and Etching

The single-track deposits were sectioned in three different areas of the build plate for analyzing geometry, microstructure and hardness properties of AB deposit: at 5 mm, 22.5 mm and 50 mm from the deposition start as shown in Figure 5. Three sections were made between the length of 5mm and 22.5mm for the extraction

EXPERIMENTS

25

of specimens for heat treatment. A series of grinding (240 µm and 40 µm) and polishing steps (9 µm, 3 µm and 0.05 µm-silica polishing agent) are carried out post sectioning and mounting of samples. For Microstructural analysis, the samples are electrolytically etched by oxalic acid etchant at 1.5 V, 0.2 A/cm2 for up to 10 s.

Figure 5: A schematic showing the sample extraction scheme followed for L-DED deposits in this work. The samples for geometrical and microstructural analyses were extracted from a distance of 5 mm, 22.5 mm and 50 mm from the start of the deposit depicted by dashed lines. The dotted lines between 5 mm and 22.5 mm depict sample locations of heat-treated specimens.

4.4.2 Microscopy and Image Analysis

Light Optical microscope is utilized to perform basic microstructural characterization. Also, porosity analysis is performed on a fully polished specimen for pores with dimensions higher than 1 µm. Image analysis is performed by using Image-J v1.52p software. The micrographs were assigned threshold values based on image brightness/contrast ratios to obtain binary images and to obtain relative porosity content. Scanning Electron Microscopy (SEM) coupled with Electron Dispersive Spectroscopy (EDS) and Electron Back-scatter Diffraction (EBSD) techniques are used to analyze various microstructural features. SEM images in Back-scattered Electron (BSE) mode are used for determination of volume fraction of phases having high Nb-content, such as laves and NbC in AB sample conditions and delta with laves and NbC in heat-treated conditions, by Image-J software. These phases are generalized as Nb-rich phases hitherto in this report and in the appended articles. EBSD analysis is performed to analyze the texture of grains with respect to selected process conditions, namely B1, B2, B3, B8 and B9 in order to study the effect of three varying process parameters considered in the experimental design.

24

Table 5: Experimental runs designed for work in Paper B as per the partial factorial approach.

Designation Laser Power, W Deposition Speed, mm/min

Offset Distance, mm

B1 1200 1100 9.5 B2 1600 900 9.5 B3 2000 1100 9.5 B4 1200 1100 13 B5 1600 900 13 B6 2000 1100 13

B7* 1800 1100 13 B8 1600 1100 9.5 B9 1600 1100 13

*B7 trial run is used to determine the optimal conditions for phase fraction analysis

4.3.2 Heat Treatments

In order to study the influence of heat treatment on phase constituents and microstructure, ST and aging steps are performed in accordance with AMS 5663N standard specification [25]. The ST was performed at 954 °C/1 hr followed by air cooling. A double aging treatment typical for Alloy 718 is performed at 718 °C /8 hr followed by furnace cooling at a rate of 56 °C/hr to 621 °C and 621 °C/8 hr, air cooling to room temperature. Three different schedules of heat treatment involving ST, ST+ double aging (STA) and DA involving just the double aging treatments are performed on AB specimens. The treatments were performed in a tube furnace with an argon atmosphere to avoid interference of active gases. The difference between different heat treatments and AB condition are analyzed through volume fraction analysis of Nb-rich precipitates, elemental analysis and hardness testing.

4.4 Material Preparation and Metallography

4.4.1 Sample Preparation and Etching

The single-track deposits were sectioned in three different areas of the build plate for analyzing geometry, microstructure and hardness properties of AB deposit: at 5 mm, 22.5 mm and 50 mm from the deposition start as shown in Figure 5. Three sections were made between the length of 5mm and 22.5mm for the extraction

EXPERIMENTS

25

of specimens for heat treatment. A series of grinding (240 µm and 40 µm) and polishing steps (9 µm, 3 µm and 0.05 µm-silica polishing agent) are carried out post sectioning and mounting of samples. For Microstructural analysis, the samples are electrolytically etched by oxalic acid etchant at 1.5 V, 0.2 A/cm2 for up to 10 s.

Figure 5: A schematic showing the sample extraction scheme followed for L-DED deposits in this work. The samples for geometrical and microstructural analyses were extracted from a distance of 5 mm, 22.5 mm and 50 mm from the start of the deposit depicted by dashed lines. The dotted lines between 5 mm and 22.5 mm depict sample locations of heat-treated specimens.

4.4.2 Microscopy and Image Analysis

Light Optical microscope is utilized to perform basic microstructural characterization. Also, porosity analysis is performed on a fully polished specimen for pores with dimensions higher than 1 µm. Image analysis is performed by using Image-J v1.52p software. The micrographs were assigned threshold values based on image brightness/contrast ratios to obtain binary images and to obtain relative porosity content. Scanning Electron Microscopy (SEM) coupled with Electron Dispersive Spectroscopy (EDS) and Electron Back-scatter Diffraction (EBSD) techniques are used to analyze various microstructural features. SEM images in Back-scattered Electron (BSE) mode are used for determination of volume fraction of phases having high Nb-content, such as laves and NbC in AB sample conditions and delta with laves and NbC in heat-treated conditions, by Image-J software. These phases are generalized as Nb-rich phases hitherto in this report and in the appended articles. EBSD analysis is performed to analyze the texture of grains with respect to selected process conditions, namely B1, B2, B3, B8 and B9 in order to study the effect of three varying process parameters considered in the experimental design.

26

4.5 Single-Track Characterization

The single-track deposits are characterized based on the differences in geometry, microstructure and hardness. A transverse cross-section of a single-track deposit is as shown in Figure 6. The geometrical analysis includes the measurements of height (H), width (W) and depth (D) of deposits. Further, the area of deposition (Ad) and the area of dilution (As) are measured and compared with different process parameters and combined process parameters. Some of the relationships between process parameters are:

Lm = M/V (g/mm) Equation 1

HI= P/V (J/mm) Equation 2

SE= HI/Ds (J/mm2) Equation 3

η= ρ x V x As/M Equation 4

Where M is feed rate, Lm is line mass, HI is heat input, V is deposition speed, P is laser power, SE is specific energy, Ds is laser spot diameter, η is powder capture efficiency and ρ is the density of Alloy 718.

Figure 6: A typical transverse cross-section of a single-track deposit indicating geometrical features and division of different regions [42].

27

Based on the differences observed in microstructural features, the cross-section of a single-track deposit is divided into three different areas: (i) Top region; (ii) Middle region; (iii) Bottom region, as shown in Figure 6. The basis of this division is further discussed in Section 5.2. The volumetric phase analysis is performed discretely in these three regions for all specimen conditions. Image analysis is used to determine the microstructural features such as grain size and primary dendritic arm spacing (PDAS). The grain-size analysis is performed in accordance with ASTM E112-13 standard specification [83]. Since the grains in the middle and bottom regions were columnar and elongated in nature, the mean intercept length is calculated in three different axes of the grain to measure the average grain-size. Vickers micro-indentation hardness testing is performed in accordance with ASTM E384-17 [84] with a load of 200 gf applied for 15 s in top, middle and bottom regions of AB and heat-treated conditions.

26

4.5 Single-Track Characterization

The single-track deposits are characterized based on the differences in geometry, microstructure and hardness. A transverse cross-section of a single-track deposit is as shown in Figure 6. The geometrical analysis includes the measurements of height (H), width (W) and depth (D) of deposits. Further, the area of deposition (Ad) and the area of dilution (As) are measured and compared with different process parameters and combined process parameters. Some of the relationships between process parameters are:

Lm = M/V (g/mm) Equation 1

HI= P/V (J/mm) Equation 2

SE= HI/Ds (J/mm2) Equation 3

η= ρ x V x As/M Equation 4

Where M is feed rate, Lm is line mass, HI is heat input, V is deposition speed, P is laser power, SE is specific energy, Ds is laser spot diameter, η is powder capture efficiency and ρ is the density of Alloy 718.

Figure 6: A typical transverse cross-section of a single-track deposit indicating geometrical features and division of different regions [42].

27

Based on the differences observed in microstructural features, the cross-section of a single-track deposit is divided into three different areas: (i) Top region; (ii) Middle region; (iii) Bottom region, as shown in Figure 6. The basis of this division is further discussed in Section 5.2. The volumetric phase analysis is performed discretely in these three regions for all specimen conditions. Image analysis is used to determine the microstructural features such as grain size and primary dendritic arm spacing (PDAS). The grain-size analysis is performed in accordance with ASTM E112-13 standard specification [83]. Since the grains in the middle and bottom regions were columnar and elongated in nature, the mean intercept length is calculated in three different axes of the grain to measure the average grain-size. Vickers micro-indentation hardness testing is performed in accordance with ASTM E384-17 [84] with a load of 200 gf applied for 15 s in top, middle and bottom regions of AB and heat-treated conditions.

28

29

5 Summary of Results

5.1 Geometry of Single-Track Deposits

A brief overview of geometrical analysis from the appended articles [42,43] shows that the height of the deposits is primarily influenced by feed rate and deposition speed. The width and depth of the deposits are predominantly influenced by laser power and moderately affected by deposition speed and feed rate. The width, depth and dilution decrease with an increase in laser stand-off distance which in turn depends on the laser spot diameter, whereas the height of deposits is unaffected by stand-off distance and spot-diameter. The summary of the effects of an increase in process parameters (depicted by ‘▲’) on height, width and depth of deposits from Paper A and Paper B are summarized in Table 6. The effect of process parameters on the geometrical measurements is categorized into three levels: (i.) significantly effective marked by ‘(+)’; (ii.) moderately effective marked by ‘(=)’; and ineffective which is unmarked.

Line mass is defined as the amount of material delivered by the deposition head per unit length (Equation 1). An increase in feed rate or a decrease in deposition speed increases the material available for deposition. If the amount of material is higher i.e., increase in line mass results in an increase in height of deposits. Laser power and laser stand-off distances influence the energy distribution, which determines the amount of powder that can be melted. If the energy density or intensity (See Table 3 in Paper B) is higher than that required to melt the powder fed, the excess energy derived wither from the increase in power or by the decrease in stand-off distance proves ineffective in determining the height of deposits [12].

The width of the deposits mainly depends on the powder focus diameter and laser spot diameter. Powder focus is kept constant throughout all the experiments as it is a function of nozzle diameter, feed rate and powder particle size. Laser spot diameter increases with an increase in laser stand-off distance. Hence increasing spot diameter increases the width of the deposit at nominal feed rates. An increase in laser power affects the larger area of the substrate and increases the size of the melt-pool due to the increased forces acting on easily deformable melt, thereby increasing the width of the deposit [71]. Speed and feed rates have a minor influence on the width of deposits. Basically, increase in speed and feed rate decreases the heat input and hence decreases the size of the melt-pool and width of the deposit.

28

29

5 Summary of Results

5.1 Geometry of Single-Track Deposits

A brief overview of geometrical analysis from the appended articles [42,43] shows that the height of the deposits is primarily influenced by feed rate and deposition speed. The width and depth of the deposits are predominantly influenced by laser power and moderately affected by deposition speed and feed rate. The width, depth and dilution decrease with an increase in laser stand-off distance which in turn depends on the laser spot diameter, whereas the height of deposits is unaffected by stand-off distance and spot-diameter. The summary of the effects of an increase in process parameters (depicted by ‘▲’) on height, width and depth of deposits from Paper A and Paper B are summarized in Table 6. The effect of process parameters on the geometrical measurements is categorized into three levels: (i.) significantly effective marked by ‘(+)’; (ii.) moderately effective marked by ‘(=)’; and ineffective which is unmarked.

Line mass is defined as the amount of material delivered by the deposition head per unit length (Equation 1). An increase in feed rate or a decrease in deposition speed increases the material available for deposition. If the amount of material is higher i.e., increase in line mass results in an increase in height of deposits. Laser power and laser stand-off distances influence the energy distribution, which determines the amount of powder that can be melted. If the energy density or intensity (See Table 3 in Paper B) is higher than that required to melt the powder fed, the excess energy derived wither from the increase in power or by the decrease in stand-off distance proves ineffective in determining the height of deposits [12].

The width of the deposits mainly depends on the powder focus diameter and laser spot diameter. Powder focus is kept constant throughout all the experiments as it is a function of nozzle diameter, feed rate and powder particle size. Laser spot diameter increases with an increase in laser stand-off distance. Hence increasing spot diameter increases the width of the deposit at nominal feed rates. An increase in laser power affects the larger area of the substrate and increases the size of the melt-pool due to the increased forces acting on easily deformable melt, thereby increasing the width of the deposit [71]. Speed and feed rates have a minor influence on the width of deposits. Basically, increase in speed and feed rate decreases the heat input and hence decreases the size of the melt-pool and width of the deposit.

30

The depth of the deposit is a function of energy available for melting the substrate and the alloying composition. The amount of total energy in the process can be divided into energy utilized for melting substrate and powder feed-stock with some energy losses [85]. The energy is a function of laser power, spot diameter and scanning speed. Increase in power or decrease in speed and/or spot diameter increases the energy density for deposition. Further, increase in feed rates increase the energy utilized in melting the powder and the low remnant energy is used for melting the substrate, thereby decreasing the depth of the deposits.

Table 6: Effect of increase in process parameters on height, width and depth of single-track deposit

Laser power

▲ Deposition speed ▲

Feed rate ▲ Laser stand-off distance ▲

Height Ineffective Decreases (+) Increases (+) Ineffective Width Increases (+) Decreases (=) Decreases (=) Increases (+) Depth Increases (+) Increases (+) Decreases (=) Decreases (=)

Note: The symbols (+) and (=) indicates significant and moderately effective parameters on geometry

5.2 Microstructure and Phase Constituents

5.2.1 As-built Deposits

3D-builds from the L-DED process are shown to have varied microstructural features in different areas along the height of the builds [86]. A similar distinction can be noticed in single-track deposits analyzed in Paper A and Paper B. A deposit can be distinguished into the top, middle and bottom regions (indicated in Figure 6). The solidification starts at the bottom region by epitaxial growth of grains from the substrate which transforms into columnar dendritic grains due to the composition of Alloy 718 and high thermal gradient as depicted in Figure 7c. This region consists of fine columnar grains with an average size of 45 µm having a dendritic microstructure as shown in Figure 7f. The PDAS of 2 µm - 5 µm is measured for the grains in the bottom region. In the middle region, the grains have an average size of 80 µm, predominantly having a columnar dendritic microstructure with a PDAS varying between 4 µm – 8.5 µm. The top region of the deposit has an equiaxed structure with an average grain size of 20 µm. A plot

SUMMARY OF RESULTS

31

depicting variation in PDAS in the bottom and middle regions are shown in Figure 8.

Figure 7: LOM images in (a) top region (b) middle region (c) bottom region; and SEM image in (d) top region (e) middle region (f) bottom region of a single-track deposit showing varied microstructural features. The arrows in Figure 7a,c indicate primary and secondary solidification directions, respectively and the dotted line in 7a indicates the Columnar to Equiaxed Transition (CET).

Two different solidification directions are marked in Figure 7a and Figure 7c, one primary front starting from the bottom and a secondary front from the top. The primary solidification front arises due to the undercooling offered by a relative difference in temperatures of substrate and molten metal. The secondary front arises due to the heterogeneous nucleation of precipitates effected by constitutional supercooling. Another possibility for heterogeneous nucleation is the precipitation of high melting phases such as TiN and TiC that could act as intrinsic inoculants [87]. The earlier phenomenon of constitutional effects is determined due to the variation of amount of Nb-rich phases present in the top region compared to the bottom and the middle regions of the deposits. However, intrinsic inoculation effects cannot be accurately determined in this work. The number of nucleation sites proves important in deciding the grain size in the top region of deposit. In the current study, small grain size at the top indicate many such nucleation sites stabilizing predominantly equiaxed grain structure as seen from EBSD representation in Figure 9. The transition along the height of the deposit where a columnar growth changes into equiaxed grains, known as CET is marked by dotted lines in Figure 7a and Figure 9.

30

The depth of the deposit is a function of energy available for melting the substrate and the alloying composition. The amount of total energy in the process can be divided into energy utilized for melting substrate and powder feed-stock with some energy losses [85]. The energy is a function of laser power, spot diameter and scanning speed. Increase in power or decrease in speed and/or spot diameter increases the energy density for deposition. Further, increase in feed rates increase the energy utilized in melting the powder and the low remnant energy is used for melting the substrate, thereby decreasing the depth of the deposits.

Table 6: Effect of increase in process parameters on height, width and depth of single-track deposit

Laser power

▲ Deposition speed ▲

Feed rate ▲ Laser stand-off distance ▲

Height Ineffective Decreases (+) Increases (+) Ineffective Width Increases (+) Decreases (=) Decreases (=) Increases (+) Depth Increases (+) Increases (+) Decreases (=) Decreases (=)

Note: The symbols (+) and (=) indicates significant and moderately effective parameters on geometry

5.2 Microstructure and Phase Constituents

5.2.1 As-built Deposits

3D-builds from the L-DED process are shown to have varied microstructural features in different areas along the height of the builds [86]. A similar distinction can be noticed in single-track deposits analyzed in Paper A and Paper B. A deposit can be distinguished into the top, middle and bottom regions (indicated in Figure 6). The solidification starts at the bottom region by epitaxial growth of grains from the substrate which transforms into columnar dendritic grains due to the composition of Alloy 718 and high thermal gradient as depicted in Figure 7c. This region consists of fine columnar grains with an average size of 45 µm having a dendritic microstructure as shown in Figure 7f. The PDAS of 2 µm - 5 µm is measured for the grains in the bottom region. In the middle region, the grains have an average size of 80 µm, predominantly having a columnar dendritic microstructure with a PDAS varying between 4 µm – 8.5 µm. The top region of the deposit has an equiaxed structure with an average grain size of 20 µm. A plot

SUMMARY OF RESULTS

31

depicting variation in PDAS in the bottom and middle regions are shown in Figure 8.

Figure 7: LOM images in (a) top region (b) middle region (c) bottom region; and SEM image in (d) top region (e) middle region (f) bottom region of a single-track deposit showing varied microstructural features. The arrows in Figure 7a,c indicate primary and secondary solidification directions, respectively and the dotted line in 7a indicates the Columnar to Equiaxed Transition (CET).

Two different solidification directions are marked in Figure 7a and Figure 7c, one primary front starting from the bottom and a secondary front from the top. The primary solidification front arises due to the undercooling offered by a relative difference in temperatures of substrate and molten metal. The secondary front arises due to the heterogeneous nucleation of precipitates effected by constitutional supercooling. Another possibility for heterogeneous nucleation is the precipitation of high melting phases such as TiN and TiC that could act as intrinsic inoculants [87]. The earlier phenomenon of constitutional effects is determined due to the variation of amount of Nb-rich phases present in the top region compared to the bottom and the middle regions of the deposits. However, intrinsic inoculation effects cannot be accurately determined in this work. The number of nucleation sites proves important in deciding the grain size in the top region of deposit. In the current study, small grain size at the top indicate many such nucleation sites stabilizing predominantly equiaxed grain structure as seen from EBSD representation in Figure 9. The transition along the height of the deposit where a columnar growth changes into equiaxed grains, known as CET is marked by dotted lines in Figure 7a and Figure 9.

32

Figure 8: Plot showing PDAS for middle and bottom regions

Figure 9: EBSD pole figure indicating grain texture of deposit with process condition B1 (Table 5) and the pole figure representing dominant (100) directional growth of

crystals. The dotted line in the picture shows CET [42].

The Nb-rich phases such as laves and NbC appear as bright white particles in BSE-mode SEM images as seen in Figure 7d-f. The volumetric analysis shows an increase in the amount of Nb-rich phases from the bottom to the top of the deposit. The result of phase fraction analysis can be seen in both articles appended (Figure 8 of Paper A and Figure 8 of Paper B). In both articles, increased laser power is found to decrease the amount of segregation in the top region of deposit as high-power conditions suppress the CET and thereby constitutional supercooling. In the bottom and middle regions, cooling rates are found to affect the segregation and hence lower segregation levels are achieved at higher deposition speeds as discussed in Paper B.

SUMMARY OF RESULTS

33

5.2.2 Heat-Treatment of Deposits

The results of three different heat treatments viz., ST, STA and DA of AB deposits (conditions mentioned under Section 4.3.2) are discussed in Paper A. The ST in aerospace applications are generally performed in lower delta solvus temperatures which yield controlled precipitation of δ-phase [24]. δ-phase aids in the control of grain size as mentioned in the earlier sections and thereby influences the tensile strength, fatigue strength and stress-rupture ductility properties.

Figure 10: SEM images showing low amounts of Nb-rich precipitates in processing condition A2 (Table 4) upon (a) solution treatment; (b) solution treatment + aging; (c) direct aging; compared to Image (d) depicting as-built deposit microstructure

The SEM images indicating the effects of ST, STA and DA are depicted in Figure 10a-c, respectively. Partial dissolution of Nb-rich phases is noticed on comparison of Nb-rich phases present in AB (Figure 10d) and heat-treated conditions. Even after aging (at relatively low temperatures of 720°C), the dissolution of Nb-rich phases is evident from both volume fraction analysis and EDS point-ID analysis of Nb element in the matrix (See Figure 8 in Paper A). The DA condition consistently showed a lower volume fraction of Nb-rich phases (Figure 10c) compared to AB specimens of respective processing conditions.

32

Figure 8: Plot showing PDAS for middle and bottom regions

Figure 9: EBSD pole figure indicating grain texture of deposit with process condition B1 (Table 5) and the pole figure representing dominant (100) directional growth of

crystals. The dotted line in the picture shows CET [42].

The Nb-rich phases such as laves and NbC appear as bright white particles in BSE-mode SEM images as seen in Figure 7d-f. The volumetric analysis shows an increase in the amount of Nb-rich phases from the bottom to the top of the deposit. The result of phase fraction analysis can be seen in both articles appended (Figure 8 of Paper A and Figure 8 of Paper B). In both articles, increased laser power is found to decrease the amount of segregation in the top region of deposit as high-power conditions suppress the CET and thereby constitutional supercooling. In the bottom and middle regions, cooling rates are found to affect the segregation and hence lower segregation levels are achieved at higher deposition speeds as discussed in Paper B.

SUMMARY OF RESULTS

33

5.2.2 Heat-Treatment of Deposits

The results of three different heat treatments viz., ST, STA and DA of AB deposits (conditions mentioned under Section 4.3.2) are discussed in Paper A. The ST in aerospace applications are generally performed in lower delta solvus temperatures which yield controlled precipitation of δ-phase [24]. δ-phase aids in the control of grain size as mentioned in the earlier sections and thereby influences the tensile strength, fatigue strength and stress-rupture ductility properties.

Figure 10: SEM images showing low amounts of Nb-rich precipitates in processing condition A2 (Table 4) upon (a) solution treatment; (b) solution treatment + aging; (c) direct aging; compared to Image (d) depicting as-built deposit microstructure

The SEM images indicating the effects of ST, STA and DA are depicted in Figure 10a-c, respectively. Partial dissolution of Nb-rich phases is noticed on comparison of Nb-rich phases present in AB (Figure 10d) and heat-treated conditions. Even after aging (at relatively low temperatures of 720°C), the dissolution of Nb-rich phases is evident from both volume fraction analysis and EDS point-ID analysis of Nb element in the matrix (See Figure 8 in Paper A). The DA condition consistently showed a lower volume fraction of Nb-rich phases (Figure 10c) compared to AB specimens of respective processing conditions.

34

The needle-like δ-phases are noticed in ST and STA conditions but not in the DA specimens, as the nucleation of δ-precipitates accelerate in temperatures range of 900 °C-950 °C [34-36]. δ-precipitate density is found to be higher in the bottom and top region of the deposits. The grain size in the top and bottom are relatively smaller, aiding the diffusion of Nb necessary for precipitation of δ-phase compared to the middle region of deposits. The eutectic precipitates in the top region provide the necessary Nb for δ-precipitates (Figure 5a and 6a in Paper A), whereas the Nb in the bottom region is provided by the partial dissolution of laves phase and Nb present in the matrix.

5.3 Hardness Testing

The results from hardness testing of deposits in different conditions summarized in Paper A is shown in Figure 11. The average hardness in AB conditions varies between 270-350 HV0.2 with top regions showing slightly higher hardness compared to the middle and the bottom regions of deposits (Figure 11a). This variation in hardness in the top region is attributable to the higher amount of Nb-rich precipitates inhibiting the deformation of material and due to relatively smaller grain size in the top region of the deposits. On ST, the average hardness level decreases (Figure 11b) due to the dissolution of Nb-rich precipitates and γʹ precipitates formed during the deposition process as compared to AB deposits [88].

The aged (STA and DA) conditions show higher hardness values due to the presence of strengthening precipitates, although the influence of process parameters is adjudged to be higher in DA condition compared to STA condition, based on the variations seen in Figure 11c,d. Homogeneous precipitation of γʹ and γʺ can be envisioned in STA condition due to the formation of near solid-solution compared to DA condition. For example, condition A2 in Figure 11d pertaining to DA condition displays very low hardness. This specimen had the highest amount of Nb-rich phases in the AB condition (Figure 8 in Paper A). Since most of the Nb present in the alloy participate in the formation of Nb-rich phases even after partial dissolution of these phases, relatively lower amount of strengthening precipitate is expected upon aging, yielding a lower hardness response. The A2 condition undergoing STA treatment (Figure 11c) shows average hardness value of 450 HV par with other specimen conditions due to higher dissolution compared to DA condition, which in turn results in a higher amount of γʹ and γʺ precipitates.

SUMMARY OF RESULTS

35

34

The needle-like δ-phases are noticed in ST and STA conditions but not in the DA specimens, as the nucleation of δ-precipitates accelerate in temperatures range of 900 °C-950 °C [34-36]. δ-precipitate density is found to be higher in the bottom and top region of the deposits. The grain size in the top and bottom are relatively smaller, aiding the diffusion of Nb necessary for precipitation of δ-phase compared to the middle region of deposits. The eutectic precipitates in the top region provide the necessary Nb for δ-precipitates (Figure 5a and 6a in Paper A), whereas the Nb in the bottom region is provided by the partial dissolution of laves phase and Nb present in the matrix.

5.3 Hardness Testing

The results from hardness testing of deposits in different conditions summarized in Paper A is shown in Figure 11. The average hardness in AB conditions varies between 270-350 HV0.2 with top regions showing slightly higher hardness compared to the middle and the bottom regions of deposits (Figure 11a). This variation in hardness in the top region is attributable to the higher amount of Nb-rich precipitates inhibiting the deformation of material and due to relatively smaller grain size in the top region of the deposits. On ST, the average hardness level decreases (Figure 11b) due to the dissolution of Nb-rich precipitates and γʹ precipitates formed during the deposition process as compared to AB deposits [88].

The aged (STA and DA) conditions show higher hardness values due to the presence of strengthening precipitates, although the influence of process parameters is adjudged to be higher in DA condition compared to STA condition, based on the variations seen in Figure 11c,d. Homogeneous precipitation of γʹ and γʺ can be envisioned in STA condition due to the formation of near solid-solution compared to DA condition. For example, condition A2 in Figure 11d pertaining to DA condition displays very low hardness. This specimen had the highest amount of Nb-rich phases in the AB condition (Figure 8 in Paper A). Since most of the Nb present in the alloy participate in the formation of Nb-rich phases even after partial dissolution of these phases, relatively lower amount of strengthening precipitate is expected upon aging, yielding a lower hardness response. The A2 condition undergoing STA treatment (Figure 11c) shows average hardness value of 450 HV par with other specimen conditions due to higher dissolution compared to DA condition, which in turn results in a higher amount of γʹ and γʺ precipitates.

SUMMARY OF RESULTS

35

36

Figure 11: Results from Micro-Vickers hardness testing (HV0.2) of various sample conditions (A1-A9) in the top, middle and bottom regions in (a) as-built; (b) solution treated; (c) solution treated + aged; (d) direct aged conditions

37

6 Conclusions and Future Work

6.1 Conclusions

The present research focusses on investigating the influence of process parameters such as laser power, deposition speed, feed rate and stand-off distance on the geometrical and microstructural characteristics of L-DED deposits. The post-build heat treatment was performed to analyze the effect of solution and aging treatments on microstructure and segregation of the alloy. The following conclusions can be drawn based on the results obtained:

1. The height of the deposit is primarily affected by feed rate and deposition speed. It increases with an increase in feed rate and decreases with an increase in deposition speed. The width and depth of the deposits increase with an increase in laser power. Width increases and depth decreases with an increase in stand-off distance and feed rates. Deposition speed has minimal effect on the width of the deposit but an increase in speed is found to increase the depth of deposits.

2. A single-track deposit can be divided into top, middle and bottom regions based on the microstructural features that arise due to the varied solidification conditions. The bottom and middle regions consist of columnar grains predominantly growing in (100) direction, whereas the top region has equiaxed grains having a varied texture. The volume of Nb-rich precipitates is higher in the top region compared to the bottom and middle regions of the deposit.

3. High power condition (1600 W and 2000 W) and high deposition speed (1100 mm/min) result in lower precipitation of Nb-rich phases and hence lower volume fractions. High power condition is particularly influential in decreasing the Nb-rich phases in the top region of the deposits. High deposition speed affects the segregation in the middle and bottom regions of the deposit.

4. High power conditions yield low process porosity in the deposited specimens. 5. Solution treatment (954 °C/1 hr), solution treatment (954 °C/1 hr) + aging

(718 °C/8 hr + 621 °C/8 hr) and direct aging (718 °C/8 hr + 621 °C/8 hr) of deposits lowers the volume of secondary Nb-rich phases compared to as-built deposits, indicating a partial dissolution of laves phase. Solution treatments aid in the nucleation of delta precipitates that grows during subsequent aging treatment. Direct aging performed at relatively low temperatures is found to be insufficient for precipitation of delta phase.

36

Figure 11: Results from Micro-Vickers hardness testing (HV0.2) of various sample conditions (A1-A9) in the top, middle and bottom regions in (a) as-built; (b) solution treated; (c) solution treated + aged; (d) direct aged conditions

37

6 Conclusions and Future Work

6.1 Conclusions

The present research focusses on investigating the influence of process parameters such as laser power, deposition speed, feed rate and stand-off distance on the geometrical and microstructural characteristics of L-DED deposits. The post-build heat treatment was performed to analyze the effect of solution and aging treatments on microstructure and segregation of the alloy. The following conclusions can be drawn based on the results obtained:

1. The height of the deposit is primarily affected by feed rate and deposition speed. It increases with an increase in feed rate and decreases with an increase in deposition speed. The width and depth of the deposits increase with an increase in laser power. Width increases and depth decreases with an increase in stand-off distance and feed rates. Deposition speed has minimal effect on the width of the deposit but an increase in speed is found to increase the depth of deposits.

2. A single-track deposit can be divided into top, middle and bottom regions based on the microstructural features that arise due to the varied solidification conditions. The bottom and middle regions consist of columnar grains predominantly growing in (100) direction, whereas the top region has equiaxed grains having a varied texture. The volume of Nb-rich precipitates is higher in the top region compared to the bottom and middle regions of the deposit.

3. High power condition (1600 W and 2000 W) and high deposition speed (1100 mm/min) result in lower precipitation of Nb-rich phases and hence lower volume fractions. High power condition is particularly influential in decreasing the Nb-rich phases in the top region of the deposits. High deposition speed affects the segregation in the middle and bottom regions of the deposit.

4. High power conditions yield low process porosity in the deposited specimens. 5. Solution treatment (954 °C/1 hr), solution treatment (954 °C/1 hr) + aging

(718 °C/8 hr + 621 °C/8 hr) and direct aging (718 °C/8 hr + 621 °C/8 hr) of deposits lowers the volume of secondary Nb-rich phases compared to as-built deposits, indicating a partial dissolution of laves phase. Solution treatments aid in the nucleation of delta precipitates that grows during subsequent aging treatment. Direct aging performed at relatively low temperatures is found to be insufficient for precipitation of delta phase.

38

6. The delta phase density is higher in the bottom and top regions of the deposit compared to the middle region, attributable to smaller diffusion lengths in top and bottom regions.

6.2 Future Work

To evaluate the influence of process parameters on deposit walls and layered deposition on the substrate. Apart from preliminary process parameters considered in this study, parameters such as the deposition sequence, dwell time and overlap conditions are influential in deciding the properties of a final 3D- builds.

The results from the heat treatment are a mere interpretation of apparent changes in the material. Hence further heat treatments including delta dumping need to be carried out in the context of investigating notch sensitivity properties of L-DED built Alloy 718. Further, mechanical property evaluation is planned to be carried out by notch-rupture tests.

Local heat treatment on L-DED repairs is envisaged using an induction treatment.

39

References

1. Kerrebrock, J.L. Aircraft Engines and Gas Turbines, Second Edition ed.; MIT Press: 1992.

2. Sims, C.T.; Stoloff, N.S.; Hagel, W.C. Superalloys Ii: High-Temperature Materials for Aerospace and Industrial Power; John Wiley & Sons: 1987.

3. The Jet Engine, Fifth ed.; The Technical Publications Department, Rolls-Royce plc: Derby, England, 1986.

4. Paulonis, D.F.; Schirra, J.J. Alloy 718 at Pratt & Whitney: Historical Perspective and Future Challenges. In Proceedings of Superalloys 718, 625, 706 and Various Derivatives. DOI: https://doi.org/10.7449/2001/superalloys_2001_13_23.

5. Barker, J.F. The Initial Years of Alloy 718: A Ge Perspective. In Proceedings of Superalloy 719--Metallurgy and Applications DOI: https://doi.org/10.7449/1989/Superalloys_1989_269_277.

6. Additive Manufacturing- General Principles- Terminology: Iso/Astm 52900:2017; Swedish Standards Institute 2017.

7. Gibson, I.; Rosen, D.W.; Stucker, B. Additive Manufacturing Technologies; 2010; 10.1007/978-1-4419-1120-9. DOI: https://doi.org/10.1007/978-1-4419-1120-9.

8. Uckelmann, I. ‘Buy-to-Fly’ Ratio Cutting Costs with Metal 3d Printing. Availabe online: https://www.materialise.com/en/manufacturing/whitepaper-buy-to-fly-ratio-cutting-costs-metal-3d-printing (accessed on 30/4/2020).

9. Corbin, D.J.; Nassar, A.R.; Reutzel, E.W.; Beese, A.M.; Kistler, N.A. Effect of Directed Energy Deposition Processing Parameters on Laser Deposited Inconel® 718: External Morphology. Journal of Laser Applications 2017, 29. DOI: https://doi.org/10.2351/1.4977476.

10. Graf, B.; Ammer, S.; Gumenyuk, A.; Rethmeier, M. Design of Experiments for Laser Metal Deposition in Maintenance, Repair and Overhaul Applications. Procedia CIRP 2013, 11, 245-248. DOI: https://doi.org/10.1016/j.procir.2013.07.031.

38

6. The delta phase density is higher in the bottom and top regions of the deposit compared to the middle region, attributable to smaller diffusion lengths in top and bottom regions.

6.2 Future Work

To evaluate the influence of process parameters on deposit walls and layered deposition on the substrate. Apart from preliminary process parameters considered in this study, parameters such as the deposition sequence, dwell time and overlap conditions are influential in deciding the properties of a final 3D- builds.

The results from the heat treatment are a mere interpretation of apparent changes in the material. Hence further heat treatments including delta dumping need to be carried out in the context of investigating notch sensitivity properties of L-DED built Alloy 718. Further, mechanical property evaluation is planned to be carried out by notch-rupture tests.

Local heat treatment on L-DED repairs is envisaged using an induction treatment.

39

References

1. Kerrebrock, J.L. Aircraft Engines and Gas Turbines, Second Edition ed.; MIT Press: 1992.

2. Sims, C.T.; Stoloff, N.S.; Hagel, W.C. Superalloys Ii: High-Temperature Materials for Aerospace and Industrial Power; John Wiley & Sons: 1987.

3. The Jet Engine, Fifth ed.; The Technical Publications Department, Rolls-Royce plc: Derby, England, 1986.

4. Paulonis, D.F.; Schirra, J.J. Alloy 718 at Pratt & Whitney: Historical Perspective and Future Challenges. In Proceedings of Superalloys 718, 625, 706 and Various Derivatives. DOI: https://doi.org/10.7449/2001/superalloys_2001_13_23.

5. Barker, J.F. The Initial Years of Alloy 718: A Ge Perspective. In Proceedings of Superalloy 719--Metallurgy and Applications DOI: https://doi.org/10.7449/1989/Superalloys_1989_269_277.

6. Additive Manufacturing- General Principles- Terminology: Iso/Astm 52900:2017; Swedish Standards Institute 2017.

7. Gibson, I.; Rosen, D.W.; Stucker, B. Additive Manufacturing Technologies; 2010; 10.1007/978-1-4419-1120-9. DOI: https://doi.org/10.1007/978-1-4419-1120-9.

8. Uckelmann, I. ‘Buy-to-Fly’ Ratio Cutting Costs with Metal 3d Printing. Availabe online: https://www.materialise.com/en/manufacturing/whitepaper-buy-to-fly-ratio-cutting-costs-metal-3d-printing (accessed on 30/4/2020).

9. Corbin, D.J.; Nassar, A.R.; Reutzel, E.W.; Beese, A.M.; Kistler, N.A. Effect of Directed Energy Deposition Processing Parameters on Laser Deposited Inconel® 718: External Morphology. Journal of Laser Applications 2017, 29. DOI: https://doi.org/10.2351/1.4977476.

10. Graf, B.; Ammer, S.; Gumenyuk, A.; Rethmeier, M. Design of Experiments for Laser Metal Deposition in Maintenance, Repair and Overhaul Applications. Procedia CIRP 2013, 11, 245-248. DOI: https://doi.org/10.1016/j.procir.2013.07.031.

40

11. Steen, W.M.; Courtney, C.G.H. Hardfacing of Nimonic 75 Using 2 Kw Continuouswave Co2laser. Metals Technology 2013, 7, 232-237. DOI: https://doi.org/10.1179/030716980803286955.

12. de Oliveira, U.; Ocelík, V.; De Hosson, J.T.M. Analysis of Coaxial Laser Cladding Processing Conditions. Surface and Coatings Technology 2005, 197, 127-136. DOI: https://doi.org/10.1016/j.surfcoat.2004.06.029.

13. Segerstark, A.; Andersson, J.; Svensson, L.-E. Investigation of Laser Metal Deposited Alloy 718 onto an En 1.4401 Stainless Steel Substrate. Optics & Laser Technology 2017, 97, 144-153. DOI: https://doi.org/10.1016/j.optlastec.2017.05.038.

14. Chen, B.; Mazumder, J. Role of Process Parameters During Additive Manufacturing by Direct Metal Deposition of Inconel 718. Rapid Prototyping Journal 2017, 23, 919-929. DOI: https://doi.org/10.1108/rpj-05-2016-0071.

15. Pinkerton, A.J. Advances in the Modeling of Laser Direct Metal Deposition. Journal of Laser Applications 2015, 27. DOI: https://doi.org/10.2351/1.4815992.

16. Zhong, C.; Gasser, A.; Kittel, J.; Schopphoven, T.; Pirch, N.; Fu, J.; Poprawe, R. Study of Process Window Development for High Deposition-Rate Laser Material Deposition by Using Mixed Processing Parameters. Journal of Laser Applications 2015, 27. DOI: https://doi.org/10.2351/1.4919804.

17. Thompson, S.M.; Bian, L.; Shamsaei, N.; Yadollahi, A. An Overview of Direct Laser Deposition for Additive Manufacturing; Part I: Transport Phenomena, Modeling and Diagnostics. Additive Manufacturing 2015, 8, 36-62. DOI: https://doi.org/10.1016/j.addma.2015.07.001.

18. Zhong, C.; Kittel, J.; Gasser, A.; Schleifenbaum, J.H. Study of Nickel-Based Super-Alloys Inconel 718 and Inconel 625 in High-Deposition-Rate Laser Metal Deposition. Optics & Laser Technology 2019, 109, 352-360. DOI: https://doi.org/10.1016/j.optlastec.2018.08.003.

19. Zhong, C.; Gasser, A.; Kittel, J.; Fu, J.; Ding, Y.; Poprawe, R. Microstructures and Tensile Properties of Inconel 718 Formed by High Deposition-Rate Laser Metal Deposition. Journal of Laser Applications 2016, 28. DOI: https://doi.org/10.2351/1.4943290.

20. Dass, A.; Moridi, A. State of the Art in Directed Energy Deposition: From Additive Manufacturing to Materials Design. Coatings 2019, 9. DOI: https://doi.org/10.3390/coatings9070418.

REFERENCES

41

21. Zhang, Y.; Li, Z.; Nie, P.; Wu, Y. Effect of Heat Treatment on Niobium Segregation of Laser-Cladded In718 Alloy Coating. Metallurgical and Materials Transactions A 2012, 44, 708-716. DOI: https://doi.org/10.1007/s11661-012-1459-z.

22. Long, Y.-t.; Nie, P.-l.; Li, Z.-g.; Huang, J.; Li, X.; Xu, X.-m. Segregation of Niobium in Laser Cladding Inconel 718 Superalloy. Transactions of Nonferrous Metals Society of China 2016, 26, 431-436. DOI: https://doi.org/10.1016/s1003-6326(16)64131-6.

23. Liu, F.; Lin, X.; Song, M.; Zhao, W.; Chen, J.; Huang, W. Effect of Intermediate Heat Treatment Temperature on Microstructure and Notch Sensitivity of Laser Solid Formed Inconel 718 Superalloy. Journal of Wuhan University of Technology-Mater. Sci. Ed. 2011, 26, 908-913. DOI: https://doi.org/10.1007/s11595-011-0335-9.

24. Mankins, W.; Lampman, S. Heat Treatment of Wrought Nickel Alloys. In Heat Treating of Nonferrous Alloys, Totten, G.E., Ed. ASM International: 2016; Vol. 4E, pp. 426-458. DOI: https://doi.org/10.31399/asm.hb.v04e.a0006265.

25. Nickel Alloy, Corrosion and Heat-Resistant, Bars, Forgings, and Rings 52.5 Ni - 19 Cr - 3.0 Mo - 5.1 Cb ( Nb) - 0.90 Ti - 0.50 Al - 18 Fe Consumable Electrode or Vacuum Induction Melted 1775 °F (968 °C) Solution and Precipitation Heat Treated: Sae-Ams; SAE International 2016. DOI: https://doi.org/https://doi.org/10.4271/AMS5663N.

26. Janaki Ram, G.D.; Venugopal Reddy, A.; Prasad Rao, K.; Reddy, G.M.; Sarin Sundar, J.K. Microstructure and Tensile Properties of Inconel 718 Pulsed Nd-Yag Laser Welds. Journal of Materials Processing Technology 2005, 167, 73-82. DOI: https://doi.org/10.1016/j.jmatprotec.2004.09.081.

27. Janaki Ram, G.D.; Venugopal Reddy, A.; Prasad Rao, K.; Madhusudhan Reddy, G. Control of Laves Phase in Inconel 718 Gta Welds with Current Pulsing. Science and Technology of Welding and Joining 2013, 9, 390-398. DOI: https://doi.org/10.1179/136217104225021788.

28. Reed, R.C. The Superalloys- Fundamentals and Applications; Cambridge University Press: Cambridge, 2006; 10.1017/cbo9780511541285. DOI: https://doi.org/10.1017/cbo9780511541285.

29. Aerospace Materials and Material Technologies; 2017; 10.1007/978-981-10-2134-3. DOI: https://doi.org/10.1007/978-981-10-2134-3.

30. Eiselstein, H.L. Age-Hardenable Nickel Alloy. US3046108A, 1962.

40

11. Steen, W.M.; Courtney, C.G.H. Hardfacing of Nimonic 75 Using 2 Kw Continuouswave Co2laser. Metals Technology 2013, 7, 232-237. DOI: https://doi.org/10.1179/030716980803286955.

12. de Oliveira, U.; Ocelík, V.; De Hosson, J.T.M. Analysis of Coaxial Laser Cladding Processing Conditions. Surface and Coatings Technology 2005, 197, 127-136. DOI: https://doi.org/10.1016/j.surfcoat.2004.06.029.

13. Segerstark, A.; Andersson, J.; Svensson, L.-E. Investigation of Laser Metal Deposited Alloy 718 onto an En 1.4401 Stainless Steel Substrate. Optics & Laser Technology 2017, 97, 144-153. DOI: https://doi.org/10.1016/j.optlastec.2017.05.038.

14. Chen, B.; Mazumder, J. Role of Process Parameters During Additive Manufacturing by Direct Metal Deposition of Inconel 718. Rapid Prototyping Journal 2017, 23, 919-929. DOI: https://doi.org/10.1108/rpj-05-2016-0071.

15. Pinkerton, A.J. Advances in the Modeling of Laser Direct Metal Deposition. Journal of Laser Applications 2015, 27. DOI: https://doi.org/10.2351/1.4815992.

16. Zhong, C.; Gasser, A.; Kittel, J.; Schopphoven, T.; Pirch, N.; Fu, J.; Poprawe, R. Study of Process Window Development for High Deposition-Rate Laser Material Deposition by Using Mixed Processing Parameters. Journal of Laser Applications 2015, 27. DOI: https://doi.org/10.2351/1.4919804.

17. Thompson, S.M.; Bian, L.; Shamsaei, N.; Yadollahi, A. An Overview of Direct Laser Deposition for Additive Manufacturing; Part I: Transport Phenomena, Modeling and Diagnostics. Additive Manufacturing 2015, 8, 36-62. DOI: https://doi.org/10.1016/j.addma.2015.07.001.

18. Zhong, C.; Kittel, J.; Gasser, A.; Schleifenbaum, J.H. Study of Nickel-Based Super-Alloys Inconel 718 and Inconel 625 in High-Deposition-Rate Laser Metal Deposition. Optics & Laser Technology 2019, 109, 352-360. DOI: https://doi.org/10.1016/j.optlastec.2018.08.003.

19. Zhong, C.; Gasser, A.; Kittel, J.; Fu, J.; Ding, Y.; Poprawe, R. Microstructures and Tensile Properties of Inconel 718 Formed by High Deposition-Rate Laser Metal Deposition. Journal of Laser Applications 2016, 28. DOI: https://doi.org/10.2351/1.4943290.

20. Dass, A.; Moridi, A. State of the Art in Directed Energy Deposition: From Additive Manufacturing to Materials Design. Coatings 2019, 9. DOI: https://doi.org/10.3390/coatings9070418.

REFERENCES

41

21. Zhang, Y.; Li, Z.; Nie, P.; Wu, Y. Effect of Heat Treatment on Niobium Segregation of Laser-Cladded In718 Alloy Coating. Metallurgical and Materials Transactions A 2012, 44, 708-716. DOI: https://doi.org/10.1007/s11661-012-1459-z.

22. Long, Y.-t.; Nie, P.-l.; Li, Z.-g.; Huang, J.; Li, X.; Xu, X.-m. Segregation of Niobium in Laser Cladding Inconel 718 Superalloy. Transactions of Nonferrous Metals Society of China 2016, 26, 431-436. DOI: https://doi.org/10.1016/s1003-6326(16)64131-6.

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82. Xiao, H.; Li, S.M.; Xiao, W.J.; Li, Y.Q.; Cha, L.M.; Mazumder, J.; Song, L.J. Effects of Laser Modes on Nb Segregation and Laves Phase Formation During Laser Additive Manufacturing of Nickel-Based Superalloy. Materials Letters 2017, 188, 260-262. DOI: https://doi.org/10.1016/j.matlet.2016.10.118.

83. Standard Test Methods for Determining Average Grain Size: Astm E112-13; 2013. DOI: https://doi.org/https://doi.org/10.1520/E0112-13.

84. Astm E384-17, Standard Test Method for Microindentation Hardness of Materials; www.astm.org 2017. DOI: https://doi.org/10.1520/E0384-17.

85. Lia, F.; Park, J.; Tressler, J.; Martukanitz, R. Partitioning of Laser Energy During Directed Energy Deposition. Additive Manufacturing 2017, 18, 31-39. DOI: https://doi.org/10.1016/j.addma.2017.08.012.

86. Parimi, L.L.; A, R.G.; Clark, D.; Attallah, M.M. Microstructural and Texture Development in Direct Laser Fabricated In718. Materials Characterization 2014, 89, 102-111. DOI: https://doi.org/10.1016/j.matchar.2013.12.012.

87. Cherepanov, A.N.; Ovcharenko, V.E. Effect of Nanostructured Composite Powders on the Structure and Strength Properties of the High-Temperature Inconel 718 Alloy. The Physics of Metals and Metallography 2015, 116, 1279-1284. DOI: https://doi.org/10.1134/s0031918x1510004x.

88. Sivaprasad, K.; Ganesh Sundara Raman, S. Influence of Weld Cooling Rate on Microstructure and Mechanical Properties of Alloy 718 Weldments. Metallurgical and Materials Transactions A 2008, 39, 2115-2127. DOI: https://doi.org/10.1007/s11661-008-9553-y.

46

71. Zhang, K.; Liu, W.; Shang, X. Research on the Processing Experiments of Laser Metal Deposition Shaping. Optics & Laser Technology 2007, 39, 549-557. DOI: https://doi.org/10.1016/j.optlastec.2005.10.009.

72. Kistler, N.A.; Nassar, A.R.; Reutzel, E.W.; Corbin, D.J.; Beese, A.M. Effect of Directed Energy Deposition Processing Parameters on Laser Deposited Inconel®718: Microstructure, Fusion Zone Morphology, and Hardness. Journal of Laser Applications 2017, 29. DOI: https://doi.org/10.2351/1.4979702.

73. Güpner, M.; Patschger, A.; Bliedtner, J. Influence of Process Parameters on the Process Efficiency in Laser Metal Deposition Welding. Physics Procedia 2016, 83, 657-666. DOI: https://doi.org/10.1016/j.phpro.2016.08.068.

74. DuPont, J.N. Fundamentals of Weld Solidification. In Welding Fundamentals and Processes, Lienert T., S.T., Babu S., Acoff V., Ed. ASM International: 2011; pp. 96-114.

75. DuPont, J.N.; Lippold, J.C.; Kiser, S.D. Welding Metallurgy and Weldability of Nickel-Base Alloys; 2009.

76. Selcuk, C. Laser Metal Deposition for Powder Metallurgy Parts. Powder Metallurgy 2011, 54, 94-99. DOI: https://doi.org/10.1179/174329011X12977874589924.

77. Segerstark, A.; Andersson, J.; Svensson, L.-E. Influence of Heat Input on Grain Structure in Thin-Wall Deposits Using Laser Metal Powder Deposition. In Proceedings of The 7th International Swedish Production Symposium, SPS16, Lund, Sweden; p. 7.

78. Dinda, G.P.; Dasgupta, A.K.; Mazumder, J. Texture Control During Laser Deposition of Nickel-Based Superalloy. Scripta Materialia 2012, 67, 503-506. DOI: https://doi.org/10.1016/j.scriptamat.2012.06.014.

79. Zhao, X.; Chen, J.; Lin, X.; Huang, W. Study on Microstructure and Mechanical Properties of Laser Rapid Forming Inconel 718. Materials Science and Engineering: A 2008, 478, 119-124. DOI: https://doi.org/10.1016/j.msea.2007.05.079.

80. Zhong, C.; Gasser, A.; Kittel, J.; Wissenbach, K.; Poprawe, R. Improvement of Material Performance of Inconel 718 Formed by High Deposition-Rate Laser Metal Deposition. Materials & Design 2016, 98, 128-134. DOI: https://doi.org/10.1016/j.matdes.2016.03.006.

REFERENCES

47

81. Sui, S.; Zhong, C.; Chen, J.; Gasser, A.; Huang, W.; Schleifenbaum, J.H. Influence of Solution Heat Treatment on Microstructure and Tensile Properties of Inconel 718 Formed by High-Deposition-Rate Laser Metal Deposition. Journal of Alloys and Compounds 2018, 740, 389-399. DOI: https://doi.org/10.1016/j.jallcom.2017.11.004.

82. Xiao, H.; Li, S.M.; Xiao, W.J.; Li, Y.Q.; Cha, L.M.; Mazumder, J.; Song, L.J. Effects of Laser Modes on Nb Segregation and Laves Phase Formation During Laser Additive Manufacturing of Nickel-Based Superalloy. Materials Letters 2017, 188, 260-262. DOI: https://doi.org/10.1016/j.matlet.2016.10.118.

83. Standard Test Methods for Determining Average Grain Size: Astm E112-13; 2013. DOI: https://doi.org/https://doi.org/10.1520/E0112-13.

84. Astm E384-17, Standard Test Method for Microindentation Hardness of Materials; www.astm.org 2017. DOI: https://doi.org/10.1520/E0384-17.

85. Lia, F.; Park, J.; Tressler, J.; Martukanitz, R. Partitioning of Laser Energy During Directed Energy Deposition. Additive Manufacturing 2017, 18, 31-39. DOI: https://doi.org/10.1016/j.addma.2017.08.012.

86. Parimi, L.L.; A, R.G.; Clark, D.; Attallah, M.M. Microstructural and Texture Development in Direct Laser Fabricated In718. Materials Characterization 2014, 89, 102-111. DOI: https://doi.org/10.1016/j.matchar.2013.12.012.

87. Cherepanov, A.N.; Ovcharenko, V.E. Effect of Nanostructured Composite Powders on the Structure and Strength Properties of the High-Temperature Inconel 718 Alloy. The Physics of Metals and Metallography 2015, 116, 1279-1284. DOI: https://doi.org/10.1134/s0031918x1510004x.

88. Sivaprasad, K.; Ganesh Sundara Raman, S. Influence of Weld Cooling Rate on Microstructure and Mechanical Properties of Alloy 718 Weldments. Metallurgical and Materials Transactions A 2008, 39, 2115-2127. DOI: https://doi.org/10.1007/s11661-008-9553-y.

Appended Papers

Appended Papers

Paper A

Influence of Laser-Directed Energy Deposition process parameters and Thermal Post-Treatments on Nb-rich

Secondary Phases of Alloy 718 Specimens

Suhas Sreekanth, Kjell Hurtig, Shrikant Joshi, Joel Andersson

Submitted: “Metals”

MDPI

A

Paper B

Effect of Direct Energy Deposition Process Parameters on Single-Track

Deposits of Alloy 718

Suhas Sreekanth, Ehsan Ghassemali, Kjell Hurtig, Shrikant Joshi, Joel Andersson

Published: “Metals”

MDPI, January 2020

Printed with permission

B

metals

Article

Effect of Direct Energy Deposition Process Parameterson Single-Track Deposits of Alloy 718

Suhas Sreekanth 1,* , Ehsan Ghassemali 2 , Kjell Hurtig 1 , Shrikant Joshi 1 andJoel Andersson 1

1 Department of Engineering Science, University West, 46186 Trollhättan, Sweden; kjell.hurtig@hv.se (K.H.);shrikant.joshi@hv.se (S.J.); joel.andersson@hv.se (J.A.)

2 Department of Materials and Manufacturing, Jönköping University, 55111 Jönköping, Sweden;Ehsan.Ghassemali@ju.se

* Correspondence: suhas.sreekanth@hv.se; Tel.: +46-520-223-484

Received: 5 December 2019; Accepted: 5 January 2020; Published: 7 January 2020�����������������

Abstract: The effect of three important process parameters, namely laser power, scanning speedand laser stand-off distance on the deposit geometry, microstructure and segregation characteristicsin direct energy deposited alloy 718 specimens has been studied. Laser power and laser stand-offdistance were found to notably affect the width and depth of the deposit, while the scanning speedinfluenced the deposit height. An increase in specific energy conditions (between 0.5 J/mm2 and1.0 J/mm2) increased the total area of deposit yielding varied grain morphologies and precipitationbehaviors which were comprehensively analyzed. A deposit comprising three distinct zones, namelythe top, middle and bottom regions, categorized based on the distinct microstructural features formedon account of variation in local solidification conditions. Nb-rich eutectics preferentially segregatedin the top region of the deposit (5.4–9.6% area fraction, Af) which predominantly consisted of anequiaxed grain structure, as compared to the middle (1.5–5.7% Af) and the bottom regions (2.6–4.5%Af), where columnar dendritic morphology was observed. High scan speed was more effective inreducing the area fraction of Nb-rich phases in the top and middle regions of the deposit. The <100>crystallographic direction was observed to be the preferred growth direction of columnar grainswhile equiaxed grains had a random orientation.

Keywords: laser metal deposition (LMD); columnar dendritic morphology; constitutional supercooling;columnar to equiaxed transition (CET); high deposition rate

1. Introduction

Directed Energy Deposition (DED) is an additive manufacturing (AM) process utilized for variedapplications such as rapid prototyping, cladding, building components and part features, in-situalloying, repair and refurbishing applications [1–3]. The process offers advantages such as: (a) thecapability to achieve high deposition volumes in a large build envelope [4]; (b) the ability to depositcompositionally and/or functionally graded materials [5–7]; (c) ability to deposit materials on top offlat and curved substrate surface; and (d) adaptability to existing CNCs and robotic systems, whichprovides an edge over metal powder-bed fusion (PBF) processes. The above benefits offset the relativeinadequacy of the process to produce near-net-shape part features of geometrical complexity that canbe accomplished through PBF processes.

In the present manuscript, DED refers to a variant having laser as the energy source whichuses powder form of Alloy 718 as feedstock. Also, high deposition rate -directed energy deposition(HDR-DED) process is of specific interest in this work as it can have direct implications on processingtimes and economics. Most of the work conducted thus far has involved feed rates lower than 0.5 kg/h,

Metals 2020, 10, 96; doi:10.3390/met10010096 www.mdpi.com/journal/metals

metals

Article

Effect of Direct Energy Deposition Process Parameterson Single-Track Deposits of Alloy 718

Suhas Sreekanth 1,* , Ehsan Ghassemali 2 , Kjell Hurtig 1 , Shrikant Joshi 1 andJoel Andersson 1

1 Department of Engineering Science, University West, 46186 Trollhättan, Sweden; kjell.hurtig@hv.se (K.H.);shrikant.joshi@hv.se (S.J.); joel.andersson@hv.se (J.A.)

2 Department of Materials and Manufacturing, Jönköping University, 55111 Jönköping, Sweden;Ehsan.Ghassemali@ju.se

* Correspondence: suhas.sreekanth@hv.se; Tel.: +46-520-223-484

Received: 5 December 2019; Accepted: 5 January 2020; Published: 7 January 2020�����������������

Abstract: The effect of three important process parameters, namely laser power, scanning speedand laser stand-off distance on the deposit geometry, microstructure and segregation characteristicsin direct energy deposited alloy 718 specimens has been studied. Laser power and laser stand-offdistance were found to notably affect the width and depth of the deposit, while the scanning speedinfluenced the deposit height. An increase in specific energy conditions (between 0.5 J/mm2 and1.0 J/mm2) increased the total area of deposit yielding varied grain morphologies and precipitationbehaviors which were comprehensively analyzed. A deposit comprising three distinct zones, namelythe top, middle and bottom regions, categorized based on the distinct microstructural features formedon account of variation in local solidification conditions. Nb-rich eutectics preferentially segregatedin the top region of the deposit (5.4–9.6% area fraction, Af) which predominantly consisted of anequiaxed grain structure, as compared to the middle (1.5–5.7% Af) and the bottom regions (2.6–4.5%Af), where columnar dendritic morphology was observed. High scan speed was more effective inreducing the area fraction of Nb-rich phases in the top and middle regions of the deposit. The <100>crystallographic direction was observed to be the preferred growth direction of columnar grainswhile equiaxed grains had a random orientation.

Keywords: laser metal deposition (LMD); columnar dendritic morphology; constitutional supercooling;columnar to equiaxed transition (CET); high deposition rate

1. Introduction

Directed Energy Deposition (DED) is an additive manufacturing (AM) process utilized for variedapplications such as rapid prototyping, cladding, building components and part features, in-situalloying, repair and refurbishing applications [1–3]. The process offers advantages such as: (a) thecapability to achieve high deposition volumes in a large build envelope [4]; (b) the ability to depositcompositionally and/or functionally graded materials [5–7]; (c) ability to deposit materials on top offlat and curved substrate surface; and (d) adaptability to existing CNCs and robotic systems, whichprovides an edge over metal powder-bed fusion (PBF) processes. The above benefits offset the relativeinadequacy of the process to produce near-net-shape part features of geometrical complexity that canbe accomplished through PBF processes.

In the present manuscript, DED refers to a variant having laser as the energy source whichuses powder form of Alloy 718 as feedstock. Also, high deposition rate -directed energy deposition(HDR-DED) process is of specific interest in this work as it can have direct implications on processingtimes and economics. Most of the work conducted thus far has involved feed rates lower than 0.5 kg/h,

Metals 2020, 10, 96; doi:10.3390/met10010096 www.mdpi.com/journal/metals

Metals 2020, 10, 96 2 of 16

termed as ‘low deposition rates’ (LDR) [8,9]. This shift from LDR to HDR is expected to influencevarious properties of deposits such as grain structure, morphology, defect content and geometricalcharacteristics. These characteristics of the HDR-DED process was studied by Zhong et al. [10] interms of track porosity, wettability, deposit height and width with varied process parameters such aspower, feed rate and scanning speeds. Porosity decreased with increased laser power whereas widthof the deposit increased with increased power or decreased feed rate. Both height and width of thetracks decreased with increased scanning speed. Although anisotropic properties, porosity contentand elemental segregation properties were reported by Zhong et al. [11] for HDR-DED samples, thetensile properties after Hot Isostatic Pressing (HIP) + heat treatment was comparable to wrought Alloy718. A brief account of microstructure and its mechanical properties were reported for HDR-DED ofAlloy 718 [12], but a more detailed analysis of single-track deposit is necessary for an improved processunderstanding that can be subsequently utilized for 3D-deposition of HDR builds. In the LDR-DEDprocess, it is often reported that increased feed produces increased height and decreased depth of thedeposit, but the effect on width is minimal.

Many studies have considered Alloy 718 as the deposition material in the DED process owing to acombination of factors, including the alloy’s widespread usage particularly in the aerospace industryand the attractive nature of the DED process as a cladding and repair technique. Alloy 718 is used inthe hot section region of gas turbine applications up to a temperature of 650 ◦C [13–15]; however, it isvulnerable to Nb-segregation, an element which influences the precipitation of strengthening phases asoften reported by numerous researchers [16–19]. Various strategies have been employed in controllingand minimizing these Nb-rich eutectics in both welding [20–22] and AM practices. Xiao et al. [23,24]accommodated a strategy that involved variation of laser modes between continuous-wave (CW) andquasi-continuous wave (QCW), wherein the cooling rates achieved by QCW mode were an order ofmagnitude higher compared to CW mode. An equiaxed dendritic microstructure was obtained in QCWmode as opposed to CW mode which produced columnar dendritic microstructures. Chen et al. [25]studied the formation and segregation of laves phase by manipulation of laser energy distributionin the form of flat-top and Gaussian beams. Gaussian beam distribution yielded higher fraction oflaves phase that increased along the build height (6–11% volume fraction) as opposed to flat-topdistribution (5–6% volume fraction) throughout the deposit. However, the cooling rates for Gaussianbeams predicted from dendrite arm spacing were five times higher compared to flat-top beam. Thisresulted in an increased size of precipitates and consequently increased tensile strength in case of theflat-top beam condition.

A thermal model that predicted the cooling rate at varied powers and scanning speeds wasdeveloped by Amine et al. [26] for different scanning strategies. Low power and high scanning speedconditions yielded higher cooling rates. Numerical modeling work by Nie et al. [27] investigated theformation of microstructure in Alloy 718 by prediction of cooling rates along the deposit height for variedpowder sizes and thence the secondary dendritic arm spacing along the deposit centerline. Further, theconcentration of the laves phase was predicted based on the variations in temperature gradient andcooling rates [28]. A low temperature gradient and high cooling rate was predicted to be necessary forformation of a discrete laves phase instead of precipitation of continuous bands along inter-dendriticregions. The effect of process parameters and scanning strategies were experimentally studied by Parimiet al. [29] and the study concluded that grain morphology can alter Nb-rich segregation. A similar workwas reported by Dinda et al. [30] wherein different scanning strategies were employed to control thetexture of DED deposited Alloy 718. Most of the modelling and microstructural work discussed abovepertains to LDR-DED process.

Generally, the multiple tracks in DED have been distinguished into top, middle and bottomregions based on phase fraction analysis and segregation characteristics have been studied along thebuild height [25,31,32]. As a precursor to developing a robust HDR-DED process, it is desirable tounderstand the impact of process parameters on the characteristics of these top, middle and bottomregions of a single track deposit. The present work primarily concerns identification and control of such

Metals 2020, 10, 96 3 of 16

microstructural and segregation patterns in discrete regions of single-track deposits of Alloy 718. Thecontrol of segregation is particularly important in DED repair and building part features applicationsas they are often introduced to direct aging heat treatment after deposition. Variations in phasefraction of segregation can cause heterogeneities in the deposit, rendering ineffective precipitationstrengthening in local areas of segregation, which could prove detrimental to material properties.Strategies to control and nullify such compositional and phase heterogeneities by suitable manipulationof process parameters is an intended outcome of this study. Also, the effect of process parameters onmicrostructure and texture development is considered for single-track deposits.

2. Materials and Methods

The deposited samples were built on an ISEL M40 4-Axis CNC machine bed with COAX-50(Fraunhofer ILT, Aachen, Germany), otherwise referred to as D50-coaxial nozzle. An IPG YLR-6000S, a6 kW Ytterbium (Yb) doped fiber laser (IPG Photonics, Burbach, Germany) was used as the laser sourcecoupled with a Permanova Collimator of 200 mm focal length (Permanova Lasersystems AB, Mölndal,Sweden) with Optoskand fiber of 800 µm diameter (Optoskand AB, Mölndal, Sweden). The powderfeed-rate was controlled through a dual D2-volumetric powder-feeder arrangement from UniquecoatTechnologies (Uniquecoat Technologies LLC, Oilville, VA, United States). Argon was used as both thecarrier gas and shielding gas throughout the deposition process.

Gas Atomized (GA) Alloy 718 powder with a particle size range of 40–105 µm (D10, D50 andD90 of 47 µm, 65 µm and 83 µm respectively) was used as feedstock. Table 1 presents the nominalcompositions of the powder and the substrate complying to the standard F3055-14a [33] and ASTMB670-07 [34] respectively.

Table 1. Chemical composition (wt.%) of Alloy 718 powder and substrate.

Elements Ni Cr Fe Nb + Ta Mo C Ti Al

Powder 52.89 18.7 18.52 4.9 2.94 0.05 0.92 0.61Substrate 53.57 18.7 17.58 4.97 2.89 0.04 0.91 0.59

Elements Co Ta B Cu Mn Si P S

Powder 0.11 <0.01 <0.001 <0.1 0.05 0.19 0.005 0.004Substrate 0.25 0.004 0.002 0.23 0.09 0.06 0.008 0.001

Single tracks were deposited on wrought Alloy 718 substrate with dimensions 150 mm × 60 mm ×3.2 mm. Two tracks of length 55 mm were deposited for each parametric set to check for repeatability.A reduced multi-level factorial design was considered in the present study with three variable factors:power (P), scanning speed (V) and laser stand-off distance (Lo). P was varied at three levels: 1200,1600, 2000 W; V was varied at two levels: 900, 1100 mm/min; and Lo was varied at two levels: 9.5,13 mm. Another parameter set (P = 1800 W; V = 1100 mm/min; Lo = 13 mm) close to high powercondition with was considered for verification of results. Other parameters such as powder feed rate(Mp): 1.2 kg/hr, carrier gas flow (Gc): 5 L/min and shielding gas flow rate (Gs): 12 L/min were keptconstant. The effect of V was studied at intermediate power of 1600 W alone reducing the iterations indesign. The resulting nine parametric sets, as shown in Table 2, were analyzed to study the effect ofP, V and Lo parameters on various deposit characteristics such as deposit geometry, microstructuralmorphology and segregation of niobium-rich (Nb-rich) eutectic phases. Lo of 9.5 mm and 13 mmabove the substrate in the divergence part of laser beam corresponded to a laser spot diameter (Ds) of1.8 mm and 2.3 mm, respectively.

The sample preparation for metallographic investigations involved a series of grinding andpolishing steps performed on a Buehler Eco-Met 300 machine (ITW Test and Measurement GmbH,Esslingen, Germany) followed by electrolytic etching with oxalic acid to reveal the presence of Nb-richeutectic phases that were predicted to form during solidification of Alloy 718. Optical microscopy was

Metals 2020, 10, 96 2 of 16

termed as ‘low deposition rates’ (LDR) [8,9]. This shift from LDR to HDR is expected to influencevarious properties of deposits such as grain structure, morphology, defect content and geometricalcharacteristics. These characteristics of the HDR-DED process was studied by Zhong et al. [10] interms of track porosity, wettability, deposit height and width with varied process parameters such aspower, feed rate and scanning speeds. Porosity decreased with increased laser power whereas widthof the deposit increased with increased power or decreased feed rate. Both height and width of thetracks decreased with increased scanning speed. Although anisotropic properties, porosity contentand elemental segregation properties were reported by Zhong et al. [11] for HDR-DED samples, thetensile properties after Hot Isostatic Pressing (HIP) + heat treatment was comparable to wrought Alloy718. A brief account of microstructure and its mechanical properties were reported for HDR-DED ofAlloy 718 [12], but a more detailed analysis of single-track deposit is necessary for an improved processunderstanding that can be subsequently utilized for 3D-deposition of HDR builds. In the LDR-DEDprocess, it is often reported that increased feed produces increased height and decreased depth of thedeposit, but the effect on width is minimal.

Many studies have considered Alloy 718 as the deposition material in the DED process owing to acombination of factors, including the alloy’s widespread usage particularly in the aerospace industryand the attractive nature of the DED process as a cladding and repair technique. Alloy 718 is used inthe hot section region of gas turbine applications up to a temperature of 650 ◦C [13–15]; however, it isvulnerable to Nb-segregation, an element which influences the precipitation of strengthening phases asoften reported by numerous researchers [16–19]. Various strategies have been employed in controllingand minimizing these Nb-rich eutectics in both welding [20–22] and AM practices. Xiao et al. [23,24]accommodated a strategy that involved variation of laser modes between continuous-wave (CW) andquasi-continuous wave (QCW), wherein the cooling rates achieved by QCW mode were an order ofmagnitude higher compared to CW mode. An equiaxed dendritic microstructure was obtained in QCWmode as opposed to CW mode which produced columnar dendritic microstructures. Chen et al. [25]studied the formation and segregation of laves phase by manipulation of laser energy distributionin the form of flat-top and Gaussian beams. Gaussian beam distribution yielded higher fraction oflaves phase that increased along the build height (6–11% volume fraction) as opposed to flat-topdistribution (5–6% volume fraction) throughout the deposit. However, the cooling rates for Gaussianbeams predicted from dendrite arm spacing were five times higher compared to flat-top beam. Thisresulted in an increased size of precipitates and consequently increased tensile strength in case of theflat-top beam condition.

A thermal model that predicted the cooling rate at varied powers and scanning speeds wasdeveloped by Amine et al. [26] for different scanning strategies. Low power and high scanning speedconditions yielded higher cooling rates. Numerical modeling work by Nie et al. [27] investigated theformation of microstructure in Alloy 718 by prediction of cooling rates along the deposit height for variedpowder sizes and thence the secondary dendritic arm spacing along the deposit centerline. Further, theconcentration of the laves phase was predicted based on the variations in temperature gradient andcooling rates [28]. A low temperature gradient and high cooling rate was predicted to be necessary forformation of a discrete laves phase instead of precipitation of continuous bands along inter-dendriticregions. The effect of process parameters and scanning strategies were experimentally studied by Parimiet al. [29] and the study concluded that grain morphology can alter Nb-rich segregation. A similar workwas reported by Dinda et al. [30] wherein different scanning strategies were employed to control thetexture of DED deposited Alloy 718. Most of the modelling and microstructural work discussed abovepertains to LDR-DED process.

Generally, the multiple tracks in DED have been distinguished into top, middle and bottomregions based on phase fraction analysis and segregation characteristics have been studied along thebuild height [25,31,32]. As a precursor to developing a robust HDR-DED process, it is desirable tounderstand the impact of process parameters on the characteristics of these top, middle and bottomregions of a single track deposit. The present work primarily concerns identification and control of such

Metals 2020, 10, 96 3 of 16

microstructural and segregation patterns in discrete regions of single-track deposits of Alloy 718. Thecontrol of segregation is particularly important in DED repair and building part features applicationsas they are often introduced to direct aging heat treatment after deposition. Variations in phasefraction of segregation can cause heterogeneities in the deposit, rendering ineffective precipitationstrengthening in local areas of segregation, which could prove detrimental to material properties.Strategies to control and nullify such compositional and phase heterogeneities by suitable manipulationof process parameters is an intended outcome of this study. Also, the effect of process parameters onmicrostructure and texture development is considered for single-track deposits.

2. Materials and Methods

The deposited samples were built on an ISEL M40 4-Axis CNC machine bed with COAX-50(Fraunhofer ILT, Aachen, Germany), otherwise referred to as D50-coaxial nozzle. An IPG YLR-6000S, a6 kW Ytterbium (Yb) doped fiber laser (IPG Photonics, Burbach, Germany) was used as the laser sourcecoupled with a Permanova Collimator of 200 mm focal length (Permanova Lasersystems AB, Mölndal,Sweden) with Optoskand fiber of 800 µm diameter (Optoskand AB, Mölndal, Sweden). The powderfeed-rate was controlled through a dual D2-volumetric powder-feeder arrangement from UniquecoatTechnologies (Uniquecoat Technologies LLC, Oilville, VA, United States). Argon was used as both thecarrier gas and shielding gas throughout the deposition process.

Gas Atomized (GA) Alloy 718 powder with a particle size range of 40–105 µm (D10, D50 andD90 of 47 µm, 65 µm and 83 µm respectively) was used as feedstock. Table 1 presents the nominalcompositions of the powder and the substrate complying to the standard F3055-14a [33] and ASTMB670-07 [34] respectively.

Table 1. Chemical composition (wt.%) of Alloy 718 powder and substrate.

Elements Ni Cr Fe Nb + Ta Mo C Ti Al

Powder 52.89 18.7 18.52 4.9 2.94 0.05 0.92 0.61Substrate 53.57 18.7 17.58 4.97 2.89 0.04 0.91 0.59

Elements Co Ta B Cu Mn Si P S

Powder 0.11 <0.01 <0.001 <0.1 0.05 0.19 0.005 0.004Substrate 0.25 0.004 0.002 0.23 0.09 0.06 0.008 0.001

Single tracks were deposited on wrought Alloy 718 substrate with dimensions 150 mm × 60 mm ×3.2 mm. Two tracks of length 55 mm were deposited for each parametric set to check for repeatability.A reduced multi-level factorial design was considered in the present study with three variable factors:power (P), scanning speed (V) and laser stand-off distance (Lo). P was varied at three levels: 1200,1600, 2000 W; V was varied at two levels: 900, 1100 mm/min; and Lo was varied at two levels: 9.5,13 mm. Another parameter set (P = 1800 W; V = 1100 mm/min; Lo = 13 mm) close to high powercondition with was considered for verification of results. Other parameters such as powder feed rate(Mp): 1.2 kg/hr, carrier gas flow (Gc): 5 L/min and shielding gas flow rate (Gs): 12 L/min were keptconstant. The effect of V was studied at intermediate power of 1600 W alone reducing the iterations indesign. The resulting nine parametric sets, as shown in Table 2, were analyzed to study the effect ofP, V and Lo parameters on various deposit characteristics such as deposit geometry, microstructuralmorphology and segregation of niobium-rich (Nb-rich) eutectic phases. Lo of 9.5 mm and 13 mmabove the substrate in the divergence part of laser beam corresponded to a laser spot diameter (Ds) of1.8 mm and 2.3 mm, respectively.

The sample preparation for metallographic investigations involved a series of grinding andpolishing steps performed on a Buehler Eco-Met 300 machine (ITW Test and Measurement GmbH,Esslingen, Germany) followed by electrolytic etching with oxalic acid to reveal the presence of Nb-richeutectic phases that were predicted to form during solidification of Alloy 718. Optical microscopy was

Metals 2020, 10, 96 4 of 16

carried out with the aid of a Zeiss Axioscope 5 coupled with Axiocam 506 (Carl Zeiss AB, Stockholm,Sweden). For higher magnifications needed for phase fraction analyses and phase identification,a Hitachi TM3000 table-top SEM (Hitachi High-Technologies Europe GmbH, Stockholm, Sweden),coupled with a Bruker Scan Generator and an Oxford Instruments X-Max 20 Electron DiffractionSpectroscopy (EDS) detector (Oxford Instruments NanoAnalysis, Stockholm, Sweden) were utilized.Electron Backscatter Diffraction (EBSD) analysis was performed with JEOL7001F SEM/EDAX-Hikari(super) detector (JEOL Nordic AB, Sollentuna, Sweden) combination. TEAM 4.5 and OIM (TSL) version8 (Ametek GmbH, Weiterstadt, Germany) were used for data acquisition and analyses.

Table 2. Parameter sets used in the designed experiments.

Laser Power,P (W)

Scanning Speed, V(mm/min)

Laser Stand-OffDistance, Lo (mm) Designation

1200 1100 9.5 11600 900 9.5 22000 1100 9.5 31200 1100 13 41600 900 13 52000 1100 13 61800 1100 13 71600 1100 9.5 81600 1100 13 9

A Gaussian power distribution was obtained for the parametric conditions considered (Figure 1a).A typical beam profile and power distribution is shown in Figure 1 for P = 1200 W and Lo = 9.5 mm.The power intensity at Lo = 9.5 mm was higher compared to Lo = 13 mm for all laser powers (1200,1600 and 2000 W) as recorded in Table 3.

Metals 2020, 10, x FOR PEER REVIEW 4 of 16

The sample preparation for metallographic investigations involved a series of grinding and polishing steps performed on a Buehler Eco-Met 300 machine (ITW Test and Measurement GmbH, Esslingen, Germany) followed by electrolytic etching with oxalic acid to reveal the presence of Nb-rich eutectic phases that were predicted to form during solidification of Alloy 718. Optical microscopy was carried out with the aid of a Zeiss Axioscope 5 coupled with Axiocam 506 (Carl Zeiss AB, Stockholm, Sweden). For higher magnifications needed for phase fraction analyses and phase identification, a Hitachi TM3000 table-top SEM (Hitachi High-Technologies Europe GmbH, Stockholm, Sweden), coupled with a Bruker Scan Generator and an Oxford Instruments X-Max 20 Electron Diffraction Spectroscopy (EDS) detector (Oxford Instruments NanoAnalysis, Stockholm, Sweden) were utilized. Electron Backscatter Diffraction (EBSD) analysis was performed with JEOL7001F SEM/EDAX-Hikari (super) detector (JEOL Nordic AB, Sollentuna, Sweden) combination. TEAM 4.5 and OIM (TSL) version 8 (Ametek GmbH, Weiterstadt, Germany) were used for data acquisition and analyses.

A Gaussian power distribution was obtained for the parametric conditions considered (Figure 1a). A typical beam profile and power distribution is shown in Figure 1 for P = 1200 W and Lo = 9.5 mm. The power intensity at Lo = 9.5 mm was higher compared to Lo = 13 mm for all laser powers (1200, 1600 and 2000 W) as recorded in Table 3.

(a) (b)

Figure 1. (a) Gaussian distribution of power intensity; (b) laser beam profile for P of 1200 W and Lo of 9.5 mm above the focal point.

Table 3. Peak power intensity for different Lo at varying P.

Laser Power P in W Peak Power Intensity

Lo = 9.5 mm Lo = 13 mm Peak Power, kW/cm2 Peak Power, kW/cm2

1200 85.6 52.6 1600 114 70 2000 142.5 87.7

A typical cross-section of a single-track DED deposit depicting geometric measurements of interest is illustrated in Figure 2. The DED material showed a distinct appearance compared to the wrought microstructure present in the substrate region. The horizontal dotted line marked in the figure is indicative of the original substrate surface before deposition. The values of height (H), width (W) and depth (D) were obtained by measuring dimensions as in Figure 2. The measurement of area of deposition and dilution was done using Image-J software (version 1.5). The area above the fine dotted line indicated the area of deposit (Ad) and the area below indicated area of dilution (As), measured in mm2. Dilution percentage was obtained by the equation;

Figure 1. (a) Gaussian distribution of power intensity; (b) laser beam profile for P of 1200 W and Lo of9.5 mm above the focal point.

Table 3. Peak power intensity for different Lo at varying P.

Laser Power P in WPeak Power Intensity

Lo = 9.5 mm Lo = 13 mm

Peak Power, kW/cm2 Peak Power, kW/cm2

1200 85.6 52.61600 114 702000 142.5 87.7

Metals 2020, 10, 96 5 of 16

A typical cross-section of a single-track DED deposit depicting geometric measurements of interestis illustrated in Figure 2. The DED material showed a distinct appearance compared to the wroughtmicrostructure present in the substrate region. The horizontal dotted line marked in the figure isindicative of the original substrate surface before deposition. The values of height (H), width (W)and depth (D) were obtained by measuring dimensions as in Figure 2. The measurement of area ofdeposition and dilution was done using Image-J software (version 1.5). The area above the fine dottedline indicated the area of deposit (Ad) and the area below indicated area of dilution (As), measured inmm2. Dilution percentage was obtained by the equation;

%Dilution =As

As + Ad(1)

Metals 2020, 10, x FOR PEER REVIEW 5 of 16

%�i���i�n � ���� � �� (1)

The powder capture efficiency (η) can provide useful insights into the different geometrical aspects of deposits, which can be calculated by the equation [10];

� � ρ � � � ��Mp (2)

where ρ represents the density of Alloy 718. As Mp in the present set of experiments was kept constant, a specific energy term (measured in

J/mm2) can be determined by the following relationship between process parameters [31];

Specific Energy � P� � �� (3)

Figure 2. A typical directed energy deposition (DED) deposit depicting geometric measurements of interest.

Quantification of Nb-rich phases was performed using Image-J software. The threshold value for brightness/contrast ratios were assigned based on the SEM images. Further, a binary mask was applied for quantification of results.

The determination of non-equiaxed grain-size was performed as per the ASTM E112-13 standard [35]. The mean length of intercept (𝑙𝑙) was calculated in three co-ordinate axes along the geometry of the deposit to measure the average grain size. ASTM number (G) was calculated by using the relation;

G = (−6.644 log10 𝑙𝑙) − 3.288 (4)

3. Results and Discussions

All the parameter sets considered in this experimental design (see Table 2) yielded sound and continuous deposition of material. In the following section geometrical characteristics, morphological characteristics and segregation observed at varied processing conditions are reported and discussed.

3.1. Deposit Geometry

A desirable attribute of AM is to minimize material removal through machining processes and hence calls for accurate deposit geometries, attainable by controlling process parameters. Geometric measurements were compiled by averaging six different sections. The values of H varied between 900 µm and 1300 µm, W varied between 1900 µm and 2600 µm and D varied between 120 µm and 720 µm. It can be noted from Figure 3a, that the total area (As + Ad) and Ad increased with increase in specific energy. Figure 3b shows the variation in η at Lo = 9.5 mm and 13 mm. It shows an increasing trend with a product of specific energy and laser spot diameter. As the stand-off distance increased, the area of heat input increased due to larger spot diameter, thereby melting more powder.

Figure 2. A typical directed energy deposition (DED) deposit depicting geometric measurementsof interest.

The powder capture efficiency (η) can provide useful insights into the different geometrical aspectsof deposits, which can be calculated by the equation [10];

η =ρ×V×Ad

Mp(2)

where ρ represents the density of Alloy 718.As Mp in the present set of experiments was kept constant, a specific energy term (measured in

J/mm2) can be determined by the following relationship between process parameters [31];

Specific Energy =P

V×Ds(3)

Quantification of Nb-rich phases was performed using Image-J software. The threshold valuefor brightness/contrast ratios were assigned based on the SEM images. Further, a binary mask wasapplied for quantification of results.

The determination of non-equiaxed grain-size was performed as per the ASTM E112-13 standard [35].The mean length of intercept (l) was calculated in three co-ordinate axes along the geometry of thedeposit to measure the average grain size. ASTM number (G) was calculated by using the relation;

G =(−6.644 log10 l

)− 3.288 (4)

3. Results and Discussions

All the parameter sets considered in this experimental design (see Table 2) yielded sound andcontinuous deposition of material. In the following section geometrical characteristics, morphologicalcharacteristics and segregation observed at varied processing conditions are reported and discussed.

Metals 2020, 10, 96 4 of 16

carried out with the aid of a Zeiss Axioscope 5 coupled with Axiocam 506 (Carl Zeiss AB, Stockholm,Sweden). For higher magnifications needed for phase fraction analyses and phase identification,a Hitachi TM3000 table-top SEM (Hitachi High-Technologies Europe GmbH, Stockholm, Sweden),coupled with a Bruker Scan Generator and an Oxford Instruments X-Max 20 Electron DiffractionSpectroscopy (EDS) detector (Oxford Instruments NanoAnalysis, Stockholm, Sweden) were utilized.Electron Backscatter Diffraction (EBSD) analysis was performed with JEOL7001F SEM/EDAX-Hikari(super) detector (JEOL Nordic AB, Sollentuna, Sweden) combination. TEAM 4.5 and OIM (TSL) version8 (Ametek GmbH, Weiterstadt, Germany) were used for data acquisition and analyses.

Table 2. Parameter sets used in the designed experiments.

Laser Power,P (W)

Scanning Speed, V(mm/min)

Laser Stand-OffDistance, Lo (mm) Designation

1200 1100 9.5 11600 900 9.5 22000 1100 9.5 31200 1100 13 41600 900 13 52000 1100 13 61800 1100 13 71600 1100 9.5 81600 1100 13 9

A Gaussian power distribution was obtained for the parametric conditions considered (Figure 1a).A typical beam profile and power distribution is shown in Figure 1 for P = 1200 W and Lo = 9.5 mm.The power intensity at Lo = 9.5 mm was higher compared to Lo = 13 mm for all laser powers (1200,1600 and 2000 W) as recorded in Table 3.

Metals 2020, 10, x FOR PEER REVIEW 4 of 16

The sample preparation for metallographic investigations involved a series of grinding and polishing steps performed on a Buehler Eco-Met 300 machine (ITW Test and Measurement GmbH, Esslingen, Germany) followed by electrolytic etching with oxalic acid to reveal the presence of Nb-rich eutectic phases that were predicted to form during solidification of Alloy 718. Optical microscopy was carried out with the aid of a Zeiss Axioscope 5 coupled with Axiocam 506 (Carl Zeiss AB, Stockholm, Sweden). For higher magnifications needed for phase fraction analyses and phase identification, a Hitachi TM3000 table-top SEM (Hitachi High-Technologies Europe GmbH, Stockholm, Sweden), coupled with a Bruker Scan Generator and an Oxford Instruments X-Max 20 Electron Diffraction Spectroscopy (EDS) detector (Oxford Instruments NanoAnalysis, Stockholm, Sweden) were utilized. Electron Backscatter Diffraction (EBSD) analysis was performed with JEOL7001F SEM/EDAX-Hikari (super) detector (JEOL Nordic AB, Sollentuna, Sweden) combination. TEAM 4.5 and OIM (TSL) version 8 (Ametek GmbH, Weiterstadt, Germany) were used for data acquisition and analyses.

A Gaussian power distribution was obtained for the parametric conditions considered (Figure 1a). A typical beam profile and power distribution is shown in Figure 1 for P = 1200 W and Lo = 9.5 mm. The power intensity at Lo = 9.5 mm was higher compared to Lo = 13 mm for all laser powers (1200, 1600 and 2000 W) as recorded in Table 3.

(a) (b)

Figure 1. (a) Gaussian distribution of power intensity; (b) laser beam profile for P of 1200 W and Lo of 9.5 mm above the focal point.

Table 3. Peak power intensity for different Lo at varying P.

Laser Power P in W Peak Power Intensity

Lo = 9.5 mm Lo = 13 mm Peak Power, kW/cm2 Peak Power, kW/cm2

1200 85.6 52.6 1600 114 70 2000 142.5 87.7

A typical cross-section of a single-track DED deposit depicting geometric measurements of interest is illustrated in Figure 2. The DED material showed a distinct appearance compared to the wrought microstructure present in the substrate region. The horizontal dotted line marked in the figure is indicative of the original substrate surface before deposition. The values of height (H), width (W) and depth (D) were obtained by measuring dimensions as in Figure 2. The measurement of area of deposition and dilution was done using Image-J software (version 1.5). The area above the fine dotted line indicated the area of deposit (Ad) and the area below indicated area of dilution (As), measured in mm2. Dilution percentage was obtained by the equation;

Figure 1. (a) Gaussian distribution of power intensity; (b) laser beam profile for P of 1200 W and Lo of9.5 mm above the focal point.

Table 3. Peak power intensity for different Lo at varying P.

Laser Power P in WPeak Power Intensity

Lo = 9.5 mm Lo = 13 mm

Peak Power, kW/cm2 Peak Power, kW/cm2

1200 85.6 52.61600 114 702000 142.5 87.7

Metals 2020, 10, 96 5 of 16

A typical cross-section of a single-track DED deposit depicting geometric measurements of interestis illustrated in Figure 2. The DED material showed a distinct appearance compared to the wroughtmicrostructure present in the substrate region. The horizontal dotted line marked in the figure isindicative of the original substrate surface before deposition. The values of height (H), width (W)and depth (D) were obtained by measuring dimensions as in Figure 2. The measurement of area ofdeposition and dilution was done using Image-J software (version 1.5). The area above the fine dottedline indicated the area of deposit (Ad) and the area below indicated area of dilution (As), measured inmm2. Dilution percentage was obtained by the equation;

%Dilution =As

As + Ad(1)

Metals 2020, 10, x FOR PEER REVIEW 5 of 16

%�i���i�n � ���� � �� (1)

The powder capture efficiency (η) can provide useful insights into the different geometrical aspects of deposits, which can be calculated by the equation [10];

� � ρ � � � ��Mp (2)

where ρ represents the density of Alloy 718. As Mp in the present set of experiments was kept constant, a specific energy term (measured in

J/mm2) can be determined by the following relationship between process parameters [31];

Specific Energy � P� � �� (3)

Figure 2. A typical directed energy deposition (DED) deposit depicting geometric measurements of interest.

Quantification of Nb-rich phases was performed using Image-J software. The threshold value for brightness/contrast ratios were assigned based on the SEM images. Further, a binary mask was applied for quantification of results.

The determination of non-equiaxed grain-size was performed as per the ASTM E112-13 standard [35]. The mean length of intercept (𝑙𝑙) was calculated in three co-ordinate axes along the geometry of the deposit to measure the average grain size. ASTM number (G) was calculated by using the relation;

G = (−6.644 log10 𝑙𝑙) − 3.288 (4)

3. Results and Discussions

All the parameter sets considered in this experimental design (see Table 2) yielded sound and continuous deposition of material. In the following section geometrical characteristics, morphological characteristics and segregation observed at varied processing conditions are reported and discussed.

3.1. Deposit Geometry

A desirable attribute of AM is to minimize material removal through machining processes and hence calls for accurate deposit geometries, attainable by controlling process parameters. Geometric measurements were compiled by averaging six different sections. The values of H varied between 900 µm and 1300 µm, W varied between 1900 µm and 2600 µm and D varied between 120 µm and 720 µm. It can be noted from Figure 3a, that the total area (As + Ad) and Ad increased with increase in specific energy. Figure 3b shows the variation in η at Lo = 9.5 mm and 13 mm. It shows an increasing trend with a product of specific energy and laser spot diameter. As the stand-off distance increased, the area of heat input increased due to larger spot diameter, thereby melting more powder.

Figure 2. A typical directed energy deposition (DED) deposit depicting geometric measurementsof interest.

The powder capture efficiency (η) can provide useful insights into the different geometrical aspectsof deposits, which can be calculated by the equation [10];

η =ρ×V×Ad

Mp(2)

where ρ represents the density of Alloy 718.As Mp in the present set of experiments was kept constant, a specific energy term (measured in

J/mm2) can be determined by the following relationship between process parameters [31];

Specific Energy =P

V×Ds(3)

Quantification of Nb-rich phases was performed using Image-J software. The threshold valuefor brightness/contrast ratios were assigned based on the SEM images. Further, a binary mask wasapplied for quantification of results.

The determination of non-equiaxed grain-size was performed as per the ASTM E112-13 standard [35].The mean length of intercept (l) was calculated in three co-ordinate axes along the geometry of thedeposit to measure the average grain size. ASTM number (G) was calculated by using the relation;

G =(−6.644 log10 l

)− 3.288 (4)

3. Results and Discussions

All the parameter sets considered in this experimental design (see Table 2) yielded sound andcontinuous deposition of material. In the following section geometrical characteristics, morphologicalcharacteristics and segregation observed at varied processing conditions are reported and discussed.

Metals 2020, 10, 96 6 of 16

3.1. Deposit Geometry

A desirable attribute of AM is to minimize material removal through machining processes andhence calls for accurate deposit geometries, attainable by controlling process parameters. Geometricmeasurements were compiled by averaging six different sections. The values of H varied between 900µm and 1300 µm, W varied between 1900 µm and 2600 µm and D varied between 120 µm and 720 µm.It can be noted from Figure 3a, that the total area (As + Ad) and Ad increased with increase in specificenergy. Figure 3b shows the variation in η at Lo = 9.5 mm and 13 mm. It shows an increasing trendwith a product of specific energy and laser spot diameter. As the stand-off distance increased, the areaof heat input increased due to larger spot diameter, thereby melting more powder.

Metals 2020, 10, x FOR PEER REVIEW 6 of 16

Energy from the laser source is primarily utilized for melting the feedstock powder continuously fed into a melt-pool in a DED process. However, a part of the energy is also utilized for melting the substrate to create the melt-pool which determines the amount of dilution. When the feed rate is low, a higher amount of energy is available for melting the substrate yielding high dilutions, given other parametric conditions remain the same [36]. The values of dilution measured are depicted in Figure 3a for various process parameters and found to vary between 9% and 39%.

(a) (b)

Figure 3. (a) Plot of energy input vs. area of deposit (Ad) and area of dilution (As); (b) plot of % powder capture efficiency vs specific energy x laser spot diameter. Note: numbers adjacent to data-points in (a,b) correspond to parameter set numbers summarized in Table 2.

3.1.1. Effect of Laser Power and Laser Stand-Off Distance on the Geometry of Deposits

Three different laser powers: 1200, 1600 and 2000 W were employed at two different laser stand-off distances of 9.5 mm and 13 mm. An additional parameter of P = 1800 W with Lo = 13 mm was also considered as shown in parameter set 7. Figure 4 shows the effect of P on H, W and D of deposits having processing conditions as tabulated in Table 2.

H decreased when P increased from 1200 W to 1600 W and increased when P was further increased to 2000 W. 𝜂𝜂 depicted in Figure 3c were 66% and 72% for Lo of 9.5 mm and 13 mm respectively at P of 1600 W (decreased H conditions), which is quite low, hence the decrease in H measured. Decreased 𝜂𝜂𝜂decreased Mp, which in turn decreased the value of H; as H is a variable dependent on Mp. Similar results were obtained by Chen et al. [31], wherein the effect of P on H was not ascertained. A study by Segerstark et al. [32] concluded that decrease in V and increase in Mp

increased H, but P was found to have negligible effect on H. A study by Corbin et al. [37] showed that H remained fairly constant over a certain P (250–400 W) and gradually increased with further increase in P when other process parameters remained unchanged, which ascertains the result obtained in the present study. Change in Lo did not show any significant effect on H as seen in Figure 4a.

W increased with P in both cases of Lo as shown in Figure 4b. A relatively large amount of energy was available at 2000 W compared to 1200 W and 1600 W, increasing the shape and size of the melt-pool resulting in increased W. The power distribution in case of Lo = 13 mm acts upon a larger area due to Ds (2.3 mm) compared to Lo = 9.5 mm (Ds = 1.8 mm) and thereby increases W. The powder capture efficiency has minimal effects on W, and same is seen in case of P = 1600 W as Ds was the dominant factor. As reported in the introduction, decreased Mp leads to increased W, as more lateral heat transfer takes place for melting the substrate and providing a larger area for the melt flow.

Figure 4c indicates that Lo of 9.5 mm had higher D compared to the Lo of 13 mm due to higher energy density at 9.5 mm (Table 3) which melted more material. The specific energy had a noticeable effect on D, as deposits 2 and 3 indicated higher dilution and in the other extreme, deposits 1 and 4 showed lower levels of dilution (Figure 3a).

Specific Energy, J/mm2

Mea

sure

d A

rea,

mm

2

Figure 3. (a) Plot of energy input vs. area of deposit (Ad) and area of dilution (As); (b) plot of % powdercapture efficiency vs specific energy x laser spot diameter. Note: numbers adjacent to data-points in(a,b) correspond to parameter set numbers summarized in Table 2.

Energy from the laser source is primarily utilized for melting the feedstock powder continuouslyfed into a melt-pool in a DED process. However, a part of the energy is also utilized for melting thesubstrate to create the melt-pool which determines the amount of dilution. When the feed rate is low, ahigher amount of energy is available for melting the substrate yielding high dilutions, given otherparametric conditions remain the same [36]. The values of dilution measured are depicted in Figure 3afor various process parameters and found to vary between 9% and 39%.

3.1.1. Effect of Laser Power and Laser Stand-Off Distance on the Geometry of Deposits

Three different laser powers: 1200, 1600 and 2000 W were employed at two different laser stand-offdistances of 9.5 mm and 13 mm. An additional parameter of P = 1800 W with Lo = 13 mm was alsoconsidered as shown in parameter set 7. Figure 4 shows the effect of P on H, W and D of depositshaving processing conditions as tabulated in Table 2.

H decreased when P increased from 1200 W to 1600 W and increased when P was further increasedto 2000 W. η depicted in Figure 3c were 66% and 72% for Lo of 9.5 mm and 13 mm respectively at P of1600 W (decreased H conditions), which is quite low, hence the decrease in H measured. Decreased ηdecreased Mp, which in turn decreased the value of H; as H is a variable dependent on Mp. Similarresults were obtained by Chen et al. [31], wherein the effect of P on H was not ascertained. A study bySegerstark et al. [32] concluded that decrease in V and increase in Mp increased H, but P was foundto have negligible effect on H. A study by Corbin et al. [37] showed that H remained fairly constantover a certain P (250–400 W) and gradually increased with further increase in P when other processparameters remained unchanged, which ascertains the result obtained in the present study. Change inLo did not show any significant effect on H as seen in Figure 4a.

W increased with P in both cases of Lo as shown in Figure 4b. A relatively large amount of energywas available at 2000 W compared to 1200 W and 1600 W, increasing the shape and size of the melt-poolresulting in increased W. The power distribution in case of Lo = 13 mm acts upon a larger area due to

Metals 2020, 10, 96 7 of 16

Ds (2.3 mm) compared to Lo = 9.5 mm (Ds = 1.8 mm) and thereby increases W. The powder captureefficiency has minimal effects on W, and same is seen in case of P = 1600 W as Ds was the dominantfactor. As reported in the introduction, decreased Mp leads to increased W, as more lateral heat transfertakes place for melting the substrate and providing a larger area for the melt flow.

Figure 4c indicates that Lo of 9.5 mm had higher D compared to the Lo of 13 mm due to higherenergy density at 9.5 mm (Table 3) which melted more material. The specific energy had a noticeableeffect on D, as deposits 2 and 3 indicated higher dilution and in the other extreme, deposits 1 and 4showed lower levels of dilution (Figure 3a).Metals 2020, 10, x FOR PEER REVIEW 7 of 16

(a)

(b)

(c)

Figure 4. Influence of laser power of 1200, 1600, 1800 and 2000 W on (a) height; (b) width; and (c) depth of deposits at laser stand-off distances of 9.5 mm and 13 mm at a constant scanning speed of 1100 mm/min.

3.1.2. Effect of Scanning Speed and Laser Stand-Off Distance on the Geometry of Deposits

For studying the effects of scanning speed and laser stand-off distance on geometry, single-tracks processed at two different V of 900 mm/min and 1100 mm/min with Lo of 9.5 mm and 13 mm at P of 1600 W were considered, as in parameter sets 2, 5, 8 and 9.

It can be seen in Figure 5a that the H was insensitive to the changes in Lo. However, with increased V, H decreased. W showed a similar trend as H with respect to V. An increased V showed a decline in W, although the decline was less significant as compared to H. The effect of Lo on W was significant as seen in Figure 5b. A lower Lo of 9.5 mm yielded lower W compared to Lo of 13 mm. D showed a varied relationship as compared to H and W depicted in Figure 5c.

Figure 4. Influence of laser power of 1200, 1600, 1800 and 2000 W on (a) height; (b) width; and (c) depth ofdeposits at laser stand-off distances of 9.5 mm and 13 mm at a constant scanning speed of 1100 mm/min.

3.1.2. Effect of Scanning Speed and Laser Stand-Off Distance on the Geometry of Deposits

For studying the effects of scanning speed and laser stand-off distance on geometry, single-tracksprocessed at two different V of 900 mm/min and 1100 mm/min with Lo of 9.5 mm and 13 mm at P of1600 W were considered, as in parameter sets 2, 5, 8 and 9.

It can be seen in Figure 5a that the H was insensitive to the changes in Lo. However, with increasedV, H decreased. W showed a similar trend as H with respect to V. An increased V showed a decline in

Metals 2020, 10, 96 6 of 16

3.1. Deposit Geometry

A desirable attribute of AM is to minimize material removal through machining processes andhence calls for accurate deposit geometries, attainable by controlling process parameters. Geometricmeasurements were compiled by averaging six different sections. The values of H varied between 900µm and 1300 µm, W varied between 1900 µm and 2600 µm and D varied between 120 µm and 720 µm.It can be noted from Figure 3a, that the total area (As + Ad) and Ad increased with increase in specificenergy. Figure 3b shows the variation in η at Lo = 9.5 mm and 13 mm. It shows an increasing trendwith a product of specific energy and laser spot diameter. As the stand-off distance increased, the areaof heat input increased due to larger spot diameter, thereby melting more powder.

Metals 2020, 10, x FOR PEER REVIEW 6 of 16

Energy from the laser source is primarily utilized for melting the feedstock powder continuously fed into a melt-pool in a DED process. However, a part of the energy is also utilized for melting the substrate to create the melt-pool which determines the amount of dilution. When the feed rate is low, a higher amount of energy is available for melting the substrate yielding high dilutions, given other parametric conditions remain the same [36]. The values of dilution measured are depicted in Figure 3a for various process parameters and found to vary between 9% and 39%.

(a) (b)

Figure 3. (a) Plot of energy input vs. area of deposit (Ad) and area of dilution (As); (b) plot of % powder capture efficiency vs specific energy x laser spot diameter. Note: numbers adjacent to data-points in (a,b) correspond to parameter set numbers summarized in Table 2.

3.1.1. Effect of Laser Power and Laser Stand-Off Distance on the Geometry of Deposits

Three different laser powers: 1200, 1600 and 2000 W were employed at two different laser stand-off distances of 9.5 mm and 13 mm. An additional parameter of P = 1800 W with Lo = 13 mm was also considered as shown in parameter set 7. Figure 4 shows the effect of P on H, W and D of deposits having processing conditions as tabulated in Table 2.

H decreased when P increased from 1200 W to 1600 W and increased when P was further increased to 2000 W. 𝜂𝜂 depicted in Figure 3c were 66% and 72% for Lo of 9.5 mm and 13 mm respectively at P of 1600 W (decreased H conditions), which is quite low, hence the decrease in H measured. Decreased 𝜂𝜂𝜂decreased Mp, which in turn decreased the value of H; as H is a variable dependent on Mp. Similar results were obtained by Chen et al. [31], wherein the effect of P on H was not ascertained. A study by Segerstark et al. [32] concluded that decrease in V and increase in Mp

increased H, but P was found to have negligible effect on H. A study by Corbin et al. [37] showed that H remained fairly constant over a certain P (250–400 W) and gradually increased with further increase in P when other process parameters remained unchanged, which ascertains the result obtained in the present study. Change in Lo did not show any significant effect on H as seen in Figure 4a.

W increased with P in both cases of Lo as shown in Figure 4b. A relatively large amount of energy was available at 2000 W compared to 1200 W and 1600 W, increasing the shape and size of the melt-pool resulting in increased W. The power distribution in case of Lo = 13 mm acts upon a larger area due to Ds (2.3 mm) compared to Lo = 9.5 mm (Ds = 1.8 mm) and thereby increases W. The powder capture efficiency has minimal effects on W, and same is seen in case of P = 1600 W as Ds was the dominant factor. As reported in the introduction, decreased Mp leads to increased W, as more lateral heat transfer takes place for melting the substrate and providing a larger area for the melt flow.

Figure 4c indicates that Lo of 9.5 mm had higher D compared to the Lo of 13 mm due to higher energy density at 9.5 mm (Table 3) which melted more material. The specific energy had a noticeable effect on D, as deposits 2 and 3 indicated higher dilution and in the other extreme, deposits 1 and 4 showed lower levels of dilution (Figure 3a).

Specific Energy, J/mm2

Mea

sure

d A

rea,

mm

2

Figure 3. (a) Plot of energy input vs. area of deposit (Ad) and area of dilution (As); (b) plot of % powdercapture efficiency vs specific energy x laser spot diameter. Note: numbers adjacent to data-points in(a,b) correspond to parameter set numbers summarized in Table 2.

Energy from the laser source is primarily utilized for melting the feedstock powder continuouslyfed into a melt-pool in a DED process. However, a part of the energy is also utilized for melting thesubstrate to create the melt-pool which determines the amount of dilution. When the feed rate is low, ahigher amount of energy is available for melting the substrate yielding high dilutions, given otherparametric conditions remain the same [36]. The values of dilution measured are depicted in Figure 3afor various process parameters and found to vary between 9% and 39%.

3.1.1. Effect of Laser Power and Laser Stand-Off Distance on the Geometry of Deposits

Three different laser powers: 1200, 1600 and 2000 W were employed at two different laser stand-offdistances of 9.5 mm and 13 mm. An additional parameter of P = 1800 W with Lo = 13 mm was alsoconsidered as shown in parameter set 7. Figure 4 shows the effect of P on H, W and D of depositshaving processing conditions as tabulated in Table 2.

H decreased when P increased from 1200 W to 1600 W and increased when P was further increasedto 2000 W. η depicted in Figure 3c were 66% and 72% for Lo of 9.5 mm and 13 mm respectively at P of1600 W (decreased H conditions), which is quite low, hence the decrease in H measured. Decreased ηdecreased Mp, which in turn decreased the value of H; as H is a variable dependent on Mp. Similarresults were obtained by Chen et al. [31], wherein the effect of P on H was not ascertained. A study bySegerstark et al. [32] concluded that decrease in V and increase in Mp increased H, but P was foundto have negligible effect on H. A study by Corbin et al. [37] showed that H remained fairly constantover a certain P (250–400 W) and gradually increased with further increase in P when other processparameters remained unchanged, which ascertains the result obtained in the present study. Change inLo did not show any significant effect on H as seen in Figure 4a.

W increased with P in both cases of Lo as shown in Figure 4b. A relatively large amount of energywas available at 2000 W compared to 1200 W and 1600 W, increasing the shape and size of the melt-poolresulting in increased W. The power distribution in case of Lo = 13 mm acts upon a larger area due to

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Ds (2.3 mm) compared to Lo = 9.5 mm (Ds = 1.8 mm) and thereby increases W. The powder captureefficiency has minimal effects on W, and same is seen in case of P = 1600 W as Ds was the dominantfactor. As reported in the introduction, decreased Mp leads to increased W, as more lateral heat transfertakes place for melting the substrate and providing a larger area for the melt flow.

Figure 4c indicates that Lo of 9.5 mm had higher D compared to the Lo of 13 mm due to higherenergy density at 9.5 mm (Table 3) which melted more material. The specific energy had a noticeableeffect on D, as deposits 2 and 3 indicated higher dilution and in the other extreme, deposits 1 and 4showed lower levels of dilution (Figure 3a).Metals 2020, 10, x FOR PEER REVIEW 7 of 16

(a)

(b)

(c)

Figure 4. Influence of laser power of 1200, 1600, 1800 and 2000 W on (a) height; (b) width; and (c) depth of deposits at laser stand-off distances of 9.5 mm and 13 mm at a constant scanning speed of 1100 mm/min.

3.1.2. Effect of Scanning Speed and Laser Stand-Off Distance on the Geometry of Deposits

For studying the effects of scanning speed and laser stand-off distance on geometry, single-tracks processed at two different V of 900 mm/min and 1100 mm/min with Lo of 9.5 mm and 13 mm at P of 1600 W were considered, as in parameter sets 2, 5, 8 and 9.

It can be seen in Figure 5a that the H was insensitive to the changes in Lo. However, with increased V, H decreased. W showed a similar trend as H with respect to V. An increased V showed a decline in W, although the decline was less significant as compared to H. The effect of Lo on W was significant as seen in Figure 5b. A lower Lo of 9.5 mm yielded lower W compared to Lo of 13 mm. D showed a varied relationship as compared to H and W depicted in Figure 5c.

Figure 4. Influence of laser power of 1200, 1600, 1800 and 2000 W on (a) height; (b) width; and (c) depth ofdeposits at laser stand-off distances of 9.5 mm and 13 mm at a constant scanning speed of 1100 mm/min.

3.1.2. Effect of Scanning Speed and Laser Stand-Off Distance on the Geometry of Deposits

For studying the effects of scanning speed and laser stand-off distance on geometry, single-tracksprocessed at two different V of 900 mm/min and 1100 mm/min with Lo of 9.5 mm and 13 mm at P of1600 W were considered, as in parameter sets 2, 5, 8 and 9.

It can be seen in Figure 5a that the H was insensitive to the changes in Lo. However, with increasedV, H decreased. W showed a similar trend as H with respect to V. An increased V showed a decline in

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W, although the decline was less significant as compared to H. The effect of Lo on W was significantas seen in Figure 5b. A lower Lo of 9.5 mm yielded lower W compared to Lo of 13 mm. D showed avaried relationship as compared to H and W depicted in Figure 5c.Metals 2020, 10, x FOR PEER REVIEW 8 of 16

(a) (b)

(c) (d)

Figure 5. Influence of scanning speeds of 900 mm/min and 1100 mm/min on (a) height, (b) width, (c) depth, (d) %dilution of deposits at laser stand-off distances of 9.5 mm and 13 mm at a constant laser power of 1600 W.

With increased V, the specific energy decreased due decrease in interaction time of laser beam on substrate. This reduced the amount of material melted per unit time resulting in decreased H and W, thereby lowering powder capture efficiency (within parentheses) as shown in Figure 5d. The change in W can be explained by increased Ds associated with increased Lo. The energy for substrate melting increased with increased V as the Mp was constant, and the amount of powder available for melting depleted with increased V (defined as line mass, Mp/V).

3.2. Grain Morphology

The cross-sectional morphology for one-half of a single-track deposit is shown in Figure 6a, as the other half of the cross-section was symmetrical. The deposit can be broadly distinguished into three separate regions, the top, the middle and the bottom areas based on morphological differences observed in microstructural feature present in these discrete regions depicted by Figure 6b–d respectively.

Figure 6. (a) A typical morphology of a deposit; magnified images of (b) top; (c) middle; (d) bottom regions of the bead.

Figure 5. Influence of scanning speeds of 900 mm/min and 1100 mm/min on (a) height, (b) width, (c)depth, (d) %dilution of deposits at laser stand-off distances of 9.5 mm and 13 mm at a constant laserpower of 1600 W.

With increased V, the specific energy decreased due decrease in interaction time of laser beam onsubstrate. This reduced the amount of material melted per unit time resulting in decreased H and W,thereby lowering powder capture efficiency (within parentheses) as shown in Figure 5d. The change inW can be explained by increased Ds associated with increased Lo. The energy for substrate meltingincreased with increased V as the Mp was constant, and the amount of powder available for meltingdepleted with increased V (defined as line mass, Mp/V).

3.2. Grain Morphology

The cross-sectional morphology for one-half of a single-track deposit is shown in Figure 6a, as theother half of the cross-section was symmetrical. The deposit can be broadly distinguished into threeseparate regions, the top, the middle and the bottom areas based on morphological differences observedin microstructural feature present in these discrete regions depicted by Figure 6b–d respectively.

The bottom and middle regions closer to the substrate can dissipate heat much faster compared tothe top region, and hence yielded columnar dendritic microstructures and relatively lower segregationof Nb-rich phases [38]. An interplay between thermal gradient (G) and growth rate (R) determined thecooling rate and substructure formation during solid-liquid (S/L) transformation. The G is higher in thebottom region where relatively cooler substrate was in contact with the melt-pool which led to a highG/R ratio which preferentially formed columnar dendritic structures as the planar and cellular growthfronts terminated prematurely [39,40]. The G value gradually decreased towards the top region ofthe deposit where equiaxed grain morphology was noticed [38,40]. The dendritic growth at the topedge of the deposit indicated a secondary growth front that nucleated from partially melted powders

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and also due to shielding gas interaction with the melt-pool, that provided a suitable condition forheterogeneous nucleation away from the solidification front [32]. The heat dissipation by partiallymelted powder particles is minimal as compared to the substrate, and the cooling rates achieved inthis region are low compared to the bottom of the deposit.

Metals 2020, 10, x FOR PEER REVIEW 8 of 16

(a) (b)

(c) (d)

Figure 5. Influence of scanning speeds of 900 mm/min and 1100 mm/min on (a) height, (b) width, (c) depth, (d) %dilution of deposits at laser stand-off distances of 9.5 mm and 13 mm at a constant laser power of 1600 W.

With increased V, the specific energy decreased due decrease in interaction time of laser beam on substrate. This reduced the amount of material melted per unit time resulting in decreased H and W, thereby lowering powder capture efficiency (within parentheses) as shown in Figure 5d. The change in W can be explained by increased Ds associated with increased Lo. The energy for substrate melting increased with increased V as the Mp was constant, and the amount of powder available for melting depleted with increased V (defined as line mass, Mp/V).

3.2. Grain Morphology

The cross-sectional morphology for one-half of a single-track deposit is shown in Figure 6a, as the other half of the cross-section was symmetrical. The deposit can be broadly distinguished into three separate regions, the top, the middle and the bottom areas based on morphological differences observed in microstructural feature present in these discrete regions depicted by Figure 6b–d respectively.

Figure 6. (a) A typical morphology of a deposit; magnified images of (b) top; (c) middle; (d) bottom regions of the bead. Figure 6. (a) A typical morphology of a deposit; magnified images of (b) top; (c) middle; (d) bottomregions of the bead.

‘Fan-like’ microstructures similar to those observed in the top region of the deposit were reportedby Antonsson et al. [41], that narrated the effect of cooling rates and solidification structures formed ofAlloy 718 by Levitation Casting (LC) technique having an undercooling of 40 ◦C at the highest growthtemperature of 1300 ◦C, leading to an estimated cooling rate of 15,000 ◦C/s. Dendrites formed in thetop region of DED deposits considered in this study are well-developed compared to the LC technique,indicating lower undercooling and relatively low cooling rates of 100–600 ◦C/s [27,42].

3.2.1. EBSD Mapping

The inverse pole figure (IPF) maps and respective (001) pole figures depicting the growth orientationand textures formed in different region of deposits 1, 2, 3, 8 and 9 are as shown in Figure 7. Columnargrains were observed in the middle and bottom regions of the deposit for all five conditions considered.The grains at the bottom that grew epitaxially from the substrate were comparatively smaller comparedto those found in the middle region of the deposit [43,44]. Cooling rates (ε) are considered critical ininfluencing the microstructural scales formed during solidification [38]. The cooling rate decreasesgradually with the height of the deposit, thereby producing relatively larger columnar structures awayfrom the substrate.

Another commonality in most deposits analyzed was the orientation of grains. Apart from deposit1, most grains were oriented in<110> and<111> directions for samples 2, 3, 8 and 9. The IPFs in Figure 7showed moderate textures (3.5 to 5.6) along <110> directions in (001) pole figures. The orientation ofthe grains was influenced primarily by the direction of heat flow under high cooling rate conditions.Maximum heat-flux occurs in the direction perpendicular to the S/L growth front. The S/L growth frontrepresented the fusion line between the melt-pool and the substrate at the start of solidification [40].Hence, a columnar grain evolved from the interaction surface at the substrate in the opposite directionto the heat flux. Further, since the melt-pool shape was dynamic and curved at the fusion-line betweenthe substrate and melt-pool, the values of G and R constantly changed. This influenced the growthdirection of columnar grains that originated from the fusion-line to the deposit center-line [38]. Depositsof 2, 3, 8 and 9 conditions had high dilutions as opposed to the deposit 1, which showed the lowestdilution and depicted the strongest texture of the five conditions considered (at 7.8) in the <100>direction, the preferred growth direction in the case of cubic crystals. The interaction angle (α) whichwas the angle between the substrate surface before deposition and a tangent to fusion line at the depositedge (Figure 7a), changed with the dilution and is equal to 20◦ for deposit 1 which was considerablysmaller compared to deposits 2, 3, 8 and 9 (43◦, 46◦, 51◦, 45◦ respectively). The tendency of grain

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W, although the decline was less significant as compared to H. The effect of Lo on W was significantas seen in Figure 5b. A lower Lo of 9.5 mm yielded lower W compared to Lo of 13 mm. D showed avaried relationship as compared to H and W depicted in Figure 5c.Metals 2020, 10, x FOR PEER REVIEW 8 of 16

(a) (b)

(c) (d)

Figure 5. Influence of scanning speeds of 900 mm/min and 1100 mm/min on (a) height, (b) width, (c) depth, (d) %dilution of deposits at laser stand-off distances of 9.5 mm and 13 mm at a constant laser power of 1600 W.

With increased V, the specific energy decreased due decrease in interaction time of laser beam on substrate. This reduced the amount of material melted per unit time resulting in decreased H and W, thereby lowering powder capture efficiency (within parentheses) as shown in Figure 5d. The change in W can be explained by increased Ds associated with increased Lo. The energy for substrate melting increased with increased V as the Mp was constant, and the amount of powder available for melting depleted with increased V (defined as line mass, Mp/V).

3.2. Grain Morphology

The cross-sectional morphology for one-half of a single-track deposit is shown in Figure 6a, as the other half of the cross-section was symmetrical. The deposit can be broadly distinguished into three separate regions, the top, the middle and the bottom areas based on morphological differences observed in microstructural feature present in these discrete regions depicted by Figure 6b–d respectively.

Figure 6. (a) A typical morphology of a deposit; magnified images of (b) top; (c) middle; (d) bottom regions of the bead.

Figure 5. Influence of scanning speeds of 900 mm/min and 1100 mm/min on (a) height, (b) width, (c)depth, (d) %dilution of deposits at laser stand-off distances of 9.5 mm and 13 mm at a constant laserpower of 1600 W.

With increased V, the specific energy decreased due decrease in interaction time of laser beam onsubstrate. This reduced the amount of material melted per unit time resulting in decreased H and W,thereby lowering powder capture efficiency (within parentheses) as shown in Figure 5d. The change inW can be explained by increased Ds associated with increased Lo. The energy for substrate meltingincreased with increased V as the Mp was constant, and the amount of powder available for meltingdepleted with increased V (defined as line mass, Mp/V).

3.2. Grain Morphology

The cross-sectional morphology for one-half of a single-track deposit is shown in Figure 6a, as theother half of the cross-section was symmetrical. The deposit can be broadly distinguished into threeseparate regions, the top, the middle and the bottom areas based on morphological differences observedin microstructural feature present in these discrete regions depicted by Figure 6b–d respectively.

The bottom and middle regions closer to the substrate can dissipate heat much faster compared tothe top region, and hence yielded columnar dendritic microstructures and relatively lower segregationof Nb-rich phases [38]. An interplay between thermal gradient (G) and growth rate (R) determined thecooling rate and substructure formation during solid-liquid (S/L) transformation. The G is higher in thebottom region where relatively cooler substrate was in contact with the melt-pool which led to a highG/R ratio which preferentially formed columnar dendritic structures as the planar and cellular growthfronts terminated prematurely [39,40]. The G value gradually decreased towards the top region ofthe deposit where equiaxed grain morphology was noticed [38,40]. The dendritic growth at the topedge of the deposit indicated a secondary growth front that nucleated from partially melted powders

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and also due to shielding gas interaction with the melt-pool, that provided a suitable condition forheterogeneous nucleation away from the solidification front [32]. The heat dissipation by partiallymelted powder particles is minimal as compared to the substrate, and the cooling rates achieved inthis region are low compared to the bottom of the deposit.

Metals 2020, 10, x FOR PEER REVIEW 8 of 16

(a) (b)

(c) (d)

Figure 5. Influence of scanning speeds of 900 mm/min and 1100 mm/min on (a) height, (b) width, (c) depth, (d) %dilution of deposits at laser stand-off distances of 9.5 mm and 13 mm at a constant laser power of 1600 W.

With increased V, the specific energy decreased due decrease in interaction time of laser beam on substrate. This reduced the amount of material melted per unit time resulting in decreased H and W, thereby lowering powder capture efficiency (within parentheses) as shown in Figure 5d. The change in W can be explained by increased Ds associated with increased Lo. The energy for substrate melting increased with increased V as the Mp was constant, and the amount of powder available for melting depleted with increased V (defined as line mass, Mp/V).

3.2. Grain Morphology

The cross-sectional morphology for one-half of a single-track deposit is shown in Figure 6a, as the other half of the cross-section was symmetrical. The deposit can be broadly distinguished into three separate regions, the top, the middle and the bottom areas based on morphological differences observed in microstructural feature present in these discrete regions depicted by Figure 6b–d respectively.

Figure 6. (a) A typical morphology of a deposit; magnified images of (b) top; (c) middle; (d) bottom regions of the bead. Figure 6. (a) A typical morphology of a deposit; magnified images of (b) top; (c) middle; (d) bottomregions of the bead.

‘Fan-like’ microstructures similar to those observed in the top region of the deposit were reportedby Antonsson et al. [41], that narrated the effect of cooling rates and solidification structures formed ofAlloy 718 by Levitation Casting (LC) technique having an undercooling of 40 ◦C at the highest growthtemperature of 1300 ◦C, leading to an estimated cooling rate of 15,000 ◦C/s. Dendrites formed in thetop region of DED deposits considered in this study are well-developed compared to the LC technique,indicating lower undercooling and relatively low cooling rates of 100–600 ◦C/s [27,42].

3.2.1. EBSD Mapping

The inverse pole figure (IPF) maps and respective (001) pole figures depicting the growth orientationand textures formed in different region of deposits 1, 2, 3, 8 and 9 are as shown in Figure 7. Columnargrains were observed in the middle and bottom regions of the deposit for all five conditions considered.The grains at the bottom that grew epitaxially from the substrate were comparatively smaller comparedto those found in the middle region of the deposit [43,44]. Cooling rates (ε) are considered critical ininfluencing the microstructural scales formed during solidification [38]. The cooling rate decreasesgradually with the height of the deposit, thereby producing relatively larger columnar structures awayfrom the substrate.

Another commonality in most deposits analyzed was the orientation of grains. Apart from deposit1, most grains were oriented in<110> and<111> directions for samples 2, 3, 8 and 9. The IPFs in Figure 7showed moderate textures (3.5 to 5.6) along <110> directions in (001) pole figures. The orientation ofthe grains was influenced primarily by the direction of heat flow under high cooling rate conditions.Maximum heat-flux occurs in the direction perpendicular to the S/L growth front. The S/L growth frontrepresented the fusion line between the melt-pool and the substrate at the start of solidification [40].Hence, a columnar grain evolved from the interaction surface at the substrate in the opposite directionto the heat flux. Further, since the melt-pool shape was dynamic and curved at the fusion-line betweenthe substrate and melt-pool, the values of G and R constantly changed. This influenced the growthdirection of columnar grains that originated from the fusion-line to the deposit center-line [38]. Depositsof 2, 3, 8 and 9 conditions had high dilutions as opposed to the deposit 1, which showed the lowestdilution and depicted the strongest texture of the five conditions considered (at 7.8) in the <100>direction, the preferred growth direction in the case of cubic crystals. The interaction angle (α) whichwas the angle between the substrate surface before deposition and a tangent to fusion line at the depositedge (Figure 7a), changed with the dilution and is equal to 20◦ for deposit 1 which was considerablysmaller compared to deposits 2, 3, 8 and 9 (43◦, 46◦, 51◦, 45◦ respectively). The tendency of grain

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growth towards the weld centerline was affected by the amount of dilution indicated by α. The grainsoriented in <110>were observed to originate from the fusion line with α nearly equal to 45◦. From(100) pole figures, it is apparent that the <110> direction is at an angle of 45◦ to the <001> direction.Hence if the reference plane is changed with respect to the point of origin of the grain, the direction ofgrowth of most crystals would be close to the <001> direction.

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The grain size measured in the middle regions for columnar grains showed a strong dependence on the laser power irrespective of scanning speeds and stand-off distances. At 1200 W corresponding to deposit 1, an average grain size of 58 µm (ASTM 4.9) was measured. Deposit 3 processed at a laser power of 2000 W had a larger grain size of 88 µm (ASTM 3.7). The average grain size at 1600 W was equal to 75 µm (ASTM 4.2), 75 µm (ASTM 4.2) and 78 µm (ASTM 4.1) for deposit 2, 8 and 9 respectively. The grain heights and widths were measured along A1 and A2 directions as shown in Figure 7, along two different coordinate axes and along the length of the deposit. In order to incorporate most grains, the grain sizes were measured close to the deposit centerline where the power density was highest as per Gaussian form of energy distribution (as indicated in Figure 1a). Hence laser power principally influenced the power density factor of various parameter sets, thereby influencing the average grain size in the middle region of the deposit.

(a) (b) (c)

(d) (e)

Figure 7. (100) IPFs (above) showing grain orientation along growth direction and respective pole figures (below) for process parameter conditions as in Table 2. The dotted line in (a,b,d) indicate columnar to equiaxed transition (CET), which is confined to a very small region in (c,e). α depicts the interaction angle made by the deposit with the substrate.

Much finer equiaxed grains were observed at the top part of the deposit and a bimodal grain structure existed in an intermediate layer between the columnar and equiaxed grains in deposits 1, 2 and 8. The dotted line in Figure 7a,b,d represents this change in morphology known as Columnar to Equiaxed Transition (CET) [45,46]. This is arguably due to constitutional supercooling, wherein the liquid in front of the S/L interface becomes progressively enriched with elements having low partition co-efficients (predominantly Nb in case of Alloy 718), providing the undercooling necessary for

Figure 7. (100) IPFs (above) showing grain orientation along growth direction and respective polefigures (below) for process parameter conditions as in Table 2. The dotted line in (a,b,d) indicatecolumnar to equiaxed transition (CET), which is confined to a very small region in (c,e). α depicts theinteraction angle made by the deposit with the substrate.

The grain size measured in the middle regions for columnar grains showed a strong dependenceon the laser power irrespective of scanning speeds and stand-off distances. At 1200 W correspondingto deposit 1, an average grain size of 58 µm (ASTM 4.9) was measured. Deposit 3 processed at a laserpower of 2000 W had a larger grain size of 88 µm (ASTM 3.7). The average grain size at 1600 W wasequal to 75 µm (ASTM 4.2), 75 µm (ASTM 4.2) and 78 µm (ASTM 4.1) for deposit 2, 8 and 9 respectively.The grain heights and widths were measured along A1 and A2 directions as shown in Figure 7, alongtwo different coordinate axes and along the length of the deposit. In order to incorporate most grains,the grain sizes were measured close to the deposit centerline where the power density was highest asper Gaussian form of energy distribution (as indicated in Figure 1a). Hence laser power principallyinfluenced the power density factor of various parameter sets, thereby influencing the average grainsize in the middle region of the deposit.

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Much finer equiaxed grains were observed at the top part of the deposit and a bimodal grainstructure existed in an intermediate layer between the columnar and equiaxed grains in deposits 1, 2and 8. The dotted line in Figure 7a,b,d represents this change in morphology known as Columnarto Equiaxed Transition (CET) [45,46]. This is arguably due to constitutional supercooling, whereinthe liquid in front of the S/L interface becomes progressively enriched with elements having lowpartition co-efficients (predominantly Nb in case of Alloy 718), providing the undercooling necessaryfor heterogeneous nucleation of grains in the solute-rich melt-pool, observed in the top region of DEDdeposits [44,47].

Comparison of IPFs of deposits pertaining to parameter sets 1, 8 and 3 provides an insight intothe effect of varied laser power on grain structure. Notably, the diminished presence of the topregion in deposits 3 and 9 was of interest as it is believed to have implications for the segregationof Nb-rich phases when subsequent layers of deposit are built on top of the existing single trackspecimen. The secondary growth front in these cases was hindered by the high-temperature gradientsthat remelted any existing dendrites and heterogeneously nucleated grains present at the top regionof the deposit. The effect of varied scanning speed was negligible as deposit 2 (V = 900 mm/min)and deposit 8 (V = 1100 mm/min) did not show appreciable variation. Deposit 9 (Lo = 13 mm) wasmore successful in eliminating the top region segregation, and hence the formation of equiaxed grainscompared to deposit 8 (Lo = 9.5 mm).

3.2.2. Nb-Rich Eutectic Phases

The SEM micrographs in Figure 8a–c depict the segregation patterns of Nb-rich eutectic phasesrepresented by white particulates in the top, middle and bottom regions respectively. There are twomain eutectic reactions known to occur in Alloy 718: 1280 ◦C to 1250 ◦C wherein γ/NbC is precipitatedand the γ/laves phase occurring in the temperature range of 1075 ◦C to 1200 ◦C [19,41], dependingon the specific alloy composition and processing conditions. Local solidification conditions at theinter-dendritic Nb rich regions formed NbCs and laves phases (Figure 8). The EDS point analysesrevealed that the white particulates were indeed Nb-rich precipitates as shown in Figure 8d, for whichthe elemental analysis is tabulated in Table 4. Along the height of the deposit further away from thesubstrate, the segregation was higher and the cooling rates also decreased due to limited heat fluxconditions, leading to higher inter-dendritic spacing and precipitation of more Nb-rich phases [25,28].It is apparent from Figure 8a–c that the top regions have higher segregation of such Nb-rich eutecticphases compared to other regions, which was confirmed by the area fraction analyses as shown inFigure 8e. The modelling work performed by Nie et al. [27] provides insight into the expected coolingrates along the height of the bead and consequently the variation in secondary dendrite arm spacingfor Alloy 718.

Table 4. Elemental analysis of point spectrum locations depicted in Figure 8d.

ElementWeight %

Spectrum 1 Spectrum 2 Spectrum 3 Spectrum 4

Ni 46.86 38.14 45.11 40.67Cr 18.03 13.02 11.77 13.54Fe 17.95 11.05 10.05 11.43Nb 2.2 13.06 16.3 17.89Al 0.47 0.34 0.28 0.3Ti 0.52 1.19 1.33 1.65

Mo 2.79 3.99 - -C 0.22 5.3 3.73 4.68

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growth towards the weld centerline was affected by the amount of dilution indicated by α. The grainsoriented in <110>were observed to originate from the fusion line with α nearly equal to 45◦. From(100) pole figures, it is apparent that the <110> direction is at an angle of 45◦ to the <001> direction.Hence if the reference plane is changed with respect to the point of origin of the grain, the direction ofgrowth of most crystals would be close to the <001> direction.

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The grain size measured in the middle regions for columnar grains showed a strong dependence on the laser power irrespective of scanning speeds and stand-off distances. At 1200 W corresponding to deposit 1, an average grain size of 58 µm (ASTM 4.9) was measured. Deposit 3 processed at a laser power of 2000 W had a larger grain size of 88 µm (ASTM 3.7). The average grain size at 1600 W was equal to 75 µm (ASTM 4.2), 75 µm (ASTM 4.2) and 78 µm (ASTM 4.1) for deposit 2, 8 and 9 respectively. The grain heights and widths were measured along A1 and A2 directions as shown in Figure 7, along two different coordinate axes and along the length of the deposit. In order to incorporate most grains, the grain sizes were measured close to the deposit centerline where the power density was highest as per Gaussian form of energy distribution (as indicated in Figure 1a). Hence laser power principally influenced the power density factor of various parameter sets, thereby influencing the average grain size in the middle region of the deposit.

(a) (b) (c)

(d) (e)

Figure 7. (100) IPFs (above) showing grain orientation along growth direction and respective pole figures (below) for process parameter conditions as in Table 2. The dotted line in (a,b,d) indicate columnar to equiaxed transition (CET), which is confined to a very small region in (c,e). α depicts the interaction angle made by the deposit with the substrate.

Much finer equiaxed grains were observed at the top part of the deposit and a bimodal grain structure existed in an intermediate layer between the columnar and equiaxed grains in deposits 1, 2 and 8. The dotted line in Figure 7a,b,d represents this change in morphology known as Columnar to Equiaxed Transition (CET) [45,46]. This is arguably due to constitutional supercooling, wherein the liquid in front of the S/L interface becomes progressively enriched with elements having low partition co-efficients (predominantly Nb in case of Alloy 718), providing the undercooling necessary for

Figure 7. (100) IPFs (above) showing grain orientation along growth direction and respective polefigures (below) for process parameter conditions as in Table 2. The dotted line in (a,b,d) indicatecolumnar to equiaxed transition (CET), which is confined to a very small region in (c,e). α depicts theinteraction angle made by the deposit with the substrate.

The grain size measured in the middle regions for columnar grains showed a strong dependenceon the laser power irrespective of scanning speeds and stand-off distances. At 1200 W correspondingto deposit 1, an average grain size of 58 µm (ASTM 4.9) was measured. Deposit 3 processed at a laserpower of 2000 W had a larger grain size of 88 µm (ASTM 3.7). The average grain size at 1600 W wasequal to 75 µm (ASTM 4.2), 75 µm (ASTM 4.2) and 78 µm (ASTM 4.1) for deposit 2, 8 and 9 respectively.The grain heights and widths were measured along A1 and A2 directions as shown in Figure 7, alongtwo different coordinate axes and along the length of the deposit. In order to incorporate most grains,the grain sizes were measured close to the deposit centerline where the power density was highest asper Gaussian form of energy distribution (as indicated in Figure 1a). Hence laser power principallyinfluenced the power density factor of various parameter sets, thereby influencing the average grainsize in the middle region of the deposit.

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Much finer equiaxed grains were observed at the top part of the deposit and a bimodal grainstructure existed in an intermediate layer between the columnar and equiaxed grains in deposits 1, 2and 8. The dotted line in Figure 7a,b,d represents this change in morphology known as Columnarto Equiaxed Transition (CET) [45,46]. This is arguably due to constitutional supercooling, whereinthe liquid in front of the S/L interface becomes progressively enriched with elements having lowpartition co-efficients (predominantly Nb in case of Alloy 718), providing the undercooling necessaryfor heterogeneous nucleation of grains in the solute-rich melt-pool, observed in the top region of DEDdeposits [44,47].

Comparison of IPFs of deposits pertaining to parameter sets 1, 8 and 3 provides an insight intothe effect of varied laser power on grain structure. Notably, the diminished presence of the topregion in deposits 3 and 9 was of interest as it is believed to have implications for the segregationof Nb-rich phases when subsequent layers of deposit are built on top of the existing single trackspecimen. The secondary growth front in these cases was hindered by the high-temperature gradientsthat remelted any existing dendrites and heterogeneously nucleated grains present at the top regionof the deposit. The effect of varied scanning speed was negligible as deposit 2 (V = 900 mm/min)and deposit 8 (V = 1100 mm/min) did not show appreciable variation. Deposit 9 (Lo = 13 mm) wasmore successful in eliminating the top region segregation, and hence the formation of equiaxed grainscompared to deposit 8 (Lo = 9.5 mm).

3.2.2. Nb-Rich Eutectic Phases

The SEM micrographs in Figure 8a–c depict the segregation patterns of Nb-rich eutectic phasesrepresented by white particulates in the top, middle and bottom regions respectively. There are twomain eutectic reactions known to occur in Alloy 718: 1280 ◦C to 1250 ◦C wherein γ/NbC is precipitatedand the γ/laves phase occurring in the temperature range of 1075 ◦C to 1200 ◦C [19,41], dependingon the specific alloy composition and processing conditions. Local solidification conditions at theinter-dendritic Nb rich regions formed NbCs and laves phases (Figure 8). The EDS point analysesrevealed that the white particulates were indeed Nb-rich precipitates as shown in Figure 8d, for whichthe elemental analysis is tabulated in Table 4. Along the height of the deposit further away from thesubstrate, the segregation was higher and the cooling rates also decreased due to limited heat fluxconditions, leading to higher inter-dendritic spacing and precipitation of more Nb-rich phases [25,28].It is apparent from Figure 8a–c that the top regions have higher segregation of such Nb-rich eutecticphases compared to other regions, which was confirmed by the area fraction analyses as shown inFigure 8e. The modelling work performed by Nie et al. [27] provides insight into the expected coolingrates along the height of the bead and consequently the variation in secondary dendrite arm spacingfor Alloy 718.

Table 4. Elemental analysis of point spectrum locations depicted in Figure 8d.

ElementWeight %

Spectrum 1 Spectrum 2 Spectrum 3 Spectrum 4

Ni 46.86 38.14 45.11 40.67Cr 18.03 13.02 11.77 13.54Fe 17.95 11.05 10.05 11.43Nb 2.2 13.06 16.3 17.89Al 0.47 0.34 0.28 0.3Ti 0.52 1.19 1.33 1.65

Mo 2.79 3.99 - -C 0.22 5.3 3.73 4.68

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heterogeneous nucleation of grains in the solute-rich melt-pool, observed in the top region of DED deposits [44,47].

Comparison of IPFs of deposits pertaining to parameter sets 1, 8 and 3 provides an insight into the effect of varied laser power on grain structure. Notably, the diminished presence of the top region in deposits 3 and 9 was of interest as it is believed to have implications for the segregation of Nb-rich phases when subsequent layers of deposit are built on top of the existing single track specimen. The secondary growth front in these cases was hindered by the high-temperature gradients that remelted any existing dendrites and heterogeneously nucleated grains present at the top region of the deposit. The effect of varied scanning speed was negligible as deposit 2 (V = 900 mm/min) and deposit 8 (V = 1100 mm/min) did not show appreciable variation. Deposit 9 (Lo = 13 mm) was more successful in eliminating the top region segregation, and hence the formation of equiaxed grains compared to deposit 8 (Lo = 9.5 mm).

3.2.2. Nb-Rich Eutectic Phases

The SEM micrographs in Figure 8a–c depict the segregation patterns of Nb-rich eutectic phases represented by white particulates in the top, middle and bottom regions respectively. There are two main eutectic reactions known to occur in Alloy 718: 1280 °C to 1250 °C wherein γ/NbC is precipitated and the γ/laves phase occurring in the temperature range of 1075 °C to 1200 °C [19,41], depending on the specific alloy composition and processing conditions. Local solidification conditions at the inter-dendritic Nb rich regions formed NbCs and laves phases (Figure 8). The EDS point analyses revealed that the white particulates were indeed Nb-rich precipitates as shown in Figure 8d, for which the elemental analysis is tabulated in Table 4. Along the height of the deposit further away from the substrate, the segregation was higher and the cooling rates also decreased due to limited heat flux conditions, leading to higher inter-dendritic spacing and precipitation of more Nb-rich phases [25,28]. It is apparent from Figure 8a–c that the top regions have higher segregation of such Nb-rich eutectic phases compared to other regions, which was confirmed by the area fraction analyses as shown in Figure 8e. The modelling work performed by Nie et al. [27] provides insight into the expected cooling rates along the height of the bead and consequently the variation in secondary dendrite arm spacing for Alloy 718.

(a) (b) (c)

(d) (e)

Figure 8. SEM Micrographs depicting variation in Nb-rich phase fraction in (a) top; (b) middle; (c)bottom regions in a deposition pertaining to parameter set 1 (refer to Table 2). (d) EDS mapping ofNb-rich precipitates (Spectrum 2–4) and matrix (Spectrum 1); (e) plot showing area fraction of Nb-richeutectic phases for various parameters.

3.3. Effect of Laser Power, Scanning Speed and Laser Stand-Off Distance on Nb-Rich Phase Fraction

A qualitative conclusion can be reached by studying the effect of laser power on Nb-rich phaseformation as depicted in Figure 9a–c. The Nb-rich phase fraction decreased with increased P in thecase of the higher Lo of 13 mm. An opposite trend was noticed in the 9.5 mm stand-off distancecondition in top and bottom regions. The segregation in the top region was primarily determined bythe CET and secondary growth fronts. CET was controlled by the solute rejection and build up aheadof S/L front combined with convective currents present in the melt-pool. At higher specific energies,the thermal gradient was high enough to persist until the end of solidification of the entire depositdeterring heterogeneous nucleation of equiaxed grains. Also, the convective currents aid in bettersolute dissolution in the melt ahead of the growth front, thereby hindering the formation of Nb-richphases with increasing power throughout the deposit [38,48].

The effect of scanning speeds was much clearer at a constant laser power of 1600 W. Increasedscanning speed yielded fewer Nb-rich phases as in Figure 9d, with only exception of line indicatingLo= 9.5 mm in the bottom region of deposit. At speeds of 900 mm/min, the conditions having a stand-offdistance of 9.5 mm had higher Nb-rich area fraction compared to Lo of 13 mm. At 1100 mm/min,stand-off distances had a minimal effect on Nb-rich area fraction.

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Figure 8. SEM Micrographs depicting variation in Nb-rich phase fraction in (a) top; (b) middle; (c) bottom regions in a deposition pertaining to parameter set 1 (refer to Table 2). (d) EDS mapping of Nb-rich precipitates (Spectrum 2–4) and matrix (Spectrum 1); (e) plot showing area fraction of Nb-rich eutectic phases for various parameters.

Table 4. Elemental analysis of point spectrum locations depicted in Figure 8d.

Element Weight %

Spectrum 1 Spectrum 2 Spectrum 3 Spectrum 4 Ni 46.86 38.14 45.11 40.67 Cr 18.03 13.02 11.77 13.54 Fe 17.95 11.05 10.05 11.43 Nb 2.2 13.06 16.3 17.89 Al 0.47 0.34 0.28 0.3 Ti 0.52 1.19 1.33 1.65

Mo 2.79 3.99 - - C 0.22 5.3 3.73 4.68

3.3. Effect of Laser Power, Scanning Speed and Laser Stand-Off Distance on Nb-Rich Phase Fraction

(a) (b)

(c) (d)

Figure 9. Plots depicting the effect of laser power and laser stand-off distance on (a) top region; (b) middle region; (c) bottom region. (d) Effect of scanning speed and laser stand-off distance on Nb-rich eutectic phase fractions.

4. Conclusions

Laser power, scanning speed and laser stand-off distance were found to influence the geometry, microstructure and texture in DED single track specimens. The height of the deposits decreased with increased scanning speeds. Stand-off distance had no influence on height of the deposits. The width increased with increased laser power and stand-off distance but decreased with increased scanning speeds. Depth of the deposit increased with increased laser power and decreased with increase in

Figure 9. Plots depicting the effect of laser power and laser stand-off distance on (a) top region; (b)middle region; (c) bottom region. (d) Effect of scanning speed and laser stand-off distance on Nb-richeutectic phase fractions.

4. Conclusions

Laser power, scanning speed and laser stand-off distance were found to influence the geometry,microstructure and texture in DED single track specimens. The height of the deposits decreased withincreased scanning speeds. Stand-off distance had no influence on height of the deposits. The widthincreased with increased laser power and stand-off distance but decreased with increased scanningspeeds. Depth of the deposit increased with increased laser power and decreased with increasein stand-off distance. Deposits 3, 6 through 9 had aspect ratio (W/H) of more than two. The areaof dilution increased with increased specific energy. The powder capture efficiency increased withthe increase inproduct of specific energy and spot diameter and minutely increased with increasedstand-off distance. Deposits 3, 5, 6 and 7 pertaining to high power and stand-off distance (except 3)showed characteristics of effective powder utilization (more than 80%).

A single-track deposit had three distinct regions formed by varied cooling rates. Segregation ofNb-rich phases was found to be dominant in the top region of the deposit in all cases. Deposits 6–9associated with moderate specific energy conditions proved to be effective in controlling segregationas they had Nb-rich phase fraction of less than 4% in the middle and bottom regions and under 6% inthe top region. Scanning speed was found to be very influential in decreased area fraction of Nb-richphases formed. Higher stand-off distances also showed indication of reducing the segregation.

Columnar grains dominated the deposits in all cases with a narrow region of equiaxed grains atthe top. In cases of deposits that had better aspect ratio (as in deposit 3 and 9 from EBSD mapping),the equiaxed grains were confined to a minute region at the top regions of the deposit. Differentmorphology and orientation can be achieved by varying energy conditions. The preferred growthdirection for columnar grains is <100>. A center-line solidification was noticed in cases of grains thathad <110>, <111> orientations because of their high dilution rate, with interaction angles of nearly 45◦.Dilution was influenced by specific energy, and hence the texture morphology was affected by power,speed and stand-off distances.

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heterogeneous nucleation of grains in the solute-rich melt-pool, observed in the top region of DED deposits [44,47].

Comparison of IPFs of deposits pertaining to parameter sets 1, 8 and 3 provides an insight into the effect of varied laser power on grain structure. Notably, the diminished presence of the top region in deposits 3 and 9 was of interest as it is believed to have implications for the segregation of Nb-rich phases when subsequent layers of deposit are built on top of the existing single track specimen. The secondary growth front in these cases was hindered by the high-temperature gradients that remelted any existing dendrites and heterogeneously nucleated grains present at the top region of the deposit. The effect of varied scanning speed was negligible as deposit 2 (V = 900 mm/min) and deposit 8 (V = 1100 mm/min) did not show appreciable variation. Deposit 9 (Lo = 13 mm) was more successful in eliminating the top region segregation, and hence the formation of equiaxed grains compared to deposit 8 (Lo = 9.5 mm).

3.2.2. Nb-Rich Eutectic Phases

The SEM micrographs in Figure 8a–c depict the segregation patterns of Nb-rich eutectic phases represented by white particulates in the top, middle and bottom regions respectively. There are two main eutectic reactions known to occur in Alloy 718: 1280 °C to 1250 °C wherein γ/NbC is precipitated and the γ/laves phase occurring in the temperature range of 1075 °C to 1200 °C [19,41], depending on the specific alloy composition and processing conditions. Local solidification conditions at the inter-dendritic Nb rich regions formed NbCs and laves phases (Figure 8). The EDS point analyses revealed that the white particulates were indeed Nb-rich precipitates as shown in Figure 8d, for which the elemental analysis is tabulated in Table 4. Along the height of the deposit further away from the substrate, the segregation was higher and the cooling rates also decreased due to limited heat flux conditions, leading to higher inter-dendritic spacing and precipitation of more Nb-rich phases [25,28]. It is apparent from Figure 8a–c that the top regions have higher segregation of such Nb-rich eutectic phases compared to other regions, which was confirmed by the area fraction analyses as shown in Figure 8e. The modelling work performed by Nie et al. [27] provides insight into the expected cooling rates along the height of the bead and consequently the variation in secondary dendrite arm spacing for Alloy 718.

(a) (b) (c)

(d) (e)

Figure 8. SEM Micrographs depicting variation in Nb-rich phase fraction in (a) top; (b) middle; (c)bottom regions in a deposition pertaining to parameter set 1 (refer to Table 2). (d) EDS mapping ofNb-rich precipitates (Spectrum 2–4) and matrix (Spectrum 1); (e) plot showing area fraction of Nb-richeutectic phases for various parameters.

3.3. Effect of Laser Power, Scanning Speed and Laser Stand-Off Distance on Nb-Rich Phase Fraction

A qualitative conclusion can be reached by studying the effect of laser power on Nb-rich phaseformation as depicted in Figure 9a–c. The Nb-rich phase fraction decreased with increased P in thecase of the higher Lo of 13 mm. An opposite trend was noticed in the 9.5 mm stand-off distancecondition in top and bottom regions. The segregation in the top region was primarily determined bythe CET and secondary growth fronts. CET was controlled by the solute rejection and build up aheadof S/L front combined with convective currents present in the melt-pool. At higher specific energies,the thermal gradient was high enough to persist until the end of solidification of the entire depositdeterring heterogeneous nucleation of equiaxed grains. Also, the convective currents aid in bettersolute dissolution in the melt ahead of the growth front, thereby hindering the formation of Nb-richphases with increasing power throughout the deposit [38,48].

The effect of scanning speeds was much clearer at a constant laser power of 1600 W. Increasedscanning speed yielded fewer Nb-rich phases as in Figure 9d, with only exception of line indicatingLo= 9.5 mm in the bottom region of deposit. At speeds of 900 mm/min, the conditions having a stand-offdistance of 9.5 mm had higher Nb-rich area fraction compared to Lo of 13 mm. At 1100 mm/min,stand-off distances had a minimal effect on Nb-rich area fraction.

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Figure 8. SEM Micrographs depicting variation in Nb-rich phase fraction in (a) top; (b) middle; (c) bottom regions in a deposition pertaining to parameter set 1 (refer to Table 2). (d) EDS mapping of Nb-rich precipitates (Spectrum 2–4) and matrix (Spectrum 1); (e) plot showing area fraction of Nb-rich eutectic phases for various parameters.

Table 4. Elemental analysis of point spectrum locations depicted in Figure 8d.

Element Weight %

Spectrum 1 Spectrum 2 Spectrum 3 Spectrum 4 Ni 46.86 38.14 45.11 40.67 Cr 18.03 13.02 11.77 13.54 Fe 17.95 11.05 10.05 11.43 Nb 2.2 13.06 16.3 17.89 Al 0.47 0.34 0.28 0.3 Ti 0.52 1.19 1.33 1.65

Mo 2.79 3.99 - - C 0.22 5.3 3.73 4.68

3.3. Effect of Laser Power, Scanning Speed and Laser Stand-Off Distance on Nb-Rich Phase Fraction

(a) (b)

(c) (d)

Figure 9. Plots depicting the effect of laser power and laser stand-off distance on (a) top region; (b) middle region; (c) bottom region. (d) Effect of scanning speed and laser stand-off distance on Nb-rich eutectic phase fractions.

4. Conclusions

Laser power, scanning speed and laser stand-off distance were found to influence the geometry, microstructure and texture in DED single track specimens. The height of the deposits decreased with increased scanning speeds. Stand-off distance had no influence on height of the deposits. The width increased with increased laser power and stand-off distance but decreased with increased scanning speeds. Depth of the deposit increased with increased laser power and decreased with increase in

Figure 9. Plots depicting the effect of laser power and laser stand-off distance on (a) top region; (b)middle region; (c) bottom region. (d) Effect of scanning speed and laser stand-off distance on Nb-richeutectic phase fractions.

4. Conclusions

Laser power, scanning speed and laser stand-off distance were found to influence the geometry,microstructure and texture in DED single track specimens. The height of the deposits decreased withincreased scanning speeds. Stand-off distance had no influence on height of the deposits. The widthincreased with increased laser power and stand-off distance but decreased with increased scanningspeeds. Depth of the deposit increased with increased laser power and decreased with increasein stand-off distance. Deposits 3, 6 through 9 had aspect ratio (W/H) of more than two. The areaof dilution increased with increased specific energy. The powder capture efficiency increased withthe increase inproduct of specific energy and spot diameter and minutely increased with increasedstand-off distance. Deposits 3, 5, 6 and 7 pertaining to high power and stand-off distance (except 3)showed characteristics of effective powder utilization (more than 80%).

A single-track deposit had three distinct regions formed by varied cooling rates. Segregation ofNb-rich phases was found to be dominant in the top region of the deposit in all cases. Deposits 6–9associated with moderate specific energy conditions proved to be effective in controlling segregationas they had Nb-rich phase fraction of less than 4% in the middle and bottom regions and under 6% inthe top region. Scanning speed was found to be very influential in decreased area fraction of Nb-richphases formed. Higher stand-off distances also showed indication of reducing the segregation.

Columnar grains dominated the deposits in all cases with a narrow region of equiaxed grains atthe top. In cases of deposits that had better aspect ratio (as in deposit 3 and 9 from EBSD mapping),the equiaxed grains were confined to a minute region at the top regions of the deposit. Differentmorphology and orientation can be achieved by varying energy conditions. The preferred growthdirection for columnar grains is <100>. A center-line solidification was noticed in cases of grains thathad <110>, <111> orientations because of their high dilution rate, with interaction angles of nearly 45◦.Dilution was influenced by specific energy, and hence the texture morphology was affected by power,speed and stand-off distances.

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Moderate to high powers (P = 1800 and 2000 W), high speed (V = 1100 mm/min) and high stand-offdistances (Lo = 13 mm) corresponding to deposits 3, 6 and 7 proved to be influential in achieving goodgeometrical, microstructural and segregation properties. Hence, they are more suitable for buildingmultilayered deposits as they offer better characteristics in terms of aspect ratio, dilution and powdercapture efficiency. The segregation of Nb-rich phases at the top of these deposits are minimal and theequiaxed region at the top is minimal in these deposits.

Author Contributions: Conceptualization, methodology, investigation, writing—original draft preparation, S.S.;supervision, methodology, writing—review and editing, J.A., S.J.; EBSD investigation, E.G.; experimental work,K.H. All authors have read and agreed to the published version of the manuscript.

Funding: This research was funded by Vinnova, KK- SUMAN-Next and CAM2.

Acknowledgments: The authors acknowledge the support of GKN Aerospace AB.

Conflicts of Interest: The authors declare no conflict of interest.

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Moderate to high powers (P = 1800 and 2000 W), high speed (V = 1100 mm/min) and high stand-offdistances (Lo = 13 mm) corresponding to deposits 3, 6 and 7 proved to be influential in achieving goodgeometrical, microstructural and segregation properties. Hence, they are more suitable for buildingmultilayered deposits as they offer better characteristics in terms of aspect ratio, dilution and powdercapture efficiency. The segregation of Nb-rich phases at the top of these deposits are minimal and theequiaxed region at the top is minimal in these deposits.

Author Contributions: Conceptualization, methodology, investigation, writing—original draft preparation, S.S.;supervision, methodology, writing—review and editing, J.A., S.J.; EBSD investigation, E.G.; experimental work,K.H. All authors have read and agreed to the published version of the manuscript.

Funding: This research was funded by Vinnova, KK- SUMAN-Next and CAM2.

Acknowledgments: The authors acknowledge the support of GKN Aerospace AB.

Conflicts of Interest: The authors declare no conflict of interest.

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Laser-Directed Energy Deposition: Influence of Process Parameters and Heat-TreatmentsLaser-Directed Energy Deposition (L-DED), an Additive Manufacturing (AM) process used for the fabrication of parts in a layer-wise approach has displayed an immense potential over the last decade. The aerospace industry stands as the primary beneficiary due to the L-DED process capability to build near-net-shape components with mini-mal tooling and thereby producing minimum wastage because of reduced machining. The widespread use of Alloy 718 in the aero-engine application has prompted huge research interest in the development of L-DED processing of this superalloy. AM processes are hindered by low build rates and high cycle times which directly affects the process costs. To overcome these issues, the present work focusses on obtaining high deposition rates through a high material feed.Studying the influence of process parameters during the L-DED process is of prime importance as they determine the performance of in-service structures. In the present work, process parameters such as laser power, scanning speed, feed rate and stand-off distances are varied and their influence on geometry and microstructure of Alloy 718 single-track deposits are analyzed. The geometry of deposits is measured in terms of height, width and depth; and the powder capture efficiency is determined by measuring areas of deposition and dilution. The microstructure of the deposits shows a columnar dendritic structure in the middle and bottom region of the deposits and equiaxed grains in the top region. Nb-rich segregation involving laves and NbC phases, typical of Alloy 718 is found in the interdendritic regions and grain boundaries. The segregation increases along the height of the deposit with the bottom region having the least and the top region showing the highest concentration of Nb-rich phases due to the variation in cooling rates. A high laser power (1600 W – 2000 W) and a high scanning speed (1100 mm/min) are found to be the preferable processing conditions for minimizing segregation.Another approach to minimize segregation is by performing post-build heat treatments. The solution treatment (954 °C/1 hr) and double aging (718 °C/8 hr + 621 °C/ 8 hr) standardized for the wrought form of Alloy 718 is applied to as-built deposits which showed a reduction in segregation due to the dissolution of Nb-rich phases. Upon solution treatment, this reduction is accompanied by precipitation of the delta phase, found predominantly in top and bottom regions and sparsely in the middle region of the deposit.

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