Laser Additive Manufacturing of Oxide Dispersion ...

Laser Additive Manufacturing of

Oxide Dispersion Strengthened Steels

and Cu-Cr-Nb Alloys

Von der Fakultät für Georessourcen und Materialtechnik

der Rheinisch-Westfälischen Technischen Hochschule Aachen

zur Erlangung des akademischen Grades eines

Doktors der Ingenieurwissenschaften

genehmigte Dissertation

vorgelegt von M.Sc.

Anoop Raghunath Kini

aus Manipal Karnataka, Indien

Berichter: Prof. Dr.-Ing. Dierk Raabe

Univ.-Prof. Jochen M. Schneider, Ph.D.

Tag der mündlichen Prüfung: 07. Juni 2019

Diese Dissertation ist auf den Internetseiten der Universitätsbibliothek online verfügbar

Berichte aus der Materialwissenschaft

Anoop Raghunath Kini

Laser Additive Manufacturing of

Oxide Dispersion Strengthened Steels

and Cu-Cr-Nb Alloys

“Study hard what interests you the most in the most undisciplined, irreverent and

original manner possible.”

― Richard Feynmann

Zusammenfassung

Die Laseradditive Fertigung (LAM) von metallischen Legierungen hat sich als vielversprechender

Fertigungsprozess mit geometrischer Gestaltungsfreiheit und hoher Produktivität von Bauteilen

etabliert. Die Domäne eröffnet vielversprechende Möglichkeiten als neuartiger Weg der

Materialsynthese, um Mikrostrukturen, die auch standortspezifisch in einer Legierung sein

könnten, zu maßschneidern oder sogar fortschrittliche Legierungen für LAM zu entwickeln. Zwei

weit verbreitete Bearbeitungswege in LAM sind das selektive Laserschmelzen (SLM) und das

Lasermetallabscheidung (LMD). Die vorliegende Arbeit untersucht Materialsynthesewege durch

LAM in zwei Materialsystemen: oxiddispersionsverstärkte (ODS) Stähle und Legierungen auf Cu-

Cr-Nb-Basis. Diese beiden Materialien sind für Hochtemperaturanwendungen (> 1000°C) wie

landseitige Gasturbinen bzw. mechanisch belastete elektrische Geräte von Interesse.

Erstens konzentriert sich die Arbeit an der LAM-Herstellung von ferritischen ODS-Stählen darauf,

eine mechanische Legierung von Ausgangspulvern auszuschließen, was ein kosten- und

zeitintensiver Prozessschritt ist. Hier versuche ich, Mischpulver (Oxid und Ferrit) zu erforschen,

mit der Absicht, die Marangoni-Konvektion in der Schmelze zur Dispersion von Oxidpartikeln der

zweiten Phase zu nutzen. Dies sind entweder Y2O3 oder La2O3, die in der vorliegenden Studie

verwendet werden. Die synthetisierten Materialien mit Y2O3 weisen einen signifikanten

Oxidpartikelverlust (0,3 Gew.-%) auf, verglichen mit dem ursprünglich zugegebenen (0,5 Gew.-

%). Die Yttriumoxid-Partikelagglomeration schreitet schneller voran als ihre Retention durch

schnelle Legierungsverfestigung während LMD und SLM. Im gefertigten Material mit La2O3 von

LMD wird die Dispersionshomogenität in der Legierungsmatrix für ein

Probenahmesondenvolumen in der Größenordnung von µm3 beobachtet. Bei Verringerung des

Probenvolumens für die Charakterisierung durch Rastertransmissionselektronenmikroskopie

(STEM) und durch Atomsonden-Tomographie (APT) bleibt die räumliche Homogenität jedoch

nicht erhalten. Die Vermeidung des mechanischen Legierungsprozesses scheint derzeit für die

ODS-Stahlherstellung nicht förderlich zu sein.

Im zweiten Teil der Arbeit schneide ich eine neuartige Mikrostruktur in einer konstruierten Cu-

3.4Cr-0.6Nb (at.%) Legierung, die durch LMD gehärtet wird. Die Mikrostruktur besteht aus

kohärenten Nano-Chrom-Präzipitaten, die in-situ (4 nm Durchmesser; Zahlendichte 8x1023 m-3) in

Verbindung mit den zuvor bekannten dispergierten Laves-Phasenpartikeln (< 1 µm) zum Härten

gebildet wurden. In-situ kohärente Fällung ist für Chrom innerhalb der breiten Klasse der Cu-

Basislegierungen bisher nicht bekannt. Die in-situ kohärente Niederschlagsmenge wird auf eine

synergetische Kombination der beiden folgenden Faktoren zurückgeführt. Erstens verhält sich das

Cu-Cr-Nb-System aufgrund des Chromgehalts der Legierung wie ein quasi binäres Cu-Cr-System.

Denn Chrom in der Legierung zur Fällung ist aufgrund seiner Unverwechselbarkeit in der Cu-

Basis unabhängig von dem in den Laves-Phasenpartikeln. Zweitens ist die prozessbegleitende

Abkühlrate während der LMD (103-104 K/s) geeignet, die Niederschlagsgröße auf die des

gewünschten kohärenten Regimes (< 10 nm) zu beschränken. Die Kohärenzhärtung trägt zu einem

signifikanten Wert von 78 Vickers Härte (Hv) bei. Ebenso beträgt der Beitrag 22 Hv für Laves-

Phasenpartikel, wie sie in der Orowan-Ashby-Formulierung vorhergesagt wurden. Die Summe

dieser Werte mit der Härte des Basiskupfers stimmt eng mit der gemessenen Materialhärte von

146 Hv überein; sie ist um 11% höher als die stärkste Cu-Cr-Nb-Legierung (Cu-8Cr-4Nb (at.%)).

Das 2D-Nanohärteprofil rechtfertigt die räumliche Homogenität der Aushärtung in der

hergestellten Probe.

Abstract

Laser additive manufacturing (LAM) of metallic alloys has emerged as a promising

manufacturing process featuring geometric design freedom and high productivity of fabricated

parts. The domain brings forth promising opportunities as a novel material synthesis route to

tailor microstructures which could also be in a site-specific manner in an alloy, or even design

advanced alloys suited for LAM. Two widely used processing routes in LAM are selective laser

melting (SLM), and laser metal deposition (LMD). The present work explores material synthesis

routes by LAM in two material systems; Oxide dispersion strengthened (ODS) steels and Cu-Cr-

Nb based alloys. These two materials are of interest in high-temperature applications (> 1000°C)

such as land-based gas turbines and in mechanically loaded electric devices respectively.

First, the work on LAM fabrication of ODS ferritic steels focusses on precluding mechanical

alloying of feedstock powders, which is a cost and time intensive process step. Here, I attempt to

explore mixed powders (oxide and the ferrite) with the intent of exploiting Marangoni

convection in the melt for dispersion of second phase oxide particles. These are either Y2O3 or

La2O3 used in the present study. The synthesized materials with Y2O3 is noted to undergo a

significant oxide particle loss (0.3 wt.%) compared to that initially added (0.5 wt.%). The yttria

particle agglomeration progresses faster than their retension by rapid alloy solidification during

LMD as well as SLM. In the fabricated material with La2O3 by LMD, the dispersion homogeneity

in the alloy matrix for a sampling probe volume is observed on the order of µm3. On decreasing

the sampling volume for characterization by scanning transmission electron microscopy (STEM)

and by atom probe tomography (APT), the spatial homogeneity however does not remain

preserved. Avoiding of mechanical alloying process step does not currently appear to be

conducive for ODS steel fabrication.

In the second part of the-work, I tailor a novel microstructure in a designed Cu-3.4Cr-0.6Nb

(at.%) alloy hardened by LMD. The microstructure constitutes nano-chromium coherent

precipitates formed in-situ (4 nm in diameter; number density 8x1023 m-3) in conjunction with the

previously known dispersed Laves phase particles (< 1 µm) for hardening. In-situ coherent

precipitation has not been known hitherto for chromium within the broad class of Cu-based

alloys. The in-situ coherent precipitation is attributed to a synergetic combination of the

following two factors. First, due to the alloy’s chromium content, the Cu-Cr-Nb system behaves

as a quasi-binary Cu-Cr system. This is because chromium in the alloy for precipitation is

independent from that in the Laves phase particles because of its immiscibility in the Cu-base.

Second, the in-process cooling rate during LMD (103-104 K/s) is appropriate to restrict the

precipitate size to that of the desired coherent regime (< 10 nm). The coherency hardening

contributes to a significant value of 78 Vickers hardness (Hv). Similarly, the contribution

amounts to 22 Hv for Laves phase particles as predicted by the Orowan-Ashby formulation. The

sum of these values with that of the hardness of base copper agrees closely with the measured

material hardness of 146 Hv; it is higher by 11% than the strongest Cu-Cr-Nb alloy (Cu-8Cr-4Nb

(at.%)). The 2D nano-hardness profile justifies the spatial homogeneity of hardening in the

fabricated sample.

Acknowledgements

It was a dream, since my Master’s student days to work at a place that harbored curiosity driven

research. I am profusely grateful to Prof. Dierk Raabe for this valuable opportunity to work for

this doctoral thesis at the Max-Planck-Institut für Eisenforschung GmbH. I am greatly inspired by

his ability to grasp multitude of topics in materials science and beyond, in addition to the way in

which he makes connections between them to synthesize new ideas. Each time I visited him at his

office, glad to also have noted, some of his habits and beliefs with which he leads the organization.

I believe both these aspects are the key takeaways in terms of learnings for the rest of my life.

I am highly thankful to Dr. Eric Jägle for supervising my work. His focus on laser additive

manufacturing resonated with my interests which prompted me to work for him. I am also thankful

the financial support from the AProLAM project, funded by the strategic collaboration between

the Fraunhofer Society and the Max Planck Society. This project proposal was initiated by Eric

whom I sincerely thank. From the project partner side at Fraunhofer-Institut für Lasertechnik,

would like to thank Dr. Andreas Weisheit, Mr. Markus Wilms and Ms. Dora Masichner for their

kind support.

I must greatly indebted to Dr. Baptiste Gault for enriching and insightful discussion on various

aspects on Atom Probe Tomography. His two decades of experience and the generated intuition is

truly commendable. I feel lucky to have met him. I also thank Dr. Dirk Ponge and Dr. Stefan

Zaeferrer for insightful discussions in their passionate fields on steels and electron microscopy

respectively.

I am grateful for technical support extended by Mr. Uwe Tezins and Mr. Andreas Sturm for the

atom probe tomography and the focused ion beam facilities. I would like to thank Ms. Monika

Nellessen and Ms. Katja Angenendt for the support with scanning electron microscopes. I greatly

appreciate and thank Ms. Heidi Bögershausen for nano-indentation and hardness measurements,

Mr. Daniel Kurz with ICP-OES chemical analysis and Mr. Benjamin Breitbach for XRD

measurements.

The local eco-system for interaction decides each person’s experience at the institute. Glad to have

identified people with whom I had lengthy, thought provoking and enriching conversations for

which I am grateful. They are Shyam Katnagallu, Alison Da Silva, Dr. Seok Su Sohn, Avinash

Hariharan, Supriya Nandy, Dr. Arka Lahiri, Ankit Kumar, Dr. Surendra Makineni, Aniruddha

Dutta, Viswanadh Arigela, Dr. Leigh Stephenson, Dr. Andrew Breen, Dr. Christoph Freysoldt, Dr.

Pratheek Shanthraj, Dr. Matthew Kasemer, Philipp Kürnsteiner and Priyanshu Bajaj. I am glad to

update new perspectives to my thinking via the interactions with them. I am glad that I have been

able to have made friends with many in the department including Yanhong Chang, Huan Zhao,

Wei Ye, Dayong An and many more.

I feel fortunate to have the strongest backing and unshaken support from my parents and younger

brother without which I could not have crossed the most challenging of times. It goes without

saying that many other people have contributed for this work, I will remain grateful to them.

Anoop Kini

[email protected]

mailto:[email protected]

List of Symbols and Abbreviations

Symbols Units

Hv Vicker’s hardness kg/mm2 or HV

Φ Laser beam diameter mm or µm

v Laser scan speed mm/s or m/s

𝑷 Laser power Watt

EV Volumetric energy density J/mm3

𝑯𝑺 Hatch spacing distance mm or µm

Δz Layer height mm or µm

G Shear modulus of matrix copper GPa

r Precipitate radius nm

f Precipitate mean volume fraction -

M Taylor's factor -

𝝂𝒑 Poisson’s ratio of precipitate -

𝜹 Misfit strain parameter -

aCu Lattice parameter of Copper Å

aCr Lattice parameter of chromium Å

Gp Shear Modulus of precipitate GPa

𝝀 Mean Particle Spacing nm

b Burgers vector for the copper matrix nm

Abbreviations

AM Additive Manufacturing

APT Atom Probe Tomography

EBSD Electron Backscattered Diffraction

EDS Electron Dispersive X-Ray Spectroscopy

FCC Face Centered Cubic

FEG Field Emission Gun

FIB Focused Ion Beam

GIS Gas Injection System

IPF Inverse Pole Figure

LAM Laser Additive Manufacturing

LEAP Local Electron Atom Probe

LMD Laser Metal Deposition

LOM Light Optical Microscopy

MA Mechanical Alloying

ODS Oxide Dispersion Strengthened

SEM Scanning Electron Microscopy

SLM Selective Laser Melting

STEM Scanning Transmission Electron Microscopy

XRD X-Ray Diffraction

Table of Contents

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

2. Literature Review ................................................................................................................................... 7

2.1. Laser Additive Manufacturing (LAM) .......................................................................................... 7

2.1.1. Laser Metal Deposition (LMD) ................................................................................................ 7

2.1.2. Selective Laser Melting (SLM) ................................................................................................ 9

2.2. Laser Additive Manufacturing for Fabrication of Functional and Structural Materials ....... 11

2.3. Immiscible Alloy Systems: Fundamentals and Examples .......................................................... 14

2.4. Alloy Design for Additive Manufacturing ................................................................................... 16

2.4.1. Oxide Dispersion Strengthened Steels (ODS) Steels ............................................................ 18

2.4.2. Cu-Cr-Nb Alloys ..................................................................................................................... 23

3. Experimental Methods ......................................................................................................................... 27

3.1. Additive Manufacturing ................................................................................................................ 27

3.1.1. Laser Metal Deposition (LMD) .............................................................................................. 27

3.1.2. Selective Laser Melting (SLM) .............................................................................................. 28

3.1.3. Powders for LMD and SLM .................................................................................................. 29

3.2. Microstructural Characterization ................................................................................................ 29

3.2.1. Inductive Coupled Plasma Optical Emission Spectroscopy (ICP-OES) ............................ 29

3.2.2. Optical Microscopy and Sample Preparation ...................................................................... 30

3.2.3. Scanning Electron Microscopy (SEM) and Electron Backscattered Diffraction (EBSD) 30

3.2.4. Atom Probe Tomography ....................................................................................................... 31

3.2.5. Transmission Electron Microscopy (TEM) .......................................................................... 32

3.2.6. Focused Ion Beam (FIB) Micromachining ........................................................................... 33

3.3. Mechanical Property Characterization........................................................................................ 35

3.3.1. Nano-indentation Testing ....................................................................................................... 35

3.3.2. Hardness Testing ..................................................................................................................... 36

4. ODS Steel Produced by Laser Additive Manufacturing ................................................................... 37

4.1. ODS Steels for Laser Additive Manufacturing ........................................................................... 37

4.2. Feedstock Powder Preparation ..................................................................................................... 39

4.3. Microstructural Characterization of Dense Samples ................................................................. 42

4.3.1. LMD of Ferrite Powders mixed with Yttrium oxide ........................................................... 42

4.3.2. LMD of Ferrite Powders mixed with Yttrium oxide: Yttria loss challenge....................... 45

4.3.3. SLM of Ferrite powders mixed with Yttrium oxide ............................................................ 52

4.3.4. LMD of Ferrite Powders mixed with Lanthanum oxide ..................................................... 55

4.4. Challenges and Comments ............................................................................................................ 58

5. Cu-Cr-Nb Alloy Designed for Laser Metal Deposition ..................................................................... 61

5.1. Alloy Design .................................................................................................................................... 61

5.2. Microstructural Characterization of Dense Samples ................................................................. 63

5.2.1. Dense Sample Fabrication ...................................................................................................... 63

5.2.2. Dispersed Laves Phase Particles ............................................................................................ 68

5.2.3. Chromium Nano-precipitates ................................................................................................ 71

5.4 Hardening Assessment .................................................................................................................... 72

5.4.1. Nano-indentation Measurements ........................................................................................... 72

5.4.2. Hardness Measurements ........................................................................................................ 75

6. Discussions ............................................................................................................................................. 79

6.1. ODS Steels ....................................................................................................................................... 79

6.1.1 ODS Steels containing Yttria particles ................................................................................... 79

6.1.2 ODS Steels containing Lanthana particles ............................................................................ 82

6.2. Cu-Cr-Nb Alloy .............................................................................................................................. 84

7. Summary and Concluding Remarks ................................................................................................... 91

Appendix 1 ................................................................................................................................................. 95

Appendix 2 ................................................................................................................................................. 97

References .................................................................................................................................................. 99

Curriculum Vitae .................................................................................................................................... 109

1

1. Introduction

Laser Additive Manufacturing (LAM) [1] of metallic materials has recently evolved as a promising

process for fabricating engineering components, even with those with complex geometries [2,3].

This is viable with feedstock powders or wires [4] and a Computer Aided Design (CAD) file

containing the designed three dimensional part geometry [2]. The layer-wise addition of the alloy

material is thus facilitated to build the component. Some of these components are presently serving

different industrial applications.

Industrial components fabricated by LAM are currently catering to various sectors of the global

economy; these include the energy (including industrial gas turbines) [5,6], the aviation sector

[1,6], the biomedical sector [3,7] as well as the civil engineering sector [1,6]. The materials for

such applications belong to a broad range; including steel, nickel alloys, aluminum alloys and

titanium alloys [5,7,8].

2

LAM has been gaining significant interest in academic research apart from that in industry

[1,9,10]. LAM presents opportunities for novel material development in mechanically loaded,

high temperature components, structural and biomedical materials.

The novel material development comprises manipulating the alloy microstructure specifically by

LAM in known alloys, or in newly designed alloy systems. For example, in designed titanium

based alloys particularly, Ti-xMo and Ti-xV (x = 0-25 wt.%) [11], novel microstructures with

finely precipitated α grains in large columnar β-Ti grains of size as large as 10 mm can be realized.

Such microstructures are expected to be beneficial for serving structural and biomedical

applications. This could sometime require a functionally graded microstructures [3].

Compositionally or functionally graded microstructures can be accessed more conveniently by

LAM than by other processing routes. This has been attempted also in other metallic systems for

example, Ni-Al, Ni-Ti [2], and Fe-Ni-Al based maraging steels are some examples highlighting

this plausibility of microstructural tailoring. On a similar note, site specific microstructural

manipulation, more specifically crystallographic grain orientation was shown to be feasible by

LAM in a nickel based superalloy [12]. Similarly site specific property design via alloy design by

LAM is being considered promising [13].

In a recent work, Harrison et al. [14] demonstrated exploring alloy design approach to address

segregation related challenges to reduce crack density in the microstructure of a nickel superalloy.

The present work focuses on design of alloys specifically for LAM in two materials systems; Oxide

Dispersion Strengthened (ODS) steels and Cu-Cr-Nb alloys.

First, in the ODS steel materials, a microstructure with fine (< 50 nm) and dispersed oxide particles

(Y2O3 ; 0.5 wt.%) is desired for high temperature applications particularly under creep loading

[15–17]. Conventional processing mandates component fabrication via mechanically alloyed

3

feedstock powders. Mechanical alloying is a cost and time intensive step which has limited the use

of such materials from being used more widely than that used currently; the materials are used

sparingly in gas turbines and heat exchangers [18–20]. To prevent the material from being extinct,

or even development of novel ODS steel materials, calls for alternative processing routes. Melting

based processing of such materials, despite being less laborious, is faced with undesired challenges

of oxide particle agglomeration or coarsening [21,22]. Recently, Walker et al. [16] and Boegelein

et al. [17] demonstrated the fabrication by LAM of ODS ferritic steels with Y2O3 as the oxide

particle chemistry.

The present work on ODS steel attempts to exploit Marangoni convection in the melt pool during

LAM, for oxide particle dispersion. The key processing goal aims to preclude mechanical alloying

process step which is cost and time intensive.

. The objectives of this work are the following:

1. Fabrication of ODS ferritic steels by LAM with non-mechanically alloyed feedstock

powders. Here, the atomized ferrite alloy powders are mixed with either yttrium oxide, or

with lanthanum oxide (0.5wt.%).

2. Achieving a microstructure with homogeneous oxide particle dispersion (<100 nm; 0.5

wt.%) in the ferrite steel matrix by optimization of LAM process parameters.

In the second part, design a lean Cu-Cr-Nb ternary alloy (< 6 at.% alloying) hardened by a novel

microstructure, feasible specifically by LMD is the aim of this work. The previously reported

microstructures in the hardened alloys viz. Cu-8Cr-4Nb and Cu-4Cr-2Nb (at.%), have relied

entirely on dispersed particles (< 1 µm) of Cr2Nb Laves phase for hardening. A maximum hardness

of 132 on Vicker’s scale was obtained [23–25]. Recently, LAM for manufacturing of functional

components with Cu-8Cr-4Nb (at.%) was demonstrated [23]. Here, an alternative microstructure

4

suitable for material hardening constituting chromium nano-precipitates although at the expense

of the volume fraction of previously reliant Laves phase, is presented.

The objectives of the work on Cu-Cr-Nb are the following:

1. The identification of a novel lean alloying regime which can favor hardening (> 130 Hv)

in this ternary system by exploring chromium nano-precipitates apart from the known

Laves phase dispersed particles.

2. Obtaining the above mentioned microstructure by identification of appropriate LMD

processing parameters.

The present work identifies a novel alloy, Cu-3.4Cr-0.6Nb (at.%) hardened specifically by LMD.

The key contribution of this work is a novel way of hardening the microstructure by in-situ

coherent Cr nano-precipitates, as well as by dispersed Laves phase particles. The substantial

hardening in the present alloy is comparable to that in known Cu-8Cr-4Nb and Cu-4Cr-2Nb (at.%)

alloys also belonging to such ternary system. The hardness measurements confirms the above,

while its spatial homogeneity attested by nano-indentation. Note that previous chromium

precipitation in copper based alloys like in the binary Cu-Cr systems has demanded at least an

ageing treatment if processed by rapid solidification [26], or frequently a prior additional

solutionizing step, also by conventional processing [27].

The layout of the thesis is presented in the following sequence. The Chapter 2 describes the

fundamentals on immiscible alloys while also reviewing the relevant literature on the two systems

belonging to this category, namely ODS Steels and Cu-Cr-Nb alloys. Methods describing

experimentation is delivered in Chapter 3. This constitutes both microstructural and mechanical

property characterization in addition to the parameters used for the LAM process. This is followed

5

by the characterization results in designed ODS steels with La2O3 as the oxide particle dispersion

chemistry in Chapter 4. In an analogous manner, the results on the Cu-Cr-Nb system are outlined

in Chapter 5. Chapter 6 discusses the underlying reasons for the novel microstructures in the

designed alloys; specifically the co-relation with the mechanical properties signifying alloy

hardening.

7

2. Literature Review

2.1. Laser Additive Manufacturing (LAM)

The capability of fabricating high performance metallic components with controllable

microstructures and mechanical properties is distinct across different LAM processes. The LAM

classification also suites the different mechanisms of laser-powder material interaction. These may

be bifurcated as laser metal deposition (LMD) and laser powder bed fusion which is also referred

to as selective laser melting (SLM) [2]. The following section detail the SLM and LMD processes.

2.1.1. Laser Metal Deposition (LMD)

Laser Metal Deposition (LMD) is a method to build engineering components by melting a surface

while simultaneously applying the metal powder. The melt pool is typically protected against

oxidation by supplying a shielding carrier gas, typically argon or helium. The powder is fed with


8

a co-axial multi-jet nozzle. Fig. 2.1 displays the schematic representation of the LMD process,

which mentions some constituent parts.

LMD provides a high build rate and permits for larger build volumes (up to 300 cm3/h), as opposed

to powder-bed based processes like SLM. This is dependent on the key laser processing

parameters; spot size, scan speed, and laser power. The spot size typically spans the range 0.3 - 3

mm; scan speed in the range 0.15-1.5 m/min. The layer thickness typically varies between 380 µm

and 1 mm [7,28].

Figure 2.1. A schematic representative image of the laser metal deposition (LMD)

process [29]

Rapid developments have resulted in evolution of several different systems for LMD. Most

commonly, the fabricated piece is stationary while the deposition head is repositioned for each

layer; for example by a 5-axis Cartesian-gantry system, or a robotic arm. In other systems, the part


9

is moved under a stationary deposition head. LMD is widely used for the production or repair of

turbine blades, shafts, and gear parts. The key materials serving these applications among others

are, Ni-based superalloys [5], Ti alloys, and steels [30,31]. Such materials are also being processed

by adopting another mechanism involving laser-powder interaction, i.e. by selective laser melting.

2.1.2. Selective Laser Melting (SLM)

Selective laser melting (SLM) is a powder bed-based build process in which the alloy powder is

spread as thin layers [7], that undergoes fusion onto the preceding layer underneath, post the laser

interaction. The typical layer thickness (Ds) can span in the range, 20-100 µm [7]. A recoating

system is frequently used for the powder distribution. Prior to this, the metal powder gets fed by a

hopper or by a reservoir which is located next to the work area. Fig.2.2. reveals the schematic

image of the SLM process.

The key process parameters varied are as follows; laser power (𝑃) in the range 20 W-1 kW, and

scan speed (𝑣) across the deposited powder layer of up to 15 m/s. The lasers typically used are

single mode fiber lasers in a continuous wave mode. The emitted radiation wavelength lies near

the infrared spectral regime with a wavelength of 1060-1080 nm. Typical spot sizes of the laser

beam in the focal plane is between 50 and 100 µm [28,32].

𝐸𝑉 =𝑃

𝑣 𝐻𝑆 𝛷 (2.1)

The equation (2.1) is an expression for calculating the volumetric energy density (EV) [2]; 𝐻𝑆

stands for the hatch spacing distance. The volumetric energy supplied to the powder layer causes

the exposed material powder to melt, while a part of it reaches areas which are adjacent to the melt


10

pool by heat conduction. During the melt pool solidification, the individual melt tracks solidifies

over the layer below. After the powder exposure to the laser beam, the build platform is lowered.

The next powder layer is deposited and the process of melting the newly deposited layer is

repeated, until part completion. Post fabrication completion, the unmelted powder can be sieved

and reused into the subsequent SLM process [7]. SLM is used for producing intricate shaped

components for medical implants [3], and jet engine parts [33]. A few key examples of components

serving biomedical and structural applications are listed in the following section.

Figure 2.2. A schematic representation of the selective laser melting (SLM) process

(taken from CustomPartNet) [34].

2.2. Laser Additive Manufacturing for Fabrication of Functional and Structural Materials

11

2.2. Laser Additive Manufacturing for Fabrication of Functional and

Structural Materials

The LAM produced parts are able to retain the functional properties of fabricated components. An

instance is that of the osseointegrative structures comprising the lattice structures, considered

suitable for implants materials. Another example is their use as structural components with

superior mechanical properties.

Figure 2.3. A knee implant (tibial stem) prototype developed for additive manufacturing

(AM) with a Ti-6Al-4V alloy (in wt.%) with acceptable . (a) X-ray image for female right

knee (total knee) replacement (femur, F; tibia, T). (b) and (c) Software model; the model

in (c) is rotated 45° relative to that in (b) about the axis of stem rod. (d) Corresponding AM

fabricated samples with increasing density of mesh array stems (from right to left

corresponding to 0.86, 1.22 and 1.59 g cm−3). Details taken from Ref. [3].

Fig.2.3 shows additively manufactured tibial–knee stems with various density compatibilities.

These are the high-end trabecular regime (approx. 0.8 g cm−3) and the low-end cortical bone

regime (approx. 1.5 g cm−3). Subfigure (a) shows an X-ray image with femoral (F) attachment

device and tibial (T) stem for a total knee replacement. It holds a highly cross-linked polyethylene

block in lieu of the meniscus, on which the femoral component rides. The subfigures (b,c) reveal


12

the three dimensional geometric model to be manufactured. The design is featured by two different

geometric symmetries within osseointegrative structures.

The alloy material in this implant is the Ti-6Al-4V (in wt.%). In Fig.2.3.(d), actual fabricated

prototypes with varying outer mesh densities (left to right) of 0.86, 1.22 and 1.59 g cm−3 are

illustrated. An enlarged inset for the 0.86 g cm−3 mesh is also shown. In brief these osseointegrative

structures are suited for producing by AM. Allied benefits of such intricate part fabrication lies in

obviation of complex machining operations.

The Ti-6Al-4V alloy components demand a tensile strength of about 800 MPa and an elongation

of 10%, particularly in structural applications serving aviation. Such components include landing

gears, brackets and other structural components. The geometry of these components may be

fabricated by LAM without compromising on the combination of mechanical strength and

elongation.

Figure 2.4. Ti-6Al-4V (wt.%) bracket for Airbus A320neo and A350XB with a topology

optimized bionic design resulting in a ~ 30% weight saving compared to the conventional

milled bracket [7,8].


13

Fig. 2.4 illustrates a three dimensional design of intricate geometries for LAM sometimes

associated with allied advantages of component weight saving. The other benefits can include

enhanced productivity and reliability. Such a component design promotes a weight saving of up to

~30%, than that of a conventional milled bracket. These functional components are a part of the

junction section between the wings and the engines [8].

Figure 2.5. Parts produced at General Electric Co. (GE) via laser additive manufacturing

(LAM). (a) A housing part for the temperature sensor at the compressor inlet inside the jet

engine. (b) Fuel nozzle for the GE9X jet engine. This is the largest jet engine ever built till

date [7,35].

On a similar note on intricate part fabrication, GE is also developing 3D-printed fuel nozzles,

displayed in Fig.2.5(b), among other parts for the GE9X engine. In the fuel nozzle, an integrated

design built as one piece contains optimized interior channels. The resulting weight saving is

approximately ~25% [7]. The engine is currently used for Boeing’s new 777X aircraft. The GE

9X would be the largest built hitherto [35]. LAM driven component design is exemplified also in

the aviation pylon brackets. This is by Airbus for their A320neo and A350XB carriers. Fig.2.4.

2.3. Immiscible Alloy Systems: Fundamentals and Examples

14

shows brackets produced with Ti-6Al-4V (in wt.%) alloy by LAM. This component features a

topology optimized bionic design. These LAM produced parts are already in service.

A study performed on Ti-6Al-4V performed by Thijs et al. [36] revealed that during rapid

solidification based SLM process, the microstructure constituted martensite. Depending on laser

process parameters influencing the heat input was shown to also result in Ti3Al precipitation. This

showed that novel microstructures can be obtained when processed by LAM. Rapid solidification

processing in alloy systems exhibiting immiscible behavior (phase separating system), have been

considered previously although not by laser additive manufacturing. For instance, the rapidly

solidified microstructures in Cu-Fe system was shown to result in an egg shaped morphology of

the separating Fe rich phase in the Cu alloy matrix [37]. The present thesis focuses on two

immiscible alloy systems produced by laser metal deposition, namely, oxide dispersion

strengthened (ODS) steels and Cu-Cr-Nb alloys. Preceding this, fundamental on the immiscible

alloys systems is detailed next.


Immiscible binary alloys are those in which the constituent elements have little equilibrium

solubility in one another [38]. This is a function of temperature, at which the solubility is noted at

an atomic length scale. Note that the free energy change for mixing (to form an alloy solid solution)

is positive in immiscible systems; ΔGm > 0. Such a behavior necessitates a positive free enthalpy

of mixing (or heat of mixing), ΔHm > 0 [38].

With increasing homologous temperatures of the bulk alloy however, the entropic contribution

(ΔSm) can dominate over enthalpic contributions. This can lead to spontaneous mixing, ΔGm <0


15

which holds at a temperature of critical unmixing, Tc, or higher [38]. A corresponding miscibility

gap can be noted in the phase diagram of binary systems for example in Ag-Fe, Cu-Fe [39], Al-Bi

[40] among a few other immiscible systems.

The immiscible systems are favored thermodynamically to phase separate (T < Tc), even if retained

as a single phase regime. Specifically, the kinetics of diffusion mediating the phase separation is

dependent on temperature and time, which decides the plausibility of single phase retention which

competes against equilibrium phase formation. Different non-equilibrium processing methods may

be viable for obtaining a single phase in the microstructure. Some processing routes are liquid

quenching (LQ), vapor quenching (VQ) and mechanical alloying (MA). Note that VQ-S (LN2T)

stands for vapor quenching with liquid nitrogen.

A list of few immiscible binary systems, is presented in Table 2.1. These binary alloys are

designated as AxB100-x (where, x is in at.%) with the corresponding non-equilibrium processing

technique specified against the alloy [37].

Table 2.1. Experimentally observed binary alloy systems, designated as AxB100-x with a positive enthalpy of

mixing. The composition range (x is in at.%) exhibiting a single phase regime (prior to phase separation

via ageing) are listed. The corresponding non-equilibrium processing route is specified against the alloy.

AxB100-x x (in at.%) Processing Technique

Ag-Cu 0-100 LQ[41] and MA [42]

Cu-Cr 0-100 VQ-S (LN2T) [43]

Cu-Nb 35-74 VQ-S (LN2T) [44]

Fe-Cu 0-100 VQ-S (LN2T) [45]

Fe-W 20, 50-80 MA [46]

Mg-Ti < 12.5 MA [47]

Some applications of the afore-listed immiscible alloy systems are as follows. The Cu-Ag and Cu-

Nb system, which are used for growth of epitaxial metal films [48]. The latter among the two

systems is also essential in superconductors, and parts for robotics [49]. The Fe-Cu based alloys

2.4. Alloy Design for Additive Manufacturing

16

are widely used in automotive, electrical, and industrial machinery [37]. The Fe-W based alloys

find applications for high temperature environments [50,51]. The Mg-Ti immiscible system has

favorable biocompatibility with the human body and has been widely accepted as orthopedic

biomaterials and subchondral bone replacement materials [52]. These examples on immiscible

alloy systems are associated with a variety of microstructures achieved by non-equilibrium

processing routes including rapid solidification processing. This thesis focusses alloy systems

designed for processing by rapid solidification based laser additive manufacturing in two systems

exhibiting immiscible behavior.


The additive manufacturing domain presents opportunities for the development of new and

advanced alloys for different functional applications. Specifically, this permits the access to novel

microstructures specifically accessible by LAM. In a recent work on the design of titanium based

alloys, specifically, Ti-xMo and Ti-xV (x = 0-25 wt.%) [11], a microstructure with finely

precipitated α grains in large columnar β-Ti grains (10 mm) was demonstrated. Incipient

directional solidification during additive manufacturing was the attributed cause for the large

columnar grains. The microstructures are expected to be beneficial for serving structural and

biomedical applications. Note that the microstructures can also be compositionally graded.

Compositionally graded microstructures in a Ni-Al system was shown via in-situ synthesis using

elemental powders during additive manufacturing. A microstructural gradation via Ni, Ni/Ni3Al,

Ni3Al, Ni3Al/NiAl, NiAl was demonstrated. A graded microstructure in a binary Ti-Ni (Ti= 0-23

at.%) was successfully revealed constituting a phase evolution of α, α+β, α+β+Ti2Ni, and β/B2 +


17

Ti2Ni. A novel morphology of β/B2 + Ti2Ni accompanied a coupled eutectic growth which was

anomalous till then [53]. In a graded Fe-19Ni-xAl (x=0-25 at.%) maraging steel system, massive

nanoprecipitation of NiAl in a matrix of ferrite/martensite was demonstrated in a recent research

work. The exceptionally high precipitate number density of 1025 m-3 of NiAl precipitates of size

2-4 nm [54] was noted in this ternary system designed for LAM.

On a similar note, site specific microstructural manipulation of crystallographic grain orientation

was shown to be feasible by LAM in a nickel based superalloy [12]. Analogously for site specific

property design, the approach of alloy design via LAM is being considered promising [13],

accompanying secondary benefits of averting laborious processes which lower productivity.

In a recent research work, Harrison et al. [14] demonstrated exploring alloy design approach, by

specifically altering Mn and Si content, to resolve the key challenge of cracks in a Hastealloy-X

grade of nickel superalloy component. The authors controlled the tramp element amount,

constituting O, N, P, Cu, and Pd. These were expected to be deleterious, as these elements were

considered likely to favor crack formation even for elemental amounts not exceeding a few ppm.

This factor coupled with an increased solid solution strengthening reduced crack density by 65%.

More importantly, this investigation and inference ruled out the previous hypothesis which claimed

grain boundary segregation as the governing mechanism in cracking phenomenon in this alloy.

The above instances attest that ability to access novel microstructures via and designed alloys. The

present work explores the design of alloys is in two material systems featuring immiscibility,

namely the oxide dispersion strengthened (ODS) ferritic steels and the Cu-Cr-Nb based alloys.

The microstructures in these materials produced by conventional processing is described in the

following.


18

2.4.1. Oxide Dispersion Strengthened Steels (ODS) Steels

ODS steels are materials used for high temperature applications beyond 600°C [15,16,55] under

creep loading. These are materials essential in land based gas turbines and heat exchangers. The

microstructures in ODS steels comprises ferritic matrix [18–20] or its combination with martensite

[15], with grain size less than a µm (Fig. 2.6(a)) after an extrusion process [20,56]. The fine oxide

dispersion particles (< 50 nm; Y2O3 [57] ; 0.5 vol.% [15,56,57]) serve as the second phase

strengtheners. These constitute the key features in the desired microstructure for serving high

temperature applications [15–17].

Note that the Iron-Yttrium (Fe-Y) binary system near the Fe rich compositional domain (< 5 wt.%

Y), exhibits immiscible behavior between the two elements [58,59]. The processing of materials

constituting Fe and Y require non-equilibrium methods specifically mechanical alloying (MA), by

which these oxides are incorporated into the matrix as a solid solution. A subsequent ageing step

results in spatially homogeneous precipitate distribution with a number density exceeding a value

of 1023 m-3 (Fig. 2.6(b)) [18,19,56].

The alloying elements chromium and aluminum that are added to the iron base favor the matrix

phase to be ferrite or its co-existence with the martensite (at a functional temperature of 600°C

[15]). The alloying elements also are critical in providing oxidation and spalling resistance [15,20]

in high temperature environments. This particle chemistry choice is based on the microstructural

stability at high temperatures which is further aided by titanium addition to the alloy base. Titanium

(> 0.3 wt.% [60]) refines these oxides to a size of 1-2 nm [57] (Y-Ti-O complex oxides) and

enables grain boundary pinning effect [18,56] while remaining resistant to coarsening even at 1000

°C for prolonged exposures [18–20].


19

Fig. 2.6. Microstructure of ferritic ODS Steel. (a) TEM image revealing dispersed oxide

particles in the ferrite matrix [20]. (b) Atom probe tomography results revealing a high

number density of 1023 m-3 of Y-Ti-O rich oxides [18].

The specific oxide chemistries are Y2TiO5 and Y2Ti2O7 [61,62] precipitated upon annealing of the

consolidated powders. This is after the mechanical alloying step, essential in the route for

feedstock powders for industrial component fabrication [18,57]. The mechanical alloying process

step, however, is a cost and time intensive process and often bottlenecks the scaling up the

production of ODS Steel materials [63]. On the contrary, conventional melting based

manufacturing processes has been known to not favor processing of oxide dispersion strengthened

(ODS) steels. Such melting based processes have been known to result in agglomeration,

coarsening of oxide particles [21,22], and/or inhomogeneous particle distribution [64] including

instances of slagging [16]. The fabrication of ODS steels, therefore, has been drawing alternative

processing routes which do not necessitate mechanical alloying process step completely or

partially. From an alloy design perspective these introduce novel oxide particle chemistries.


20

Tang et al. [65] synthesized ferritic ODS steels with Ti3O5 oxide particle chemistry of size ~ 5 nm,

by in-situ oxidation. No mechanical alloying step was needed. The in-situ oxidation was

demonstrated by a combination of dispersion supply using a titanium in wire form in combination

with electromagnetic stirring of the liquid metal. However, it was not clarified if the oxide

chemistry supports creep loading at high temperatures of 1000°C [18–20].

Recently, use of MnCr2O4 as oxides particles for dispersion strengthening was shown by obviating

mechanical alloying [66] during laser metal deposition. The material matrix belonged to a 316 L

grade of stainless steel realized by controlling the partial pressure of oxygen and nitrogen in the

atmosphere during deposition. This clearly demonstrated the dispersion at room temperature,

although the rationale for the oxide choice was not addressed in addition to their suitability for

high temperature exposures.

In an analogous route involving laser melting but coupled with controlled oxygen partial pressure

during the synthesis was reported of for an ODS Steel [67]. The authors revealed the possibility to

finely disperse SiO2 oxide particles in the steel matrix. While the tensile and yield strength

respectively was shown to be 703 MPa and 456 MPa respectively. However, this was under room

temperature conditions. Hoffmann et al. [68] proposed alternative oxide chemistries like MgO,

CeO2, ZrO2. Although these were conceived to be potential candidates on cost considerations, they

did not outperform Y2O3 in the high temperature tensile behavior and microstructural stability. The

authors devised a thermodynamic criteria for comparing Gibbs free energies of different oxide

chemistries. It must be noted that the free energies were compared between oxide materials, but

not in the presence of a steel matrix. The latter prevails under service conditions of ODS materials.


21

Recently, La2O3 as a candidate oxide chemistry was proposed by Pasebani et al. [69–71]. The

conveyed rationale, was based on its promising chemistry in other high temperature materials,

specifically molybdenum alloys. The validity of the argument remains to be addressed specific to

ODS steels. Alternatively, the high temperature creep property measurement was not reported to

justify the choice of this oxide chemistry. The work by the authors also proved the possibility to

achieve a high oxide number density ~ 3.7x1024 m-3 in the feedstock powder [70]. In consolidated

material prepared by spark-plasma-sintering (SPS), the oxides enriched in La-Ti-Cr of diameter 2-

70 nm was proved by HR-TEM [71]. Note that the feedstock powders were mechanically alloyed

[70,71].

In brief, the rationale for the choice of oxide chemistry in ODS steel materials for high temperature

applications, requires systematic and detailed investigation. A comprehensive understanding

towards specificity of oxide chemistry selection is therefore essential, or alternatively their

suitability in terms of measured creep properties of the ODS materials produced with such oxide

chemistries.

LAM has recently displayed the possibility to fabricate ODS steels [16]. This has been with Y2O3

as particle chemistry, requiring mechanically alloyed powders with ferrite alloy base, compatible

to the PM2000 grade [16,17,72]. ODS steel material microstructure contained fine (< 50 nm) and

uniformly dispersed oxide particles (Y2O3 ; 0.5 vol.%) [15–17] shown in Fig. 2.7.


22

Fig. 2.7. As-SLM produced sample of a PM 2000 ferritic steel produced by selective laser

melting (SLM) process, shown in the a dark field TEM image from Ref.[17].

The fine melt pool size minimizes the risk of oxide agglomeration and could support oxide particle

dispersion due to the effect of the Marangoni convection [16,17]. The effect is dominant at fine

melt pool sizes like in LAM compared to casting. No previous work hitherto suggests the

possibility to avoid mechanically alloyed powders for LAM.

In the present work, we intend to address two aspects. First, in view of alloy design, we propose a

systematic criteria to the field of ODS steel materials specifically for oxide particle chemistry

selection. Second, if the identified candidate oxide particles can indeed be dispersed by relying on

the effects of Marangoni convection which then could obviate mechanical alloying process step.

Analogous to the ODS steels which belongs to the immiscible alloy system, this thesis also

discusses on another such system, specifically Cu-Cr-Nb.


23

2.4.2. Cu-Cr-Nb Alloys

Copper alloys are essential for applications requiring high conductivity in combination with

mechanical strength. The materials are critical for high-field pulsed magnets [49,73–75], Bitter-

type direct current magnets [76], aerospace cables [77], rocket nozzle liners [23,78,79], robotics

parts [49,73,80], circuit breaker parts, and switchgears [81,82].

One such alloy class is the Cu-Cr-Nb ternary system. In the designed alloy, Cr and Nb are

immiscible in the copper base [25,83], up to the melting temperature of the bulk copper alloy. The

two key roles of alloying elements are as follows; firstly, the elements combine to harden the alloy

via dispersed particle strengthening via Cr2Nb Laves phase (size < 1 µm) [25], shown in Fig. 2.8.

Secondly, these also permit the matrix copper to remain nearly pure and conductive, on account

of their poor solubility in copper [25,83]. Examples of materials belonging to such class, are the

Cu-(4/8)Cr-(2/4)Nb (at.%) alloys referred to as Glenn research Copper (GRCop)-series [23–25].

The alloys possess a competitive combination of thermal conductivity of 82% IACS (315 Wm-1K-

1) [23] and a tensile yield strength of (~ 300 MPa) at room temperature [24] which are functional

in the NASA’s rocket engines [78].

It must be noted that the alloys that bear a ratio of Cr and Nb (in at.%) to be 2:1, the hardening is

relied solely upon Cr2Nb dispersed particles [23–25]. A deviation from the alloying ratio, for

example a lean Cu-2Cr-0.5Nb (at.%) [84] in a melt-spun ribbon form, has not been further studied.

However, its microstructure constituted incoherent Cr precipitates, apart from a low volume

fraction of Laves phase particles than in the GRCop-alloy series. However, their work does not

substantiate if the overall alloy hardening is indeed substantial; more specifically, if the the low

volume fraction of previously reliant hardening phase Cr2Nb can be compensated by the incoherent

Cr precipitates, in terms of hardening.


24

Fig. 2.8. Microstructure of Cu-8Cr-4Nb (at.%) alloy constitutes grains of 2.7 µm in diameter) and

dispersed Cr2Nb Laves phase particles [83].

To the authors’ knowledge, no other work explores on developing Cu-Cr-Nb alloys for hardening.

This has been jointly because of the above reason pertinent to hardening, and also because of the

unfavorable high melting temperature of Nb, and its reactivity/poor oxidation resistance. Moreover

Cu-Cr-Nb component fabrication necessitates an elaborate machining as well as joining process

like friction stir welding [24]. To address such issues, a more acceptable manufacturing process as

well as the design of new Cu-Cr-Nb alloys which complement manufacturing, are important. As

these problems are being collectively addressed by rapid solidification processing based additive

manufacturing [28,85,86], the Cu-Cr-Nb ternary space can be revived for future alloy

development.


25

For design of the alloys, among Cu-Cr based systems, Nb bears other advantages over other ternary

elements, such as Ag and Zr. First, Nb has a greater relative abundance in terms of the availability

in nature [87]. Second, lower price (per weight); Nb is merely 1/4th as expensive as Ag and 1/10th

as expensive as Zr.

Laser additive manufacturing (LAM), a rapid solidification processing method [1], is currently

promoting endeavors to design new alloys [13,14]. The key underlying driver has been the

possibility to attain high in-process cooling rates (>103 K/s) [7,54,88]. LAM has lately led the

simplification of the manufacturing process chain, as exemplified in GE’s aviation fuel nozzles

[35]. Such simplifications, offer the prospects to cut down the complicated manufacturing process

steps like machining and joining [35]. In connection with the challenging manufacturing of Cu-

Cr-Nb alloys, LAM could prove to be an ideal route. The recent scrutiny of manufacturing process,

has certified LAM to produce functional components with Cu-8Cr-4Nb (at.%) alloy [89].


26

27

3. Experimental Methods

3.1. Additive Manufacturing

3.1.1. Laser Metal Deposition (LMD)

The alloy samples by LMD were prepared at the Fraunhofer-Institut für Lasertechnik (ILT),

Aachen. This was using a 5-axis handling system. It was equipped with a fiber coupled diode laser,

LDM 3000-60 from Laserline (Laserline GmbH Mülheim-Kärlich, Germany). The laser

wavelength was specified as 976 nm with a laser beam diameter of 0.6-1.8 mm. The maximum

laser output power was specified to be 3 kW. Note that the beam diameter was achieved as a

resultant of two lenses namely, the collimation lens and the focusing lens.

The interaction of the laser beam with the blown power feedstock led to the melt pool formation,

deposited above the layer located underneath. In the same manner, the deposition was continued

along the entire scan track length. The scan tracks located adjacently were spaced by a hatch

3.1. Additive Manufacturing

28

distance and a layer height distance was maintained between successive deposited layers. The first

deposited layer was over the substrate material.

For ODS steel sample fabrication in particular, the power was varied in the range 600-1250 W.

Contingent on the power, the laser scan speed was varied in the range 600-2000 mm/min. The

track offset was altered across 500-1100 µm and laser height between 300-700 µm.

For Cu-Cr-Nb alloys manufacturing, the laser beam diameter was 1.8 mm. The parameters were

set to be 1.5 kW and 600 mm/min for sample fabrication. The track offset was 800 µm and a layer

height specified was 200 µm. The chosen combination of high laser power and a low scan speed

is essential to melt the Cu based dilute alloy which demands high energy density. This is necessary

to compensate for the energy loss because of their high reflectivity, specific to pure copper or its

lean alloys [90].

3.1.2. Selective Laser Melting (SLM)

Experiments with selective laser melting (SLM) were performed using an Aconity MIDI machine,

at ILT Aachen. The laser beam source contained was an IPG fiber-laser with a specified maximum

laser power of 1 kW. Argon gas was chosen as the shielding gas with a low oxygen level of 100

ppm, during sample fabrication.

The laser power for sample fabrication was varied across 140-200 W; a laser scan speed of 600 -

1200 mm/s. The layer height for the samples was maintained to be 30 µm and a track offset of 60

µm. Note that SLM was performed on ODS steel samples only. This was to understand the oxide

amount retention in the fabricated sample by refining the size of the melt pool to achieve high

cooling rate compared to that in LMD. A detailed explanation follows subsequently.

3.2. Microstructural Characterization

29

3.1.3. Powders for LMD and SLM

The feedstock powders for LMD were prepared in a planetary ball mill by mixing at ILT Aachen.

The mixing was conducted for a duration of 4 h in total. The abrasive milling balls were yttria

stabilized zirconia (YSZ) lined. These along with the powder were placed in a milling container,

with a respective weight ratio maintained to a value of 10:1. Note that the container was sealed in

an Argon inert gas atmosphere. A rotational speed of 200 RPM was maintained during the entire

duration of milling operation.

For serving as a base case for comparison, a mechanically alloyed powder supplied from Plansee

AG, Austria was considered. It contained 0.5 wt.% yttria in a ferrite steel matrix although with a

composition Fe-26Cr-2Mo (in wt.%) as verified by ICP OES chemical analysis.


3.2.1. Inductive Coupled Plasma Optical Emission Spectroscopy (ICP-OES)

An absorption based optical emission spectroscopy based method was used for bulk chemical

compositional analysis. Specifically, it involves wet chemical analysis measured using the ICP-

OES method. An Optima 8300 model supplied by Perkin Elmer instruments was used for this

purpose. The measurements were performed at the Department of Interface Chemistry and Surface

Engineering at the Max-Planck-Institut für Eisenforschung GmbH.


30

3.2.2. Optical Microscopy and Sample Preparation

For microstructural characterization, the alloy samples were embedded in a conductive phenolic

resin with a carbon filler (from Struers Inc.). This was suited for a subsequent SEM examination.

The resin curing was performed at a temperature of 180°C for a duration of 7 min.

The sample surface preparation included grinding and polishing. Polishing was performed using a

1 µm diamond suspension, eventually followed by a 50 nm aqueous silica suspension. Further, an

eventual vibro-polishing cycle in colloidal silica medium (20 nm) MasterMetTM 2 from Bühler.

During this chemo-mechanical means of polishing, the samples oscillated on the wheel almost

horizontally under frequency of 120 Hz and an amplitude initially set at 70% of the prescribed

maximum. The procedure was conducted for a duration of at least 3 h.

The resultant finished surface was amenable to microstructural observation by light optical

microscopy (LOM), but also by scanning electron microscopy (SEM). Specifically for the powder

samples, its characteristics in terms of shape and size distributions were analyzed. These were

quantified using an open source image processing software, ImageJ.

3.2.3. Scanning Electron Microscopy (SEM) and Electron Backscattered

Diffraction (EBSD)

The sample preparation and analysis for as-LMD produced samples was identical to that for the

powder samples. SEM characterization was performed on a Zeiss Merlin system equipped with a

field emission gun (FEG) and a JEOL-6500F model from JEOL Ltd.

The probe current for image acquisition was at-most 4 nA, while the accelerating voltage was set

at 15 kV. The selected parameters suites for performing chemical compositional analysis by EDS.


31

Elemental mappings by EDS were performed with a Bruker Quantax system with the Zeiss-Merlin.

The results analyzed were using the Esprit software version 2.1. An EDAX system was used for

EDS analysis integrated with JEOL-6500F system with an analogous set of parameters as that with

the Brucker system, while the probe current was set to a value of 15.

For electron back scattered diffraction (EBSD) TSL OIM system, integrated to a JSM-6500F

microscope from JEOL Ltd was used. The measured EBSD data acquired was at an accelerating

voltage of 15 kV. The principle behind EBSD is explained in the Ref. [91].

3.2.4. Atom Probe Tomography

For spatial and chemical resolution at an atomic length scale, characterization was performed by

atom probe tomography (APT) equipped with a local electrode (LE). APT is a time of flight mass

spectroscopy technique via a destructive material characterization. LE is known for enhancing the

signal to noise ratio.

APT measurements were performed using LEAP® 5000 XS system supplied from Cameca

Instruments Inc. The ion flight path trajectory post field evaporation was straight path in this

equipment, with a specified detection efficiency of 80%. A laser source was used for assistance of

evaporation process.

In other words the samples were run in laser evaporation mode with a laser pulse energy of 40 pJ,

laser diameter of 1 µm and a wavelength of 355 nm which lies in the UV spectral regime. For all

these measurements, voltage pulse frequency was set at 200 kHz with pulse height being 15% of

the applied voltage. The detection rate was set at 0.25%, equivalent to 5 ions per 2000 pulses. The

base temperature for these trials was maintained at 80 K. Sufficient time was given for temperature


32

attainment before beginning each APT run. After finishing each run, the acquired data was

reconstructed and visualized using the IVASTM software version 3.6.12.

The principle behind the APT exploits field evaporation to successfully remove atoms from the

apex of a needle-shaped specimen. Field evaporation involves ionization of the surface atoms

whereby they are subjected to an electric field force which causes them to accelerate towards a

detector under a particular projection. The evaporation event follows immediately after the

ionization of the surface atoms. The ionization is induced by a combined effects of a standing DC

electrostatic field and a high-voltage or laser pulses that are transmitted to the surface atoms in the

specimen. Depending on the location on the detector where each ion hits post evaporation, the

material can be reconstructed and visualized [92].

3.2.5. Transmission Electron Microscopy (TEM)

Transmission electron microscopy (TEM) was performed using a Phillips CM-20 analytical

microscope operated at 200 kV. The assembly permitted a sample tilt of ±30° with reference to

the incoming beam axis. Note that the alloy sample was prepared using focused ion beam (FIB)

technique for micro-machining which was supported by a copper grid.

Scanning transmission electron microscopy (STEM) was performed with a JEOL-2100F

microscope, with the possibility to operate either in transmission mode or scanning mode. The

operating accelerating voltage was 200 kV. The microscope was equipped with a Gatan bright

field and dark field detectors. A sample tilt of ±21° was possible. A vacuum in the specimen

chamber was maintained at or below 2x10-7 mbar.


33

3.2.6. Focused Ion Beam (FIB) Micromachining

The sample preparation for APT and TEM was carried out using focused ion beam (FIB) micro-

machining. The FIB is equipped with a dual beam consisting of a column for high resolution

SEM. Helios 600 from FEI Company was employed for sample preparation. A sample for APT

with a needle like shape to promote electric field assisted evaporation process was prepared [93].

Figure 3.1. SEM image of a needle shaped sample prepared for atom probe tomography

(APT) by focused ion beam (FIB).

The APT sample preparation was begun by cutting out a triangular prism shaped trench from

the alloy. Subsequently, it was then lifted-out and transported to the location of the silicon based

micro-post array by means of tungsten manipulator. A part of the trench is affixed onto each

post, by platinum deposition by deploying the gas injection system (GIS). The final needle like

shaped-samples were prepared by annular milling until the desired tip geometry was obtained,

with a radius less than 50 nm and a shank angle of close to 20°. Gallium contamination in the

alloy which could have entered during the course of prior milling steps, was removed by a final


34

cleaning operation performed at 5 kV and 8 pA. A picture of the sample tip for APT is shown

in Fig. 3.1.

Figure 3.2. TEM lamella preparation. (a) lifted-out material deposited onto a copper grid.

(b) Different stages of a lamella thinning to produce an electron transparent lamella.

3.3. Mechanical Property Characterization

35

The sample for TEM on the other hand, must be a thin lamella (< 100 nm) which is electron

transparent. The lamella preparation involved thinning progressively along near parallel direction

(within ± 3°), until the final thickness was reached. This is shown in Fig 3.2.


3.3.1. Nano-indentation Testing

Tests were performed using a Hystrion Tribo Scope 950 nano-indenter system. It consisted of

a piezo-scanner, a transducer, apart from a 3-sided pyramidal diamond Berkovich indenter. The

indenter was positioned perpendicular to the sample surface. This was possible by controlling

the stage movement along X and Y axes with a possibility of movement also along Z axis in

the system. Acquisition of scanning probe microscopy (SPM) images, post indentation was

possible.

Load-controlled mode was chosen for carrying out the nano-indentation trials with a maximum

constant load of 5000 µN, controlled using a piezo-actuator. For loading-unloading, a trapezoid

profile was used with a holding time set to 5 s under the maximum load condition. The load-

displacement data were measured for each of the indentations in the 10x10 array in the

representative microstructural region.

For calculating the hardness and modulus values from the measured load-displacement curves,

it is critical to know the tip geometry precisely. The indenter tip was calibrated for area function

for this reason, using a quartz crystal as the standard reference. To clarify the effect of the grain

orientation on the hardness values, the inverse pole figure (IPF) maps were obtained. This was


36

done in the region containing the indents by carrying out electron back scattered diffraction

(EBSD).

3.3.2. Hardness Testing

In order to gauge hardening over a larger length scale than that by nano-indentation, Vickers

micro-hardness was also tested. The apparatus used was from LECO Instruments (AMH-43),

equipped with a diamond pyramidal indenter. A load of 500 g was used and at least a set of 12

indents were carried out on each sample for determining the hardness.

37

4. ODS Steel Produced by Laser Additive Manufacturing

The chapter presents on ODS steels produced via LAM by precluding the powder mechanical

alloying process step. Two known oxide particle chemistries i.e. Y2O3 and La2O3 are considered

for ODS steel synthesis with laser metal deposition (LMD). The results on ODS steels with Y2O3

produced by also selective laser melting (SLM) is detailed. The results are preceded by an excerpt

on the choice of Y2O3 and La2O3 as oxide particle chemistry.

4.1. ODS Steels for Laser Additive Manufacturing

The oxide particle chemistry for ODS steel fabrication by LAM in the current study are either with

yttria (Y2O3) [57] or lanthana (La2O3) [70]. Previous reports on oxide chemistries like Ti3O5 [65],

MnCr2O4 [66], SiO2 [67], MgO, CeO2, and ZrO2 [68] has not justified their suitability under high

temperature conditions.

4.1. ODS Steels for Laser Additive Manufacturing

38

For creep resistance in ODS steel materials, it is imperative to identify oxide chemistries which

support resistance to coarsening in the steel matrix. The fine and dispersed oxide particles are

essential for grain boundary pinning and thereby contribute to creep strength and low plastic strains

[55–57].

First, the oxides must have a high melting temperature with reference to that in the steel matrix;

more generally, the oxides must not undergo any phase transformation including allotropic ones

for a temperature of up to 1000°C [94] from room temperature. The oxides of following elements

viz. Sc, Sm, Tb, Tm, Th, Zr, La, and Y can be shortlisted from the Ellingham’s diagram [95].

Second, for favoring coarsening resistance of the oxide particles from a thermodynamic viewpoint,

the solubility of the element that constitutes in the oxides must be low in the steel matrix. The

solubility must remain low even at high service temperatures for such ODS steel materials. To

ensure that no solubility of the element (constituting the oxides) exists in Fe (principal element of

the matrix). Binary phase diagrams of Fe with each of these elements are therefore, examined. The

elements Sc, Sm, Tb, Tm, Th, Zr have solubility in the iron matrix at temperatures of up to 700°C;

the elements. These elements are dissolved in the matrix could subsequently aid oxide coarsening

by undergoing diffusion. The elements La and Y do not have solubility in iron matrix even beyond

800°C [39].

Third, large atomic size of such elements constituting the oxides (w.r.t. the principal matrix

element Fe) or low diffusivities in the steel matrix is essential which appends the previous criterion.

Considering this factor kinetically inhibits the oxide particles from coarsening. The coarsening

phenomenon is driven by the particle size dependent chemical potential gradient, also referred to

as the Gibbs Thompson effect. In comparison to the atomic radius of 1.40 Å for Fe which is the

matrix element, La and Y are significantly larger with radii of 1.95 and 1.80 Å respectively.

4.2. Feedstock Powder Preparation

39

Consequently, oxides of La and Y are considered suitable and chosen for the present work. Note

that the base alloy composition of the matrix is Fe-20Cr-5.5Al-0.5Ti (wt.%) which is compatible

with that of the commercial PM 2000 grade used for ODS steel. The alloy is in powder form used

as feedstock for LMD.

Figure.4.1. Pictorial representation of the criteria for oxide particle chemistry selection for

design of oxide dispersion strengthened (ODS) steels.


The yttrium oxide (yttria) or lanthanum oxide (lanthana) in powder form is mixed with the ferritic

alloy powder. The mixing is by means of milling for a short duration of 4 h (mixed powders)

performed at Fraunhofer ILT, Aachen. The approach differs from that used in conventional

processing which necessitates powders that are milled for a long duration of 80 h, until mechanical

alloying [22,96]. Fig.4.2 (a) and (b) display the X-ray diffraction pattern of yttria and lanthana

powders. The crystallite sizes (average) are estimated to be 15 nm and 30 nm respectively,

calculated using the Scherrer’s approximation [97]. This is based on peak broadening data


40

acquired from XRD. Note that the instrumental peak broadening contribution has not been

considered.

Figure 4.2. X-ray diffraction plot of oxide powder before mixing with ferrite powders. The

calculated crystallite sizes are 15 nm and 30 nm correspondingly for yttria and lanthana.

The crystallite sizes are calculated using the Scherrer’s approximation.

The SEM micrograph in Fig.4.3 indicates ferrite powders and deposition of oxide particles onto

them. The chemical composition of mixed powders (milled for a low duration 4 h) was determined

by ICP-OES chemical analysis. A 0.39 wt.%Y (0.5 wt.% of yttria), and 0.44 wt.% La (0.5 wt.%

of lanthana) was measured which is acceptable. The mixed powders are in majorly spherical in

shape subsequently serve as the feedstock for the LMD process.

The objective in the current approach, lies is exploring homogeneous oxide dispersion, but by

avoiding mechanical alloying process step which has been necessary hitherto. The underlying

reason for attempting this approach is exploiting temperature dependent surface tension driven

Marangoni convection coupled with high cooling rate [16,17]. It must be noted that the convection


41

effect becomes more predominant with a fine melt pool size. Here, in the LMD process the melt

pool width ca. 2 mm, contingent on the laser processing parameters.

Figure.4.3. SEM micrographs of powder particles. (a) Ferrite powders produced by

atomization with near-spherical shape. (b) Successful deposition of oxide (0.5 wt.% added)

onto ferrite powders which majorly remain spherical in shape. This is for a duration of 4 h

in a planetary ball mill performed at ILT Aachen.

The oxide particles upon bulk alloy re-melting during AM-process are expected to undergo

coarsening. Therefore, the size of oxides pre-AM i.e. before mixing (4 h milling) must ideally be

finer than that reported in the literature post-AM. The post-AM oxide sizes have been reported to

be in the range 25-60 nm [16,17] although for yttria. The powders used by these authors [16,17]

were commercial ferritic alloy (PM 2000 grade) with the oxides (0.5 wt.%) in the mechanically

alloyed form.

Here, the chosen yttria and lanthana powder, since the average crystallite size not exceed 30 nm is

certainly not unacceptable. By selecting fine oxide sizes seeking comparable mechanical

properties as that with yttria is intended; for instance high yield strength and the ultimate tensile

strength as that reported in Ref. [17]. In the next section, the laser processing parameters for LMD

4.3. Microstructural Characterization of Dense Samples

42

and the corresponding microstructural characterization are described for the dispersion with

lanthanum oxide.


4.3.1. LMD of Ferrite Powders mixed with Yttrium oxide

The parameters during the LMD process for a reference case, were maintained with a laser beam

diameter ‘Φ’ = 1.8 mm, and laser beam velocity ‘v’ = 600 mm/min. The powders chosen contained

0.5 wt.% of yttrium oxide (equivalent to 0.39 wt.% Yttrium) in ferrite.

Figure 4.4. EDS elemental mapping revealed for a reference case LMD trial with laser

parameters maintained at ‘Φ’= 1.8 mm, ‘v’ = 600mm/min. Note that the ferritic ODS steel

contained 0.5 wt.% of initially added yttria.

The EDS measurements by SEM for this trial is shown in Fig.4.4. It reveals a homogeneous

distribution of chromium and aluminum in the alloy matrix. Titanium segregates to certain

locations in the microstructure, with a size (diameter) of about 1 µm. Importantly, the measured

yttrium from EDS point scan (region encircled in white) is < 0.05 wt.%. This is significantly low

as compared to its initial amount of 0.39 wt.% Y, or 0.5 wt.% yttria.


43

To understand if the yttrium is present in the form of solid solution or if precipitated, a near-atomic

length scale characterization by APT was performed. The representative sampling was possible by

probing different locations in the as-LMD built sample.

Figure. 4.5 Schematic description of the sampling volume of an APT tip approximated to

be conical in shape. The red dots indicate the calculated mean number of yttria particles

per sampling volume of APT tip.

If yttrium is present in the form of precipitates its volumetric number density, or alternatively its

count for a given probe volume can be estimated. The calculation assumes their spatially

homogeneous distribution i.e. for their random distribution. Fig.4.5 reveals the probe volume

approximated to the shape of a cone with a diameter of 100 nm and a length of 300 nm,

corresponding to realistic sample dimensions. An average of 4 particles is calculated. It must be

noted the calculation assumes yttria size with a diameter equal to that estimated from XRD

measurements i.e. 15 nm.

Fig. 4.6 displays the mass spectrum plotted indicating the signal counts against the charge state

ratio (in Da). The elements constituting the matrix Fe, Cr, Al, Ti were identified in the mass


44

spectrum. More critically, the measurement via APT reveals absence of yttrium or its oxide for

different possible charge states.

Figure 4.6. Mass spectrum for a LMD produced alloy material with 0.5 wt.% of initially

added yttrium oxide particles.

The chemical composition of the bulk LMD as-produced sample was measured with ICP-OES

method. It is an absorption spectroscopy technique via wet chemical analysis. Fig. 4.7 reveals the

measured yttrium amount. It must be noted that the amount of yttrium in the sample is compared

against the initially added oxide amount corresponding to 0.39 wt.% (0.5 wt.% yttria). The

measurement reveals that yttria has undergone a significant loss. The final retained amount is about

a tenth of that initially added.


45

Figure 4.7. Bulk chemical analysis via ICP-OES measuring average Y (wt.%), pointing

out significant loss of yttria in the as-LMD material with reference to feedstock powder.

The amount retained in the as-LMD sample was about a tenth of that added.

In summary, characterization results at different length scales, by EDS, APT and bulk chemical

analysis collectively reinforces the inference that yttria has indeed undergone a significant loss

with reference to that initially added. This warrants reasons for the occurrence directing towards

possible ways to resolve this challenge. No prior research suggests the possible reason for noting

the loss of yttrium oxide particles during an additive manufacturing process.

4.3.2. LMD of Ferrite Powders mixed with Yttrium oxide: Yttria loss

challenge

A systematic study was conducted post each process step to note the yttrium content. Chemical

analysis results contribute to the understanding for identifying the cause for the yttria loss; this is

during the sequence of events within the process chain. It commences from the powder preparation


46

stage until the final LMD produced sample. In addition to compensate for the loss, another set of

experiments was conducted with an increased amount of initially added yttria, viz. 2 and 5 wt.%

yttria. Note that the chemical analysis measurements were on the samples with no specific

preparation for instance sample grinding.

Figure 4.8 Schematic representation of process steps from pre-milling of individual

powders until it enters the LMD melt pool. The amount of yttria retained at each process

step is represented in the bar graph below. This is for three levels of initially added yttria;

viz. 5, 2 and 0.5 wt.% Y2O3 denoted respectively in red, yellow, and blue (respectively 3.9,

1.56 and 0.39 wt.% Y).


47

Figure 4.9. Chemical analysis results represented in terms of measured yttrium content

retained in the melt pool. The study on the effect of following process parameters on

yttrium retained in the produced sample was conducted; (a) with laser beam diameter of

1.8 and 1.2 mm (laser scan speed of 600 mm/min); (b) with laser scan speed of 2400

mm/min (laser beam diameter 1.2 mm).

Each step of the process chain has been represented schematically in Fig 4.8. Also, in the same

figure the corresponding amount of yttria retained represented in the bar graph. The study was


48

performed for three different initial contents of yttria which have been represented by red, yellow

and blue color respectively for 5, 2 and 0.5 wt.% yttria (3.90, 0.78, 0.39 wt.% Y respectively).

In comparison to the reference trial, the intent of the next set of experiments lies in understanding

the effect of laser process parameters on yttria dispersion in the microstructure. The yttrium intake

in the as-LMD produced sample which increased to 0.2 wt.% Y with a fine laser beam diameter

of 1.2 mm (Fig.4.9 (a)). Note that the laser scan speed for the trial were same as the reference case,

600 mm/min.

Laser scan speed was independently altered to 2400 mm/min to obtain an enhanced cooling rate

compared to that in the reference case. The increase of speed was the maximum in the LMD

machine. This was to meet the objective of synthesizing ODS steel with a yttria content of 0.5

wt.% (0.39 wt.% Y). The yttria in the sample measured to be 0.44 wt.% (0.34 wt.% Y measured)

(Fig.4.9(b)).

To determine if the yttrium content was dispersed in a spatially homogeneous manner or if it was

segregated at a probe length scale (depth) of ca. 1 µm, EDS was performed. Fig 4.10 reveals the

details of the EDS elemental mapping. The results reveal the segregation of yttrium and aluminum

to form the upper most layer or along the ends of the melt pool. These are expected to be oxides

that form slag although contribute to the measured Y (wt.%) by chemical analysis.


49

Figure 4.10. Elemental mapping by EDS indicating yttrium and aluminum rich region in

the upper layer. The matrix does not reveal any presence of yttrium.

Alternative ways to achieve homogeneity of yttria can be during the AM process itself. If the yttria

can also be melted in addition to the ferrite alloy melt, then such a recipe could result in the

formation of a solution in liquid. Subsequent rapid solidification may result in its retention as a

solid solution. If this holds, then an annealing heat treatment if followed could lead to

homogeneous oxide precipitation.

The melting temperature of yttria is 2436°C [98]. Note that experimentally there exists a

measurement error during the measurement using pyrometers. This can be no better than ±1% [99]

which corresponds to ± 24°C. At such a high temperature, the iron based melt could have

dramatically high extent of superheat approaching 800°C. More importantly aluminum which is a

key alloying element essential for resisting spalling and oxidation [17], but vaporizes at 2470°C

[98]. Considering, the vaporization temperature of aluminum is also associated with a

measurement error of ± 24°C. It would not possible to control the melting of yttria experimentally

without the risk of aluminum vaporization. This is depicted using the Fig.4.11.


50

Figure 4.11 The transition temperatures for melting of yttria and vaporization of aluminum.

Second approach for tackling the challenge is during the pre-AM stage, the milling duration may

be increased. AM trials were performed with increased milling time of 10 h while keeping other

variables similar as before. Fig.4.12 (a) shows the 10 h milled powder serving as feedstock for

subsequent LMD. The initial yttria added to the ferrite powder for milling was 0.5 wt.% (0.39

wt.% Y). Although increasing milling for a high duration equivalent to that required for

mechanical alloying violates one of the key objectives of this work, LMD was performed with

powders but milled merely for 10 h. This was to serve as a benchmark case for comparison.

Fig.4.12 (b) shows the EDS elemental maps that indicate that majority of yttria is agglomerated

above the deposited material expected to be an oxide slag. Aluminum is also found to enrich the

respective region. This leads to the inference that milling duration range in the vicinity of up to 10

h is not a decisive parameter to promote spatial homogeneity of yttria by avoiding its

agglomeration or slagging.


51

Figure 4.12 (a) Feedstock powder after 10 h milling. (b) EDS elemental mappings of LMD

samples prepared the powders.

The trend of increased oxide retention and homogeneity could be expected to continue with high

solidification time. This may be strived by employing a process with noticeably smaller laser beam

diameter and faster scanning speeds. Selective Laser Melting (SLM) featuring these desired

processing parameters is discussed in section 4.3.3.


52

4.3.3. SLM of Ferrite powders mixed with Yttrium oxide

The section describes selective laser melting (SLM) with yttria as the second phase particles for

dispersion in ODS steel fabrication with feedstock powders milled for 4 h. A reference trial is

performed to note the spatial dispersion of yttrium oxide particles (diameter < 50 nm). The

parameters maintained were the following, a laser scan speed of 1200 mm/s and a beam diameter

of 90 µm (More details in Appendix 1). The two parameters favor high solidification rate (or high

cooling rate) in comparison to that by the LMD trials (0.6 mm and 40 mm/s (2400 mm/min)). Note

that the laser power was maintained at 160 W in the SLM trials.

The ‘Marangoni effect’ is likely to be more profound for fine melt pool dimensions as it is

associated with a high temperature gradient. The possibility of exploiting the ‘Marangoni effect‘

for particle dispersion in an alloy melt during SLM was reported previously in Ref. [16,17].

However, these reports mention the use of mechanically alloyed feedstock powders for SLM.

Figure 4.13 ICP-OES chemical analysis measuring Y (wt.%) in the ferritic steel matrix in

synthesized ODS material by SLM (scan speed of 600, 1200 mm/s).

A trend of increasing yttrium content is observed with laser scan speed, when increased from 600

to 1200 mm/s as shown in Fig.4.13. The measured yttium is however, is less than the nominal


53

amount of 0.39 wt.% Y (0.50 wt.% Y2O3). The yttrium intake by LMD reference case cannot be

compared against the SLM trial. This is because the sample preparation differs for the two

processes. In the latter, the as-produced samples are prepared by polishing which removes

prospective surface oxides.

The microstructure of the as-SLM produced material constituting the ferrite matrix containing

homogeneously dispersed particles. This is shown in the STEM bright field image in Fig.4.14. The

particles sizes in terms of diameter are comfortably less than 50 nm in diameter. The imaging

conditions also enabled contrast to distinguish dislocations in the matrix. The particle chemistry is

determined by APT which follows next.

Figure 4.14 STEM bright field image of particles dispersed in the ferritic steel matrix, with

a laser scan speed of 1200 mm/s.

Fig.4.15 (a) reveals the APT mass spectrum revealing the presence of Fe, Cr, Al, and Ti in the

alloy matrix. The probed volume contains yttrium (Y) enriched in the precipitate. To determine its

composition a proximity histogram was drawn with an iso-surface value of 2 at.% concentration

of Y. In the reconstructed volume titanium (Ti) and chromium (Cr), correlate with yttrium (Y) as

shown in subfigure (b). The oxygen content is expected to substantially be underestimated during

APT measurement [100] and not plotted in the proxigram.


54

Figure 4.15. (a) APT mass spectrum of the SLM-fabricated ODS steel. (b) Enrichment of

Y, Ti and Cr noted in the particles in the APT reconstructed sampling volume. (c) The

proximity histogram indicates the precipitate composition in (at.%).


55

In brief, the SLM trial reveals the possibility of obtaining dispersion of particles enriched in

yttrium, throughout the matrix of ferritic steel. STEM and APT confirm these findings, for the

SLM synthesized samples. The oxide content in the material of 0.16 wt.% Y or (0.2 wt.% Y2O3),

however remains to approach 0.39 wt.% Y (0.5 wt.% Y2O3). Prior research work on SLM of ODS

steels with yttria particles in Ref. [16,17] does not point out the possibility of yttria loss. The

authors [16,17] however, mention the use of mechanically alloyed ferrite with yttria. This

persuades a case necessitating the mechanically alloyed powders with the aim to approach the

nominal amount in the as-built sample microstructure specific to yttria oxide chemistry.

4.3.4. LMD of Ferrite Powders mixed with Lanthanum oxide

The LMD trials were performed with lanthana amount to be 0.5 wt.% (La of 0.44 wt.%) in the

feedstock powder. For dense ODS material fabrication by LMD a set of optimized processing

parameters were arrived at.

For a laser beam diameter, ‘Φ’ = 1.8 mm, a laser power ‘P’ = 600 W and a scan speed of ‘v’ = 600

mm/min was maintained during the trials. The track offset distance was controlled to 900 µm and

a layer height at 300 µm. To understand the lanthana retained in the as-LMD samples and its spatial

homogeneity, the microstructural characterization at different length scale down to near atomic

resolution was conducted via SEM-EDS, TEM and APT.

The spatial homogeneity of the dispersed particles are investigated firstly by SEM-EDS point

scans. At least 25 point scans were performed across the microstructure each with a sampling

volume (order of µm3 corresponding to the incident electron beam interaction volume). A


56

lanthanum amount reveals to be 0.35 ±0.03 wt.%. Note that the measured lanthanum amount if

considered to be in the form of La2O3 corresponds to a calculated 0.4 wt.%.

Figure 4.16. (a) Representative microstructure with measured matrix concentration for La

by SEM-EDS to be 0.35 wt.%. (b) Lanthanum rich oxide agglomerates present with a low

number density of 2.5 x 1015 m-2.

The EDS point scan measurements on the microstructure shown in Fig. 4.16 (a) indicate the

homogeneity of lanthanum concentration. However, agglomerated lanthanum rich particles

constitutes the microstructure (Fig.4.16 (b)), although with a meager number density of 5 x 102 m-


57

2. Despite the latter, it must be underscored that the spatial distribution of dispersed particles to be

homogeneous in the bulk alloy matrix. For a more detailed characterization by STEM and APT

are performed.

Figure 4.17. STEM image indicating presence of particles of size 100nm in the alloy

matrix. STEM-EDS reveals no presence of La in the particle or the matrix.

For imaging of the particles in the steel matrix, STEM characterization was performed as shown

in Fig. 4.17. Particle contrast observed, corresponds to a size (Feret diameter) of 110 nm with a

number density of 3 x 1015 m-3. However, the STEM-EDS measurement on the particle, does not

reveal enrichment of lanthanum. The possible explanation could be the limited sampling volume

for TEM analysis. A near atomic resolution by APT was also performed.

APT sampling volumes of 50 nm in diameter and 200 nm in length with needle shape were

investigated. This is repeated for 15 sampling volumes of similar volumes across the sample. The

mass spectrum shown in Fig. 4.18 was examined for presence of peaks of La+3, La+2 or La+ ions

4.4. Challenges and Comments

58

or LaO+3, LaO+2, LaO+, La2O+3, La2O

+2, La2O+1, LaO2

+3, LaO2+2, La2O2

+1. The characterization

reveals no presence of lanthanum or its oxide.

Figure 4.18. APT mass spectrum does not reveal the peaks that correspond to lanthanum

or its oxides for different charge states.


The chapter detailed on the microstructural characterization results of ODS steels produced by

LAM by milling for 4 h and obviating mechanical alloying. These were either with yttria or

lanthana as oxide particle chemistries. Yttria is found to undergo a significant loss in the bulk

LMD-produced alloy compared to that initially added. The SLM produced alloy although also

faces a significant fraction of particle loss, the remaining particles are found to be homogeneously

distributed in the bulk.

In the synthesized materials with lanthana by LMD, the particles are expected to be spatially

homogeneous for a length scale equivalent to that of the SEM electron beam interaction volume.

The homogeneity is independent of sparsely populated oxide agglomerates which are about a


59

micron in diameter and constitutes a minor fraction of lanthana in the fabricated material. A key

challenge lies in achieving a fine particle size (< 100 nm) for dispersion. It may be noted that the

particle chemistry of lanthana is more favorable for dispersion than yttria, the possible reason is

analyzed in the discussions section.


60

61

5. Cu-Cr-Nb Alloy Designed for Laser Metal Deposition

The chapter presents the microstructure and mechanical characterization of a designed Cu-3.4Cr-

0.6Nb (at.%) lean ternary alloy, hardened by LMD. First, we elucidate on the proposed alloying

regime. Second, the achievement of chromium nano-precipitates in addition to the known Laves

phase dispersed particles in the dense fabricated samples is revealed in the microstructural

characterization section (section 5.2). Subsequently, the hardness and nano-indentation

measurements attest high alloy hardening (> 130 Hv) and its spatial homogeneity is revealed.

5.1. Alloy Design

In Cu-Cr-Nb materials, the choice of Cr and Nb as the alloying elements in the Cu alloy base play

a vital role; these combine to harden the alloy via dispersed Cr2Nb Laves phase particles of sub-

micron size [25], while also permitting the matrix copper to remain nearly pure and conductive.

The latter has been achieved since Cr and Nb have poor solubility in Cu [25,83]. Note that the

5.1. Alloy Design

62

alloy design in previous functional alloys [79,83,101] have been bearing an alloying (in at.%) ratio

of Cr to Nb as 2:1. These alloys relied exclusively on Cr2Nb Laves phase among hardened alloys.

Here, compared to this ratio we propose choosing excess chromium. The intent lies in exploring

coherent nano-chromium precipitation during LMD apart from securing Laves phase for

hardening. Among lean alloys in the Cu-Cr-Nb system, this work presents a Cu-3.4Cr-0.6Nb

(at.%) designed specifically for laser metal deposition (LMD). Note that the alloying amount in

this alloy is 4 at.% (of Cr and Nb) which is lower than previous functional alloys requiring at least

(6 at.%).

Figure 5.1. A lean copper alloy compositional space in the ternary Cu-Cr-Nb system. The

present alloy base contains non-stoichiometric alloying amounts with reference to the

amounts in Cr2Nb phase. This is unlike previous hardened ternary alloys [79,83,101] which

strictly have obeyed this stoichiometry.


63

Fig.5.1 highlights the present alloy, in the Cu-Cr-Nb ternary compositional space in the lean copper

alloy base. The feedstock powder composition of Cu-3.4Cr-0.6Nb (at.%) was verified by

inductively coupled plasma optical emission spectrometry (ICP-OES), based chemical analysis.

Based on the amounts of Cr and Nb in current alloy, a 2.1 vol.% of Cr2Nb is calculated at room

temperature.

The chemical analysis also suggested the presence of 0.14 at.% Fe. Fe is expected to have entered

as an impurity in the alloy [27]. The alloy powders for the present study were produced in two

batches of 100 grams each by gas atomization process at IWT Bremen. For the melt pool formation

during LMD, the powders were directed to the laser focal distance by means of argon gas which

not only serves as the carrier gas but also resists oxidation.


5.2.1. Dense Sample Fabrication

For producing hardened alloys by additive manufacturing it is necessary to limit porosity i.e.

approach theoretical mass density (99.5% [7]). In AM, high density sample fabrication is sensitive

to feedstock powder characteristics particularly the size and the shape [2,102,103]. For this reason,

we analyze the feedstock powders for shape and size. This is followed by an analysis of as-LMD

produced samples for porosity and microstructure.

Fig.5.2 (a) shows the optical micrograph of the atomized powders which appear bright in

comparison with the mounting material, a conductive resin. The powders are quantified for the

shape and the size in terms of shape factor (4𝜋𝐴/𝑃2) and feret diameter (mean) respectively. The


64

histograms representing their distributions are displayed in Fig.5.2 (b). In the expression for the

calculation of the shape factor, 𝐴 and 𝑃 represent the area and the perimeter respectively.

Figure 5.2. (a) Optical micrograph of the as-atomized alloy (Cu-3.4Cr-0.6Nb-0.14Fe

(at.%)) powder for LMD processing. (b) Histogram representing the distribution of powder

shape and size, in terms of shape factor (4πA/P2) and feret diameter respectively. (c)

Backscattered electron (SEM) micrograph of atomized powder consisting of dispersed

particles in the copper matrix. (d) EDS elemental mapping of the corresponding region in

(c) indicating Cr and Nb enrichment in the dispersed particles, expected to be Laves phase.

In the shape factor distribution, the statistical mode lies in the range 0.8-0.9; a shape factor of 1

corresponds to a perfect spherical shape. This implies that the majority of the powders are nearly-

spherical. Similarly, in the distribution of the size, the statistical mode is 60-70 µm. For the

feedstock powder to be acceptable for LMD, its desired size range must be 40-90 µm [104,105]


65

along with near spherical shape [102,106]. The noted results indicate the suitability of the atomized

powders for subsequent LMD process.

The powder microstructure reveals the presence of second phase dispersed particles in the copper

matrix. The particles are enriched in Cr and Nb, as confirmed by SEM-EDS elemental mappings

of the corresponding microstructural region. These are shown in Fig.5.2 (c) and (d). The particles

are expected to be Laves phase intermetallic particles. The presence of Laves phase particles in

the ternary alloy powder was previously reported on the basis of SEM-EDS characterization,

although specifically in a Cu-4Cr-2Nb (at.%) [107]. In the present alloy, a minute amount of Fe

appears to have partitioned into the Laves phase. A detailed compositional characterization of the

particles by APT will be shown later.

Using the powders, sample fabrication by LMD with a low porosity level (mean) ≈ 0.12% was

possible. The porosity was measured from the micrographs of the as-built material by image

analysis. The measured porosity is lower than 0.5% (equivalent to 99.5% density) which is

considered to be a “fully dense” material in AM [7,108].

Size of pores as a metric is also sometimes used to qualify dense sample fabrication in AM. A pore

size of 300 µm [109] can be regarded as an upper size limit up to which sample fabrication is

considered dense. The pore sizes in the current sample is restricted to 30 µm. Each of the two

criteria, point out that the produced sample is acceptable in terms of density or low porosity.

It must be noted that in the microstructure, the adjacent melt pools overlap with their boundaries

outlined in red. This is shown for the uppermost deposited layer in Fig.5.3 (a). Here, a hatch

spacing distance of 800 µm which is lower than the melt pool width was maintained. Additionally,


66

the layer height increment (along build direction, ‘z’) was defined to be 200 µm. These LMD

parameters contribute to the observed low porosity.

Figure 5.3. (a) BSE image of as-LMD produced microstructure along a section orthogonal

to the laser scan direction. A low porosity level (mean) ≈0.12% is revealed. (b) Magnified

microstructure of the regions marked 1-6 in (a). The microstructure comprises large

columnar grains (~ 78 µm) grown along the build direction above the substrate. The

erstwhile melt pool boundary (marked in red) signifies the uppermost deposited layer. The

region consists of both equi-axed and columnar grains.

Fig.5.3 (a) and (b) reveal the as-LMD microstructure apart from providing the details on porosity.

The microstructure has been characterized along a section orthogonal to the laser scan direction

(xz plane). We examine the microstructure along the build direction starting from the bottom of

the build material. Subsequently, the region constituting the representative microstructure is

notified.


67

The bottom of the build material adjoining the substrate signifies the initial deposited layers

wherein large columnar grains are noticed (mean ferret diameter of 78 µm). These are nearly

parallel to the build direction, ‘z’, (subtend an angle of 81±5°with the ‘x’ direction). The

observation is typical of the LMD process [4,110,111], and occurs because of rapid heat

conduction away from the first laser scan track and into the substrate material that acts as heat

sink. Directional grain growth gets favored along counter heat flow direction, while the thermal

gradient develops.

In each of the subsequently deposited layers above (layer height ‘Δz’ = 200 µm) except for the

upper most layer (whose boundary is marked in red), the columnar growth is continued. The

columnar grain growth can be explained by epitaxial growth. This can occur by partial re-melting

of the layer deposited immediately underneath which serves as the nucleus for the directional grain

growth. Similar columnar grained structures in AM have been reported previously, for example in

nickel based superalloys by Gäumann et al. [5,112] and Dinda et al. [113].

The large columnar grains could have continued their growth in the subsequent layers, if more

than four layers (as in the present case) were deposited. Nevertheless, the large columnar grains

constitute the representative microstructure. This comprises the deposited material microstructure

except for that of the uppermost deposited layer along the build direction. The uppermost layer

reveals the erstwhile melt pool comprising both equi-axed and columnar grains (mean ferret

diameters of 17 µm and 40 µm respectively). The corresponding region is labelled as 1 and 4 in

Fig.5.3 (a) and (b). The melt pool grain morphology is analogous to that studied previously in laser

manufacturing, in Al and Ti based alloys [114,115].

The spatio-temporal solidification conditions evoking from the non-planar laser heat source, viz.

the temperature gradient ‘G’, and the solidification rate ‘R’[1,114], are expected to decide the


68

observed grain morphologies in the melt pool. Note that the equi-axed grains were most probably

also present after solidification of lower layers. But these regions near the top of the corresponding

melt pool are expected to have later re-melted to form the bottom part of the subsequent layer

above, growing as large columnar grains.

In brief, we observe low porosity in the fabricated sample which is favorable for producing

hardened alloys. This is confirmed by the microstructural examination, which also reveals the

representative microstructure.

5.2.2. Dispersed Laves Phase Particles

A contribution to the alloy hardening is expected from the Laves phase particles. The particles

constitute 2.2 ±0.1 % of the as-LMD produced microstructure, as determined by image analysis of

the SEM micrographs. Uniform particle distribution throughout the microstructure is suggested by

the SEM and EDS images in Fig.5.4 (a) and (b). At the length scale of 10 µm and beyond, particle

segregation for instance along the grain boundary is not observed. The particles are enriched in Cr,

Nb, and some Fe, as shown in the EDS elemental mapping (Fig.5.4(b)). To check if the

composition of the phase matches the Laves phase stoichiometry APT compositional

measurements were performed and will be discussed next.

A maximum particle size of ~1.23 µm is considered acceptable in particle dispersion strengthened

Cu-Cr-Nb alloys produced by conventional processing involving extrusion [83]. This would mean

that the sub-micron particles sizes post LMD are certainly small enough. Note also that the particle

sizes with reference to those in the powder have not coarsened beyond the desired size. The high

in-process cooling rates during LMD [116] is likely to have restricted such coarsening.


69

Figure 5.4. (a) SEM image acquired in secondary electron mode reveals homogeneous

distribution of particles, expected to be Laves phase. (b) The EDS elemental mapping

suggests that the particles are enriched in Cr, Nb and Fe.

We determine the particle composition by performing APT which is shown in Fig.5.5. The results

show that the particles are enriched in Cr, Fe, and Nb. This complies with the SEM-EDS result

suggesting Fe enrichment in the particles. Fe is soluble to a greater extent in both Cr [117] and Nb

[118], than in Cu [119], and explains the observed result.

The measured particle composition by APT corresponds to (Cr,Fe)2Nb and matches the A2B

stoichiometry. Hence, we conclude that the particles are indeed the Laves phase. A

crystallographic pole was noted in the detector event histogram during the APT experiment. The

(002) pole of the face centered cubic (FCC) due to the symmetry observed on the detector was

identified for the copper matrix [92]. During the reconstruction, the parameters were chosen such

that the reconstructed inter-planar distance was calibrated with the known (002) inter-planar

distance of pure copper. In the copper matrix, up to a fine length scale of about 20 nm no other

particles or precipitates were observed.


70

Figure 5.5. Compositional characterization of the particles in the copper matrix by APT.

The results suggest the enrichment of Cr, Fe and Nb in the particles. The overall

composition corresponds to A2B stoichiometry, indicative of the Laves phase. The

compositional profile has been calibrated with the known (002) inter-planar distance of

pure copper.

The previously studied high resolution compositional measurement of Laves phase particles in Cu-

Cr-Nb ternary alloys was carried out by TEM-EDS spectra by Anderson et al. [83]. They inferred

a Cr to Nb atomic ratio of 2:1 which corresponds to Cr2Nb Laves phase composition in a

conventionally produced Cu-8Cr-4Nb (at.%) alloy. To the author’s knowledge there exists no

other compositional information with high spatial resolution for example by atom probe

tomography (APT) [120,121] in Cu-Cr-Nb based alloys hitherto. APT characterizations of other

Laves phases however, have been reported in the literature. For example (Fe,Cr)2Mo [122],

(Fe,Cr,Si)2Mo [123], (Fe,Cr)2Zr [124], (Fe,Cr)2W [125], and Fe2(Mo,Ti) [126].


71

5.2.3. Chromium Nano-precipitates

In addition to the Laves phase dispersed particles another contribution to the alloy hardening is

expected from the chromium nano-precipitates formed in-situ during LMD in the produced alloy.

In the following, we present the characterization results of the nano-precipitates.

Figure 5.6. The TEM dark field image taken along the [111] zone axis of the copper matrix

in the as-LMD produced alloy. The fine bright spots suggests homogeneous distribution of

nano-precipitates. (b) APT characterization reveals the nano-precipitates to be ~ 4 nm in

size, while affirming their uniform distribution. (c) Proximity histogram based on a 10 at.%

Cr concentration value for the iso-concentration surfaces. It shows the elemental

concentration as a function of distance normal to the precipitate-matrix interface.

5.4 Hardening Assessment

72

Fig.5.6 (a) shows the dark field TEM image acquired along the [111] zone axis of copper matrix

for the (220) diffraction spot. It reveals a homogeneous distribution of fine bright spots. This is

indicative of the coherent precipitates of chromium [27,127]. APT was performed for a detailed

characterization of these coherent precipitates.

The APT characterization affirms the presence of fine nano-precipitates (number density 8x1023

m-3; 4 nm mean diameter) distributed homogeneously in the copper matrix as shown in Fig.5.6 (b).

These nano-precipitates contain chromium which is present nearly in an equiatomic amount as the

matrix element, copper. This chemical compositional information is according to the proximity

histogram [128] shown in Fig.5.6 (c). It is obtained by plotting the elemental concentration as a

function of distance normal to the precipitate-matrix interface based on a 10 at.% Cr iso-

concentration surfaces.

Note that the nomenclature followed here, distinguishes the particles from the precipitates. The

former does not dissolve in the matrix unlike the latter, at high homologous temperatures. The

individual hardening contributions arising from the Laves phase dispersed particles as well as the

nano-chromium coherent precipitates will be discussed in chapter 6.


5.4.1. Nano-indentation Measurements

This section assesses the validity of the spatial homogeneity of alloy hardening in the

microstructure arising from the coherent chromium nano-precipitates and the dispersed Laves

phase particles. The assessment is made, on the basis of nano-hardness values by performing arrays

of nano-indentations.


73

Figure 5.7. (a) Spatial sites for performing arrays of nano-indentations (box marked in

green), in the representative microstructure. (b) Nano-indentation load-displacement curve

of the as-produced material from which nano-hardness is determined using the procedure

by Oliver and Pharr [129]; scanning probe microscopy image of the nano-indent in the

inset. (c) Bar graph representation of the nano-hardness values compared with that of pure

copper from Ref.[130]. (d) Visualization of the nano-hardness data as a 2D spatial contour

plot; each indentation position is represented by a black colored dot. (e) An IPF map of the

region corresponding to that in (d).

Fig.5.7(a) shows the sites for performing nano-indentations, along the representative

microstructural regions consisting of columnar grains along the build direction, grown above the

substrate. The probe length scale for each indent is ~ 5 µm which is the size of nano-indent and


74

the surrounding strain field (cf. inset in Fig.5.7 (b)). It is sufficiently lower than the center to center

distance maintained between the closest indents of 40 µm. The prevention of hardness

overestimation due to spatial overlapping of indentation strain fields is thus ensured.

The nano-indentation load-displacement curve is displayed in Fig.5.7 (b) from which the nano-

hardness values were evaluated following the procedure established by Oliver and Pharr [129].

The evaluated nano-hardness (H) of 2.12±0.2 GPa, is up to 2.5 times that of values reported

previously for annealed and work-hardened conditions in pure copper (grade: oxygen free copper

(OFC)) [130]. The respective nano-hardness values of 0.85 and 1.7 GPa [130] are compared

against the present measurements in a bar graph representation in Fig.5.7(c). This leads to the

inference that hardening in the present material is significant.

Note that the degree of scatter of the nano-hardness in terms of one standard deviation is bound

within less than ±10% of the mean value. Fig.5.7(d) visualizes the nano-hardness data in a spatial

manner, as a 2D contour plot. At the sites of nano-indents, in order to co-relate with the grain

orientation of the matrix copper EBSD analysis was performed. This is revealed in the inverse pole

figure (IPF) map in Fig.5.7(e).

The 2D spatial contours of the nano-hardness data and the degree of scatter in the bar graph,

highlight the inference of spatial homogeneity of hardening. In fact, the grain orientation

anisotropy of copper leads solely to a scatter of 5% of the mean nano-hardness, remarked

previously in Ref. [130,131]. This implies that the scatter due to the coherent nano-precipitates

and the Laves phase particles must certainly be less than 10%.

The elastic modulus is also deduced from the measured load-displacement curves to see if it is

comparable with that of the other Cu-Cr-Nb alloys. The elastic modulus, E, is calculated from


75

equation (1) [129]. The determined reduced modulus, Er, is 126±8 GPa; Poisson’s ratio ‘υ’ is 0.35

[132]. The subscript, ‘i’ refers to the indenter properties; υi and Ei are 0.07 and 1141 GPa

respectively.

1

𝐸𝑟=

1−𝜐2

𝐸+

1−𝜐𝑖2

𝐸𝑖 (1)

The elastic modulus of the alloy is calculated to be 124±9 GPa. The value tallies closely with that

for Cu-8Cr-4Nb and Cu-4Cr-2Nb (at.%) ternary alloys each of which is about 120 GPa [79].

Analogous to the nano-hardness values, no microstructural dependence of elastic modulus values

is observed. For hardening assessment at a large length scale Vickers indentation hardness has

been performed and will be discussed next.

5.4.2. Hardness Measurements

Analogous to the nano-indentation results, a markedly high hardness of 146 ±13 Hv was measured.

This is about thrice the hardness of pure copper of 50 Hv [79]. Fig.5.8 compares the hardness of

the present alloy against those of Cu-Cr-Nb alloys studied previously.

For a Cu-4Cr-2Nb (at.%) alloy in the as-extruded condition a hardness of 117 Hv was reported

[79]. Under a similar condition for a more concentrated alloy, Cu-8Cr-4Nb (at.%), the hardness

was shown to be 128 Hv [83]. Anderson et al. [79], revealed that by refining the grains in this alloy

as compared to those in Ref. [83], the hardness increases further although marginal measuring 132

Hv. It may be noted that alloying amount has been a key factor that decides the Cu-Cr-Nb alloy

hardening.


76

Figure 5.8. Vickers hardness of the alloy in the present study, in comparison with Cu-Cr-

Nb alloys studied previously [79,83]. Additionally, the aggregate alloying amount is

plotted in the same graph.

The present finding however does not obey the general trend of hardness increase with total

alloying solute amount (of Cr and Nb). While an alloying amount of 12 at.% was required to harden

the Cu-8Cr-4Nb (at.%) alloys a mere 4 at.% suffices to comparably harden the present lean alloy.

The measured Vickers hardness compares well with other lean Cu-Cr based alloys such as Cu-Cr-

Ag [133] and Cu-Cr-Zr [27,134]. For a Cu-0.3Cr-0.1Ag (wt.%) alloy, a hardness of 144 Hv in the

peak aged condition was shown by previous authors in Ref. [133]. In a study performed on a Cu-

1Cr-0.1Zr (wt.%) alloy by Chbihi et al.[27], a hardness value of 155 Hv in the peak aged condition

was reported.

The quoted hardness values of the aforementioned lean ternary alloys [27,133,135] were

determined in samples which had been subjected to plastic straining and/or an ageing heat


77

treatment. The comparable hardness of the present alloy however, is noted in the as-LMD

produced state itself even without imposing any additional ageing heat treatment(s).

In summary, Vicker’s hardness and the nano-hardness values clarify high alloy hardening and its

spatial homogeneity. This is for a probe length scale of 5 µm and greater, at which the

microstructural characterizations show excellent homogeneity of Laves phase particles and

coherent nano-precipitates containing chromium.

Achieving such a microstructure in the Cu-Cr-Nb system which promises high hardening, is

sensitive to alloy design in terms of chromium content as well as in-process cooling rate.

Consequently, deviation of these two factors from that identified presently is expected to result in

drastically different microstructures, more importantly compromising on alloy hardening.

Therefore, the appropriate combination of these two factors is discussed in chapter 6.


78

79

6. Discussions

The chapter discusses the microstructural findings and relates to the processing conditions during

LMD and SLM, for ODS steel fabrication. In the Cu-Cr-Nb alloy, the microstructural

characterization clarified the design of a lean alloy hardened in a novel manner via LMD. This is

via nano-chromium coherent precipitates and Cr2Nb Laves phase particles. The alloy hardening is

comparable to those in this ternary system which previously have required greater alloying

amount. The present microstructure is achieved by identifying a desired combination of the LMD

processing parameters and the alloy compositional regime, elucidated in this chapter.

6.1. ODS Steels

6.1.1 ODS Steels containing Yttria particles

6.1. ODS Steels

80

The ytrria particles retained in the ferritic steel matrix of the as-SLM built sample are distributed

homogeneously as shown previously in STEM micrographs (Fig.4.14). However, this is about 0.2

wt.% of yttria as the remaning particles constituting a major fraction, agglomerates along the

uppermost built layer. The occurance is because the event of oxide agglomeration preceeds the

possibility of such particles being arrested in the as-solidified material. In other words, the time-

scale for agglomeration is faster than the time-scale for the material solidification. The explanation

is consistent with the microstructural differences noted between those obtained by LMD in

comparison to those by SLM. The mechanism underlying the loss of yttria could be due to the

following two reasons. These are oxide evaporation, or the oxide-alloy interface weakening during

laser interaction. The latter is considered plausible and explained next.

Vaporization Possibility

The possibility of evaporation is investigated by comparing the laser input energy density with the

latent heat needed to vaporize yttria. It must be noted that the yttria does not accept the entire

energy from the laser source, since its absorptivity is not 100%. The absorptivity value can be

approximated to a maximum value of 100 ppm ~ 0.01% [87,88] which is considered in the present

calculation. In this calculation, the sensible heat is considered to be twice the amount of the latent

heat [138].

The input laser energy per unit volume (J.mm-3) for a laser power of 1000 W, a beam diameter ‘Φ’

= 1.8 mm, a scan speed of ‘v’ = 600 mm/min, and a layer height ‘Δz’ = 500 µm. The laser energy

per volume evaluates be 110 J.mm-3.

For the vaporization event to occur, a latent heat is calculated for a powder material with 0.5 wt.%

of initially added yttria to ferrite. From Fig.4.8, the yttria loss is 0.2 wt.% (0.16 wt.% Y).

6.1. ODS Steels

81

Latent heat to vaporize 0.2 wt.% initial yttria: 735 J.mm-3

Maximum Laser input energy density: 110 J.mm-3

The calculations suggest that even the maximum laser energy is insufficient to meet the energy

requirements for promoting yttria vaporization. However, it could play a partial role for the

observed yttria loss.

Weakening of oxide-alloy powder interface during laser interaction:

The yttria loss from the bulk material during the laser interaction due to weakening of the interface

between the oxide and the alloy is examined. After laser interaction the bulk ferrite alloy powders

melt at a temperature lower than that needed for melting yttria. In the sequence of events, the onset

of solidification of the bulk alloy is likely to be not fast enough to precede interface weakening

between yttria-bulk alloy. This is represented in the equation 6.1. At this stage the effect of

interface weakening may also be aided by the shield gas flow.

solidification time > time for Y2O3–ferrite alloy powder interface weakening (6.1)

We test the validity of the plausibility by decreasing the solidification time. Although solidification

time has not been independently controlled or measured, it has been qualitatively decreased by

choosing the processing parameters to increase the cooling rate. This is possible via a fine laser

beam diameter and/or a high scan velocity in LMD.

The hypothesis complies with the measured yttira by LMD, revealing the increase in yttria by

decreasing the laser beam diameter ‘Φ’ from 1.8 mm to 1.2 mm. The discussion points to interface

weaking as the likely mechanism for the noted yttria loss during LMD. It also explains the

6.1. ODS Steels

82

observed microstructure processed by SLM with greater oxide particle retension, and the particle

spatial homogeniety.

6.1.2 ODS Steels containing Lanthana particles

The present LMD based rapid solidification processing route was aimed to achieve a spatial

homogeneity of lanthana (< 100 nm in diameter; 0.5 wt.%). One possible approach is if lanthana

is melted in addition to the ferrite alloy, then such a recipe can result in a liquid solution during

Marangoni convention [16,17] driven by temperature dependent surface tension. Rapid

solidification during LMD could result in its retention as a solid solution. A subsequent annealing

heat treatment, can lead to a homogeneous oxide precipitation. In the following calculation, the

energy for melting lanthana is compared against the laser energy density. While it assesses such a

plausibility it explains the observed microstructure.

In the calculation for energy density for melting of lanthana, both sensible heat and latent heat are

considered for lanthana. The former is considered to be twice the latent heat [138]. The calculated

value is compared with the input laser energy density stated in equation 2.1.

Maximum Laser input energy density: 110 J.mm-3

Minimum energy density for fusion of lanthana: 300 J.mm-3

The calculation reveals that the energy requirement is partially satisfied. More specifically, at most

a third of lanthana undergoes melting. The melting if initiates from the particle surface is likely to

melt the finer particles, belonging to the distribution of particles sizes. The remainder lanthana is

less probable to melt and are favored to undergo coarsening and/or agglomeration due to the time-

temperature combination experienced in the LMD melt pool. The argument explains the

6.1. ODS Steels

83

agglomerated lanthana (low number density of 5 x 102 m-2) in the microstructure shown in Fig.

4.16(b).

In the bulk alloy, the EDS point scans indicate the possibility of lanthanum enrichment in the solid

solution or its homogeneous dispersion in the probe volume (µm3). However, the STEM and APT

analysis rules out the possibility lanthanum oxide in the solid solution. This leads to the inference

that the time scale for solidification during LMD is not short enough to restrict the melted lanthana,

to remain in the liquid solution. Previous work during rapid solidification in an immiscible system,

although in Cu-Fe based systems [37,139] have reported similar observations. The findings were

explained Marangoni motion based on temperature dependent interfacial energy between the two

separating phases in liquid.

An approach to address this challenge requires lower solidification time. SLM could promote such

a possibility which features high cooling rate (> 104 K/s) than that by LMD (103-104 K/s) [2,7].

The approach could further be favored with feedstock powders which are mechanically alloyed as

the oxides would be present in the solid solution. This is expected as the particle coarsening from

solid solution is expected to less likely than that from the particles which exist in the melt

(nucleated or grown). Note that previous research reporting successful dispersion of oxide particles

during LAM processing, specifically by SLM have required mechanically alloyed powders

containing oxides of yttrium.

Interestingly, the present mixed powders (short milling time of 4 h), is able to reveal homogeneity

for a probe volume of µm3, although not at a finer characterization volumes. The effect of

dispersion due to the Marangoni effect in conjunction with rapid solidification process during

LMD explains how the present microstructure differs from that reported by liquid processing

routes like casting. A cooling rate of 102 K/s or lower in casting is said to result in extensive

6.2. Cu-Cr-Nb Alloy

84

agglomeration and slag formation [16,17]. More critically, the energy calculation also clarifies as

to why lanthana gets favored than yttria for dispersion by LMD process.

6.2. Cu-Cr-Nb Alloy

In the presented results in chapter 5, it is clear that the novel way of hardening a lean alloy via

nano-chromium coherent precipitates and Laves phase particles, is indeed substantial. The

hardening contribution from the achieved in-situ coherent chromium nano-precipitates is

calculated and noted to be significant. The in-situ precipitation is attributed to the synergy between

the alloy design and the LMD processing elaborated in detail here. Note that the attained hardening

is comparable with previous alloys containing 12 at.% alloying (Cr and Nb). For example the Cu-

8Cr-4Nb (at.%) which relied exclusively on the Cr2Nb Laves phase particles for hardening

[79,83,101].

In the known Cu-8Cr-4Nb (at.%) and Cu-4Cr-2Nb (at.%) alloys, the hardening contributions have

been due to Hall-Petch grain refinement (fine grains ~ 2.7 µm) [79,83,101] as well as incoherent

particle dispersion of Laves phase, present in amounts 7 - 14 vol.% [79,83,101]. When compared

with the hardening contributions in these alloys, the respective contributions are expected to be

less dominant in the present alloy. This is because, the present microstructure constitutes large

grains of size ~ 78 µm and a mere 2.2 vol.% of comparable sized Laves phase particles.

However, the hardening contribution from the nano-chromium coherent precipitates in the

microstructure must be considered. On the basis of the measured precipitate chemical composition

and its size from APT, we evaluate the coherency hardening increment [140] for the nano-

precipitates in the alloy matrix. This evaluates to 234 MPa which is equivalent to 78 Hv (converted

6.2. Cu-Cr-Nb Alloy

85

by Tabor's approximation [141]). It is a substantial increment with reference to that of the copper

base, of 50 Hv [79]. The calculation details are presented in Appendix 2.

Similarly, the hardening increment for the sub-micron sized Laves phase particles is calculated

using the Orowan-Ashby relation [79,142]. These Laves phase particle hardening increments

evaluates to 22 Hv (details in Appendix 2). This is consistent with that evaluated for nano-

precipitates of chromium, as they collectively add to 150 Hv and matches the measured hardness.

Seeking coherent nano-precipitates containing chromium in the microstructure is sensitive to the

alloy composition and the process.

The LMD processing accompanies a cooling rate which lies typically in the range, 103-104 K/s [2].

Slower solidification process (cooling rate < 102 K/s) like casting is often regarded as detrimental

to alloy hardening by chromium precipitates although in binary Cu-Cr systems. This is so, because

of the chromium precipitate coarsening (> 1 µm sized) as well as due to the undesired precipitate

segregation causing their spatial inhomogeneity [143,144]. A high cooling rate (~106 K/s) on the

other hand by rapid solidification processing, was said to result in a chromium supersaturation into

the copper matrix resulting in a supersaturated solid solution. This was in the solidified

microstructure in a binary alloy Cu-2Cr (wt.%) [26]. For a Cu-5Cr (wt.%) alloy however, also in

the same study, 50 nm sized incoherent chromium particles were shown, other than the chromium

leftover in the supersaturated solid solution [26]. This suggests that given the high cooling rate,

the alloy designed in the noted concentration regime of chromium, plays a key role in controlling

the microstructure in the Cu-Cr system.

6.2. Cu-Cr-Nb Alloy

86

Figure 6.1. (a) Schematic comparison of the solidification microstructures in the Cu-Cr

alloy system taken from Ref. [26,143,144] and the present quasi binary Cu-Cr alloy. In the

representation, chromium content available for precipitation is plotted along the abscissa

and the in-process cooling rate along the ordinate. (b) The Cu-Cr binary phase diagram in

the vicinity of the eutectic composition as taken from Ref. [145]

6.2. Cu-Cr-Nb Alloy

87

Presently, from an alloy design standpoint, we have conceived a Cu-Cr-Nb based system which

during solidification behaves like a quasi-binary [146] Cu-Cr system. This is because of the

following reasons. First, Nb is contained entirely in the Laves phase particles which are not soluble

in copper, even in its liquid form up to nearly 1600°C [147]. The insoluble Laves phase particles

contain Cr and Nb, which can reduce the alloy base to that of a binary Cu-Cr. Second, to ensure a

quasi-binary Cu-Cr base, we conceive a chromium amount which approaches ~ 1.6 at.% with

reference to the chromium contained in the Laves phase. The Cr amount corresponds to the eutectic

composition [145] in the quasi-binary Cu-Cr system (Cu-Cr phase diagram shown in Fig.6.1 (b)).

If the chromium in the liquid alloy is substantially increased with reference to the eutectic

composition i.e. in the hyper-eutectic compositional regime, then pro-eutectic chromium phase

formation is predicted in the binary Cu-Cr phase diagram [145]. The predicted phase complies

with the microstructural observations for a Cu-5Cr (wt.%) alloy in Ref. [26], reporting the primary

(pro-eutectic) chromium (> 50 nm sized). The pro-eutectic chromium can lead to a marginal

hardening increment of 12 MPa estimated according to the Orowan-Ashby formulation for

incoherent precipitates. This is equivalent to a mere 4 Hv increment in hardening (corresponding

to every 1 at.% Cr added beyond the eutectic amount, up to 5 at.% of Cr addition). Therefore its

presence in the microstructure is not desired, as the marginal hardening increment is at the cost of

relaxation of the constraint imposed by the alloying amount in dilute alloys or lean alloys.

Hardening contribution from pro-eutectic chromium can be substantial only if it firstly qualifies

for coherency, in terms of its size, which must be less than 10 nm [84,127]. This is beyond the

plausibility of the LMD process considering its cooling rate and also for other rapid solidification

processing routes. Clarification for the latter follows the findings in the work by Morris et al. [26]

revealing that even a high cooling rate of 106 K/s co-relates to a precipitate size which at best can

6.2. Cu-Cr-Nb Alloy

88

be restricted to 50 nm. Moreover, such a high cooling rate is coupled with an undesired effect

which lies in not aiding the chromium precipitation from the supersaturated solid solution (SSS),

correlating to the hypo-eutectic chromium amount in the alloy.

It follows then that for exploiting coherency hardening from hypo-eutectic chromium is

appropriate, but mandates an ageing heat treatment as also mentioned in Ref. [26]; the solute

belonging to the hyper-eutectic regime which corresponds to hardening via pro-eutectic chromium,

would remain nearly as redundant for hardening as without ageing.

These considerations were conceived for opting an alloy with a chromium amount which does not

substantially exceed the eutectic amount, whereas limited amount (< 0.2 at.%) prevents from

exploiting substantial coherency hardening. Note that for the conceived alloy, an in-process

cooling rate of < 102 K/s is undesired because of the formation of coarse chromium precipitates

[143,144]. The present findings concerning the in-situ formed coherent nano-precipitates and the

alloy hardening, illustrate that the LMD cooling rate is indeed well suited for the developed alloy

despite being sensitive to the chromium content in the alloy and the cooling rate.

Fig.6.1 (a) presents the summary of the solidification microstructures of chromium precipitates in

copper alloy. This is shown as a function of chromium content in the alloy available for

precipitation and in-process cooling rate.

In summary, the present work brings forth a new alloy to the class of dilute copper alloys and the

Cu-Cr-Nb alloys. The in-situ coherent precipitation during LMD enables to access the desired

microstructure for high alloy hardening. This is the key novelty of the current processing routine.

Even though the presence of coherent chromium precipitates in copper alloys is well known

[27,127], the current work is the first attempt at exploiting them as an additional hardening source

6.2. Cu-Cr-Nb Alloy

89

in the ternary Cu-Cr-Nb system. Previous works relied on strengthening by Cr2Nb Laves phase

only. The chromium rich coherent precipitates enables reaching higher aggregate hardening than

previous Cu-Cr-Nb alloys, despite bearing a meager alloying amount of 4 at.%.

6.2. Cu-Cr-Nb Alloy

90

91

7. Summary and Concluding Remarks

Oxide dispersion strengthened ferritic steels:

An alternative approach towards ODS steel fabrication via laser additive manufacturing, but by

obviating mechanical alloying process step has been the focused of the present work. The intent

underlying this approach was to exploit Marangoni convection in the alloy melt pool, for aiding

particle dispersion during LAM. The materials produced were either with yttria or lanthana as the

oxide particle chemistry for dispersion.

1. In the as-produced steel, yttria particles suffer a major loss compared to that initially added.

This is expected during the laser interaction, after which the bulk ferrite alloy powders melt

at a temperature lower than that needed for melting yttria. In the sequence of events, the

onset of solidification of the bulk alloy is likely to be not fast enough to precede interface

weakening between yttria-bulk alloy.

92

2. Among as-produced materials by SLM and LMD, the former showed a greater extent of

particle dispersion and retention (< 100 nm in diameter; 0.2 wt.% Y2O3) than those by

latter.

3. The LMD synthesized material with lanthana showed uniformity of measured lanthanum

concentration for a probe volume, equivalent to the electron beam interaction volume of

about 1 µm3. However, the spatial homogeneity at any finer length scale is expected to

require a higher solidification rate, for restricting the oxide particles from coarsening

beyond a size of 100 nm.

4. It could be inferred that for oxide particle dispersion in ferritic steels mechanical alloying

of feedstock powders plays a decisive role. Avoiding mechanical alloying is considered

challenging.

Cu-3.4Cr-0.6Nb (at.%) alloy:

The work reveals a lean alloy in the Cu-Cr-Nb based system by LMD hardened in a novel manner

via nano-chromium precipitates and dispersed Laves phase particles. The alloy hardness of 146

Hv is 11% higher than the strongest known alloy in this system, Cu-8Cr-4Nb (at.%). The spatial

homogeneity of alloy hardening is verified from the nano-hardness values, visualized using the 2D

nano-hardness spatial contour maps generated by performing arrays of nano-indentations.

The key inferences are as follows:

1. The current alloy is hardened by introducing fine (number density 8x1023 m-3; 4 nm in

diameter) nano-chromium coherent precipitates. This is in conjunction with the known

possibility of hardening via the Laves phase dispersed particles.

93

2. For favoring nano-precipitation, an alloy which can be considered as a quasi-binary Cu-Cr

system was conceived. This is because Nb is entirely contained in the Laves phase particles

which are insoluble in copper matrix apart from some Cr. Compared to this amount, we

choose an excess Cr amount ~1.6 at.% approaching that of the binary Cu-Cr eutectic.

Further Cr addition in the alloy is expected to result in pro-eutectic chromium particles.

These are less potent towards alloy hardening because of their coarse size (> 50 nm) which

cannot be refined further by imposition of a high cooling rate during solidification of 106

K/s [26].

3. The cooling rate during LMD suites the precipitate sizes to grow not beyond the coherent

size regime. With reference to typical cooling rates of 103-104 K/s in LMD [2], its

deviations are expected to result in drastically different microstructures which depreciates

hardening. High cooling rate of 106 K/s leads to chromium supersaturation in the solid

solution [26]; on the contrary low cooling rate of < 102 K/s results in large chromium

particles (> 1 µm) [143,144].

To harden the alloy by in-situ coherent nano-precipitates containing chromium, a delicate but a

desired combination of a cooling rate and the chromium content in the designed alloy has been

obtained. The presented recipe can have a significant implication on designing of microstructures

with nano-chromium precipitates in other Cu based systems on one hand. On the other, the alloy

hardening in this ternary system could be tuned by independently altering the Laves phase fraction

in the microstructure while retaining the quasi-binary Cu-Cr system.

95

Appendix 1

The section presents details on the ODS ferritic steel samples fabricated by selective laser melting

(SLM) with a sample dimensions (in mm) of 5x5x10. The key objective was to lower solidification

time; this was attempted with a high laser scan speed (1200 mm/s) coupled with a fine laser beam

diameter (90 µm). The parameters are more favorable than that by LMD (section 4.3) with a laser

beam diameter of 0.6 mm and a high laser scan speed of 2000 mm/min. A picture of the SLM

produced sample is shown in the following.

Fig. A1. ODS ferritic steels with yttria (0.5 wt.%) fabricated by selective laser melting (SLM).

97

Appendix 2

The hardening contribution arising from the coherent nano-chromium precipitates is examined.

This is followed by the contribution calculated for the Laves phase particles.

Firstly, we calculate the coherency hardening increment from the chromium nano-precipitates,

based on the measured APT data. This is using the coherency hardening formulation following the

equation (A1) [140].

𝛥𝜎𝐶𝑆 = 𝑀𝜒𝐺(휀)1.5 (𝑟𝑓/𝛼𝑏)0.5 (A1)

Here, M stands for the Taylor's factor, which equates to 3 for a collection of grains with no

preferred orientation in the microstructure [140], considered to be the case presently. χ is a

theoretical value which varies for each of the different theories that explain precipitate coherency

strengthening; it falls in the range from 2 to 3 [140]. With reference to Ardell [140], the value is

taken as 2.6. The shear modulus of matrix copper, G, is 42.1 GPa [140,148].

r and f represent the precipitate radius and mean volume fraction, taken from the current APT

measurements to be 2±0.3 nm and 1.1% respectively. The Burgers vector, denoted by b is equal

to 0.255 nm for the FCC copper matrix. The parameter ‘α‘ can vary in the range spanning across

0.089 and 0.5 [140]. The two values correspond to maximum theoretical strengthening and a lower

bound for strengthening, respectively. The latter holds true under the approximation of dislocation

line tension, which is assumed [149] in the present calculation.

휀 = 𝛿/[1 + 2𝐺(1 − 2𝜈𝑝)/𝐺𝑝(1 + 𝜈𝑝)] (A2)

ε in the equation, stands for the misfit strain parameter, which is sensitive to elastic constants and

the lattice parameters of the precipitate and the matrix, given by a relation in equation (A2) [140].

98

Here, the subscript ‘p’ specifically refers to the precipitate. The misfit, δ, is calculated from the

lattice parameter of the matrix copper, aCu = 3.61 Å [150] and that of the Cr precipitate, aCr = 3.68

Å [150]. The aCr value is approximated by assuming chromium precipitates to be formed in nickel

matrix, since the value for that in copper matrix has not been previously reported in literature. The

calculated misfit, δ, turns out to be 1.91%. Hence, the value of ε evaluates to 0.0145, for a poisson's

ratio, 𝜈p = 0.3 [149] and for a precipitate shear modulus Gp = 78 GPa. The latter value is predicted

by Vegard's rule [151] on the basis of measured precipitate composition by APT, which is nearly

equiatomic in Cu and Cr.

The calculated coherency strengthening increment amounts to 234 MPa, which is equivalent to 78

Hv (converted by Tabor's approximation [141]). The increment with reference to the hardness

value for pure copper of 50 Hv [152], is considerable indeed. Next, the hardening increment

calculation by Laves phase particles is shown.

The hardening increment evoking from the dispersed Laves phase particles is evaluated using the

Orowan-Ashby relation, shown in equation (A3) [79]. The formulation is considered applicable

for the incoherent precipitates. Here, 𝜆 which is the mean particle spacing given by 𝑟√2𝜋/3𝑓. The

other symbols in equation (A3) represent the same quantities as that in (A1).

𝛥𝜎𝑂𝐴 =0.84 𝑀𝐺𝑏

2𝜋(1−𝜈)0.5(𝜆−2𝑟)ln (

𝑟

𝑏) (A3)

Particles of up to 70 nm in diameter correspond to a maximum hardening increment. This equates

to a 66 MPa or 22 Hv. The individual hardening increments each of which when added to that of

reference copper, sum up to 150 Hv.

99

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Curriculum Vitae

PERSONAL DATA

Name: Anoop Raghunath Kini

Date of Birth: 25 July 1986

Place of Birth: Manipal, India

Nationality: Indian

EDUCATION

Sep 2015 onwards Max-Planck-Institut für Eisenforschung GmbH, Düsseldorf, Germany

Ph.D. in Metallurgical Engineering

Aug 2010- Jan 2013 Indian Institute of Science, Bangalore-India

MSc (Engg) in Mechanical Engineering

GPA 7.5/8.0

Aug 2005- May 2009 National Institute of Technology Karnataka, Surathkal-India

B.Tech in Metallurgical and Materials Engineering

GPA 8.09/10

PROFESSIONAL EXPERIENCE

Jun 2013-Aug 2015 General Electric Co.,

John F Welch Technology Center, Bangalore, India

Systems Engineer

Sep 2009-Jun 2010 National Facility for Semi-solid Forming (NFSSF), Bangalore-India

Project Assistant

PUBLICATIONS

A.R. Kini, D. Maischner, A. Weisheit, E.A Jägle and D. Raabe, A lean Cu-3.4Cr-0.6Nb (at.%) metal matrix

composite produced by laser metal deposition hardened by Cr nanoprecipitates and Laves phase particles,

Acta Materialia, (to be submitted).

110

C.P. Massey, D.T. Hoelzer, R.L. Seibert, P.D. Edmondson, A. Kini, B. Gault, K.A. Terrani, S.J Zinkle,

Microstructural Evaluation of a Nanostructured 12Cr ODS Alloy with Mo, Nb, and Ti additions, Acta

Materialia, (under review).

C.P. Massey, S.N. Dryepondt, P.D. Edmondson, M.G. Frith, K.C. Littrell, A. Kini, B. Gault, K.A. Terrani,

S.J. Zinkle, Multiscale Investigations of Nanoprecipitate Nucleation, Growth and Coarsening in Annealed

low-Cr Oxide Dispersion Strengthened FeCrAl Powder, Acta Materialia, (accepted).

S.K. Makineni, A.R. Kini, E.A. Jägle, H. Springer, D. Raabe, B. Gault, Synthesis and stabilization of a

new phase regime in a Mo-Si-B based alloy by laser-based additive manufacturing, Acta Materialia, 151

(2018) 31-40.

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