laser additive manufacturing of oxide dispersion
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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
anoopkini@gmail.com
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.
6
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
2.1. Laser Additive Manufacturing (LAM)
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
2.1. Laser Additive Manufacturing (LAM)
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
2.1. Laser Additive Manufacturing (LAM)
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
2.2. Laser Additive Manufacturing for Fabrication of Functional and Structural Materials
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].
2.2. Laser Additive Manufacturing for Fabrication of Functional and Structural Materials
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.
2.3. Immiscible Alloy Systems: Fundamentals and Examples
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
2.3. Immiscible Alloy Systems: Fundamentals and Examples
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.
2.4. Alloy Design for Additive Manufacturing
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 +
2.4. Alloy Design for Additive Manufacturing
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.
2.4. Alloy Design for Additive Manufacturing
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].
2.4. Alloy Design for Additive Manufacturing
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.
2.4. Alloy Design for Additive Manufacturing
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.
2.4. Alloy Design for Additive Manufacturing
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.
2.4. Alloy Design for Additive Manufacturing
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.
2.4. Alloy Design for Additive Manufacturing
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.
2.4. Alloy Design for Additive Manufacturing
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.
2.4. Alloy Design for Additive Manufacturing
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].
2.4. Alloy Design for Additive Manufacturing
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. Microstructural Characterization
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.
3.2. Microstructural Characterization
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.
3.2. Microstructural Characterization
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
3.2. Microstructural Characterization
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.
3.2. Microstructural Characterization
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
3.2. Microstructural Characterization
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. Mechanical Property Characterization
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
3.3. Mechanical Property Characterization
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.
4.2. Feedstock Powder Preparation
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
4.2. Feedstock Powder Preparation
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
4.2. Feedstock Powder Preparation
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. Microstructural Characterization of Dense Samples
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.
4.3. Microstructural Characterization of Dense Samples
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
4.3. Microstructural Characterization of Dense Samples
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.
4.3. Microstructural Characterization of Dense Samples
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
4.3. Microstructural Characterization of Dense Samples
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).
4.3. Microstructural Characterization of Dense Samples
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
4.3. Microstructural Characterization of Dense Samples
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.
4.3. Microstructural Characterization of Dense Samples
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.
4.3. Microstructural Characterization of Dense Samples
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.
4.3. Microstructural Characterization of Dense Samples
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.
4.3. Microstructural Characterization of Dense Samples
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
4.3. Microstructural Characterization of Dense Samples
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.
4.3. Microstructural Characterization of Dense Samples
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.%).
4.3. Microstructural Characterization of Dense Samples
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
4.3. Microstructural Characterization of Dense Samples
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-
4.3. Microstructural Characterization of Dense Samples
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.
4.4. Challenges and Comments
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
4.4. Challenges and Comments
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.
4.4. Challenges and Comments
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.
5.2. Microstructural Characterization of Dense Samples
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. Microstructural Characterization of Dense Samples
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
5.2. Microstructural Characterization of Dense Samples
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]
5.2. Microstructural Characterization of Dense Samples
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,
5.2. Microstructural Characterization of Dense Samples
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.
5.2. Microstructural Characterization of Dense Samples
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
5.2. Microstructural Characterization of Dense Samples
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.
5.2. Microstructural Characterization of Dense Samples
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.
5.2. Microstructural Characterization of Dense Samples
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].
5.2. Microstructural Characterization of Dense Samples
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 Hardening Assessment
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.
5.4 Hardening Assessment
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
5.4 Hardening Assessment
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
5.4 Hardening Assessment
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.
5.4 Hardening Assessment
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
5.4 Hardening Assessment
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.
5.4 Hardening Assessment
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.
94
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).
96
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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