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Theses
12-2-2019
Additive Manufacturing Materials: Fabrication of Aluminum Additive Manufacturing Materials: Fabrication of Aluminum
Matrix Composites Matrix Composites
Jakob Hamilton jdh3685@rit.edu
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Additive Manufacturing Materials: Fabrication of
Aluminum Matrix Composites
by
Jakob Hamilton
A Thesis Submitted in Partial Fulfillment of the Requirements for the
Degree of Master of Science in Industrial and Systems Engineering
Advisor: Dr. Iris V. Rivero
Department of Industrial and Systems Engineering
Kate Gleason College of Engineering
Rochester Institute of Technology
Rochester, NY
December 2nd, 2019
i
Committee Members:
Dr. Iris V. Rivero
Department Head and Kate Gleason Professor of Industrial and Systems Engineering
Rochester Institute of Technology
Dr. Denis Cormier
Earl W. Brinkman Professor of Industrial and Systems Engineering and AMPrint Center Director
Rochester Institute of Technology
Approved by:
Dr. Iris V. Rivero Date
Dr. Denis Cormier Date
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Abstract
This study aims to validate the ability of cryomilling for the production of high-quality
aluminum matrix composite (AMC) powder for powder bed fusion (PBF) additive
manufacturing (AM). The spectrum of aluminum-based materials available for AM remains
limited due to complex melting and solidification dynamics inherent to the process. To overcome
these problems, fillers are often added to aluminum matrices to create a class of materials called
AMCs that combine the ductility of aluminum with the stiffness of ceramic reinforcements.
However, producing particulate composite feedstock powder for PBF that promotes full
densification and microstructural homogeneity is nontrivial. Traditional liquid-phase processing
through atomization is not suited to produce composite powders as particle segregation
discourages composite homogeneity. AM powder production through solid-state mechanical
alloying has been studied with limited success, primarily due to poor powder spreadability and
inclusion of lubricants in the alloying process. Cryogenic mechanical alloying, termed
cryomilling, enhances homogeneity between matrix and reinforcement particles by recurrent
fracture and cold welding sans lubricants but remains unexplored for the fabrication of PBF
feedstock powder.
Herein, a method for producing homogeneous, flowable AMC powder designed for PBF
is described in detail. Various compositions, powder masses, and milling times were explored to
tune particle morphology, composition, and composite homogeneity. A representative spreading
test of cryomilled materials qualitatively indicated that distinct cryomilling parameters may
produce powder with comparable spreading characteristics to gas atomized AlSi10Mg, a
common PBF feedstock material. Cryomilled AMCs displayed superior Vickers microhardness
iii
to unmilled AlSi10Mg powder after compression and sintering. This research provides an
indication of cryomilling capabilities to become an effective production method of custom alloy
powder for PBF-AM.
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Acknowledgements
First and foremost, I would like to thank my research advisor Dr. Iris Rivero for her
support and guidance throughout the duration of this thesis project and other important research
projects throughout my time in her team. The opportunities granted and connections established
under her lead are second to none. The encouragement to pursue all research questions led to
several successful and enlightening projects. For that, I thank you.
Second, I would like to thank other faculty and staff in the Department of Industrial and
Systems Engineering at RIT. Thank you to Dr. Denis Cormier and Dr. Bruce Kahn in the AMPrint
Center for the expert guidance throughout the project and allowing my use of the glovebox and
microscopy equipment. Thank you to Rob Kraynik in the Brinkman Machine Tools Laboratory
for the generous gifts of time and knowledge to ensure equipment is both safe and effective.
Thank you to Professor Patricia Cyr for the continuing guidance in crafting professional-level
design of experiments.
I would like to extend a hearty thank you to my colleague Srikanthan Ramesh. Although
our research topics rarely crossed, the wisdom granted to me through relevant questions,
suggestions, and lighthearted banter added to this project’s success. Thank you for your support
throughout my time at ISU and RIT.
I would also like to thank my current and former research colleagues, namely Samantha
Sorondo, Andrew Greeley, Sharon Lau, and Eric Weflen. Collaborative projects and research
meetings with each of you have undoubtedly expedited my learning. The community formed
through our collaboration has enriched my experiences both at RIT and ISU.
I must also give a warm thank you to Dr. Daniel Black and Dr. LeAnn Faidley in the
Engineering Science Department at Wartburg College. The theoretical and practical aspects of
engineering imparted to me through the liberal arts engineering degree are carried with me every
day. Likewise, the opportunity to seamlessly transition into a graduate program kick-started this
thesis project. I have nothing but appreciation for the work you do.
Lastly, I would be remiss if I did not thank my parents Chris and Brenda Hamilton for
their unending love and support. Dad, the mechanical skills you have instilled in me contributed
to my success in this and all other research projects. Mom, you have given me a lifetime of literary
skills that have hastened the researching and writing processes. One paragraph is not nearly
enough words to express my gratitude to both of you adequately. I would also like to thank my
brothers Alexander and Joshua for their support, love, and humor throughout my education.
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Table of Contents
List of Figures .......................................................................................................................................... vii
List of Tables ........................................................................................................................................... viii
Nomenclature ............................................................................................................................................ ix
1. Introduction ............................................................................................................................................ 1
1.1 Research Motivation ........................................................................................................................................ 1
1.2 Problem Statement ........................................................................................................................................... 5
2. Literature Review ................................................................................................................................... 7
2.1 Additive Manufacturing.................................................................................................................................. 7
2.1.1 Additive Manufacturing of Metal ........................................................................................................... 8
2.1.2 Powder Feedstock ..................................................................................................................................... 9
2.1.3 Metallic Materials for Powder Bed Systems ........................................................................................ 10
2.2 Metal Matrix Composites .............................................................................................................................. 13
2.2.1 Strengthening Mechanisms in Particulate MMCs .............................................................................. 15
2.2.2 Metal Matrix Nanocomposites .............................................................................................................. 18
2.2.3 Additive Manufacturing of Metal Matrix Composites ...................................................................... 18
2.3 Current Methods for PBF Composite Powder Fabrication ....................................................................... 21
2.3.1 Gas Atomization...................................................................................................................................... 21
2.3.2 Reaction Synthesis .................................................................................................................................. 22
2.3.3 Surface Coating ....................................................................................................................................... 24
2.3.4 Mechanical Alloying ............................................................................................................................... 26
2.3.5 Cryomilling .............................................................................................................................................. 30
2.4 Summary and Critique of the Chapter ........................................................................................................ 33
3. Preliminary Work................................................................................................................................. 36
3.1 Cryomilling Al-TiC Study ............................................................................................................................. 36
3.1.1 Composite Materials ............................................................................................................................... 36
3.1.2 Cryomilling .............................................................................................................................................. 37
3.1.3 Powder Characterization ....................................................................................................................... 38
3.1.4 Conclusions .............................................................................................................................................. 42
4. Solid-state Synthesis of Aluminum Matrix Composites for Powder Bed Fusion Additive
Manufacturing .......................................................................................................................................... 44
Abstract .................................................................................................................................................................. 44
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4.1 Introduction .................................................................................................................................................... 44
4.2 Materials and Methods .................................................................................................................................. 48
4.3 Results .............................................................................................................................................................. 51
4.3.1 X-ray diffraction (XRD) analysis ........................................................................................................... 51
4.3.2 Morphological Characterization ........................................................................................................... 52
4.3.3 Metallographic Analysis ........................................................................................................................ 55
4.3.4 Powder Spreadability ............................................................................................................................. 58
4.3.5 Microhardness ......................................................................................................................................... 59
4.4 Discussion ........................................................................................................................................................ 60
4.5 Conclusions ..................................................................................................................................................... 63
5. General Conclusions ............................................................................................................................ 65
5.1 Conclusions ..................................................................................................................................................... 65
5.2 Review of Contributions ............................................................................................................................... 66
5.3 Future Perspectives ........................................................................................................................................ 66
References ................................................................................................................................................. 69
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List of Figures
Figure A: Impactor and powder collision in mechanical alloying (Suryanarayana, 2001). ............ 5
Figure B: Custom crystallographic orientation map for nickel used in EBM (Dehoff et al., 2015). 8
Figure C: Increase in strength based on solute content in steel (Meyers & Chawla, 2009). Solid
lines represent substitutional solutes and dashed lines represent interstitial solutes. .................. 17
Figure D: Theorized melt pool dynamics in LBM (Yuan, Gu, & Dai , 2015). .................................. 19
Figure E: Mechanically mixed Ti/nano-TiC composites (D. Gu, Wang, & Zhang, 2014). ............. 24
Figure F: Al7075/nano-TiC composite particle prepared by surface inoculation (Tasche et al.,
2019). .......................................................................................................................................................... 25
Figure G: Nanofunctionalized AlSi10Mg/nano-WC composites (Martin et al., 2018). .................. 25
Figure H: Morphological evolution of Al-Al2O3 composites after (a) 0, (b) 4, (c) 8, (d) 12, (e) 16,
and (f) 20 hours of planetary milling (Han, Setchi, & Evans, 2017). ................................................. 28
Figure I: Szegvari ball mill (Suryanarayana, 2001). ............................................................................. 31
Figure J: Side view of cryomilling vial, powder, and impactor. ....................................................... 31
Figure K: as-received (a) Al and (b) TiC powder images taken using SEM. ................................... 37
Figure L: Images of the (a) outside, (b) inside, and (c) milling vial of the SPEX 6875D
Freezer/Mill®. ........................................................................................................................................... 38
Figure M: SEM images of cryomilled AMC powder after: (a) and (d) 2 hours, (b) and (e) 4 hours,
(c) and (f) 6 hours. .................................................................................................................................... 39
Figure N: Particle size distributions for (a) 2-hour cryomilled Al-TiC, (b) 4-hour cryomilled Al-
TiC, (c) 6-hour cryomilled Al-TiC, and (d) LPW. AlSi10Mg powders ............................................. 40
Figure O: Mean particle size comparison between all cryomilled Al-TiC samples and LPW
AlSi10Mg powder. ................................................................................................................................... 40
Figure P: Example of the original and threshold image for reinforcement distribution
measurement. ........................................................................................................................................... 41
Figure Q: Interval plot of d indices of surface-level reinforcement distribution for individual
milling times. ............................................................................................................................................ 41
Figure R: As-received (a) AlSi10Mg powder from LPW-Carpenter Additive and (b) nanoscale
TiC powder from Sigma-Aldrich. .......................................................................................................... 48
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Figure S: Depiction of the cryomilling vial, impactor, and direction of impactor travel. The vial
is submerged in liquid nitrogen during the milling process. ............................................................ 49
Figure T: XRD diffractograms for as-received AlSi10Mg powder and L-10TiC and H-1TiC after 3
hours of cryomilling. ............................................................................................................................... 52
Figure U: SEM images of all samples throughout cryomilling. Red circles denote agglomerated
fine particles. ............................................................................................................................................. 53
Figure V: Porous agglomerates comprising L-1TiC after 6 hours of cryomilling. .......................... 54
Figure W: Comparison between particle size and aspect ratio for (a) cryomilling time, (b) IPR,
and (c) TiC loading. ................................................................................................................................. 55
Figure X: Powder cross sections of (a) L-10TiC and (b) H-10TiC after 3 hours of cryomilling. The
red rectangles outline incompletely welded particles. ....................................................................... 56
Figure Y: EDS images of cryomilled AMC powder displaying TiC distribution for distinct
treatment combinations........................................................................................................................... 57
Figure Z: SEM and elemental mapping of a large particle agglomerate from L-10TiC after 3
hours of cryomilling illustrating the primary and second phase striations. ................................... 57
Figure AA: Spread patterns of (a) L-1TiC-3h, (b) L-1TiC-6h, (c) H-10TiC-3h, (d) H-10TiC-6h, and
(e) as-received AlSi10Mg powder. ......................................................................................................... 58
Figure BB: Vickers microhardness measurements for the 3-hour cryomilled samples. Error bars
denote standard deviation. ..................................................................................................................... 59
List of Tables
Table 1: Metals and alloys available for PBF-AM. ............................................................................... 11
Table 2: MMC fabrication studies using AM. ...................................................................................... 20
Table 3: Design of experiment for the AlSi10Mg-TiC cryomilling study. “L” and “H” represent
low and high powder mass, respectively. ............................................................................................ 49
Table 4: Particle size distributions of the cryomilled samples at all milling times. ........................ 54
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Nomenclature
Al4C3 Aluminum Carbide
AlN Aluminum Nitride
Al2O3 Aluminum Oxide
AM Additive Manufacturing
AMC Aluminum Matrix Composite
CAD Computer Aided Design
CTE Coefficient of Thermal Expansion
DED Directed Energy Deposition
EBM Electron Beam Melting
EDS Energy-Dispersive X-ray Spectroscopy
Fe2O3 Iron Oxide
HIP Hot Isostatic Pressing
IPR Impactor-to-Powder Mass Ratio
LBM Laser Beam Melting
LMD Laser Metal Deposition
MMC Metal Matrix Composite
MMNC Metal Matrix Nanocomposite
PBF Powder Bed Fusion
PCA Process Control Agent
SEM Scanning Electron Microscopy
SiC Silicon Carbide
x
TiC Titanium Carbide
WC Tungsten Carbide
XRD X-ray Diffraction
1
1. Introduction
1.1 Research Motivation
Additive manufacturing (AM) is the process of consolidating materials to make objects
directly from 3D model data. AM’s versatility in producing complex parts with minimal process
planning, tooling, and waste make it ideal for fields where traditional subtractive and formative
manufacturing are inefficient. AM also allows for recycling of unused material (Manfredi et al.,
2014). Moreover, the lack of dies, fixtures, and other tooling makes AM ideal for custom, small-
batch applications such as biomedical implants or prototyping (Sames, List, Pannala, Dehoff, &
Babu, 2016). Further applications of AM range from complex engine blades and vanes for jets and
high-temperature fuel injectors for NASA rockets (L. J. Kumar & Krishnadas Nair, 2017; Manfredi
et al., 2014).
Despite these advantages, several limitations remain. Compared to traditional
manufacturing methods, processing speed, surface roughness, and the capital cost of machines
limit AM’s applicability (Manfredi et al., 2014). Diversity in metal materials for AM also remains
an issue as only a small fraction of the 5,500 alloys for traditional manufacturing methods are
available for metal AM (Manfredi et al., 2014). Hot cracking and large residual stresses render
6000- and 7000- series aluminum alloy components unusable after manufacturing. Likewise,
corrosion-inhibiting elements such as magnesium and zinc in traditional aluminum alloys are
known to vaporize during high-temperature processing causing altered composition and
porosity (Herzog, Seyda, Wycisk, & Emmelmann, 2016; Ly, Rubenchik, Khairallah, Guss, &
Matthews, 2017; Matthews et al., 2016). Aluminum’s high reflectivity at laser wavelengths
2
between 1 μm and 20 μm also poses heat-absorption issues in laser-based AM. Development of
lightweight aluminum alloys for metal AM lags both titanium alloys and nickel-based
superalloys.
Composite materials are one novel alternative, but their potential in metal AM remains
relatively unexplored. Metal matrix composites (MMCs) feature a metal matrix comprising the
bulk of the material, and one or more reinforcement materials in laminate, fibrous, or particulate
form. Reinforcement materials may be metallic or ceramic and often have stronger mechanical
properties than the matrix material. Combining a ductile metal matrix with hard second phase
particles allows particulate MMCs to possess higher mechanical, thermal, and fatigue properties
(Chawla & Shen, 2001).
Enhanced properties in MMCs are linked to several important physical phenomena.
Reduction in grain size has been shown to increase mechanical properties according to the Hall-
Petch relationship (Chawla & Chawla, 2013). Reinforcement size reduction also attributes to the
Orowan bowing mechanism where strength is inversely proportional to interparticle spacing
(Aversa et al., 2017). Several studies have been able to refine crystallite structure and decrease
interparticle spacing with inclusion of nanoscale reinforcement particles (Aversa et al., 2017; D.
Gu, Wang, Chang, et al., 2014; Marchese et al., 2018; Martin et al., 2017; H. Wang & Gu, 2015).
Energy absorptivity may also be tailored through the inclusion of ceramic reinforcement particles,
reducing reflectivity issues in aluminum alloys.
The potential strengthening in MMCs is attractive, however, these heightened properties
rely on two important MMC characteristics: the surface interaction between phases and
distribution homogeneity of the reinforcement phase within the matrix (Chawla & Chawla, 2013).
3
Poor wettability between matrix and reinforcement can lead to interfacial stresses and porosity,
exacerbating the residual stresses developed during AM solidification. Similarly, reinforcement
dispersion contributes to homogeneity within microstructures and heat absorptivity. Obtaining
a uniform reinforcement dispersion state can be difficult, especially with the large Van der Waals
forces between sub-micron reinforcement particles (Cabeza et al., 2017; D. Gu, Wang, Chang, et
al., 2014).
For MMCs to be feasible in AM, composite powders must retain several important
characteristics. In powder bed fusion (PBF), powder flow characteristics and packing density are
crucial and dependent on particle size distribution and individual particle morphology (D. D.
Gu, Meiners, Wissenbach, & Poprawe, 2012). Several researchers have successfully implemented
lightweight aluminum matrix composites (AMCs) into laser-based AM processes through a
variety of feedstock powder production methods (Han, Setchi, Lacan, Gu, & Evans, 2017; Martin
et al., 2018; Tasche et al., 2019; H. Wang & Gu, 2015).
Recently, traditionally unweldable aluminum alloy 7075 was successfully arc welded in a
tungsten inert gas system using AMC filler rods containing Al7075 and titanium carbide (TiC)
nanoparticles (Sokoluk, Cao, Pan, & Li, 2019). Inclusion of nanoscale particles was theorized to
encourage the growth of equiaxed grains and decrease the residual stresses known to instigate
hot cracking. The author of the present work believes that nanoscale-reinforced aluminum matrix
composites can be successfully applied to PBF to expand available AM materials. Developing
lightweight metal materials for additive manufacturing would allow for new applications of
metal AM in both biomedical and aerospace sectors.
4
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1.2 Problem Statement
While MMCs prove to be an attractive option for mitigating thermal issues in PBF, there
are several hurdles in developing feasible composite materials. MMC powders must possess
homogeneous reinforcement distribution and good flowability which is usually achieved by a
particle size distribution between 10 μm and 100 μm and near-spherical particle morphology.
Typically, gas atomization provides the ideal size distribution and morphology for single-phase
AM powders, however multi-phase powders are difficult to produce via spray atomization
owing to the high cost of equipment and difficulty in achieving homogeneous reinforcement
distribution (Chawla & Chawla, 2013).
Instead, mechanical alloying is commonly selected for powder production. Mechanical
alloying utilizes solid impactors and oscillatory motion to impart severe plastic deformation on
the subject material as depicted in Figure A. Cycles of particle fracture and cold welding in this
process encourages consistent characteristics within the resultant material (Suryanarayana, 2001).
One novel alternative is mechanical alloying at cryogenic temperatures, or cryomilling.
Cryomilling has been shown to enhance distribution homogeneity of second phase particles,
suppress microstructure recovery and recrystallization,
decrease average grain size, and prevent welding to
milling equipment (Enayati, 2017; Groza, Lavernia,
Shackelford, & Powers, 2007; Suryanarayana, 2001).
In order to minimize residual stresses generated
in metal AM consolidation, nucleation-inducing
reinforcement particles should be homogeneously Figure A: Impactor and powder collision in
mechanical alloying (Suryanarayana, 2001).
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dispersed within the PBF melt pool. The combination of cryomilling and inclusion of nanoscale
reinforcement particles is expected to develop consistent particle morphology, enhance wetting
between phases, enhance nucleation during AM consolidation, and reduce the average crystallite
size in AMC powders. Thus, this research seeks to fulfill the following objectives:
Develop cryomilling parameters to fabricate AMC powders with homogeneous
reinforcement distribution, equiaxed morphology, and critical particle size distribution
Assess feasibility of cryomilled AMC powders for PBF by relating powder properties to
flowability and densification
With key processing parameters for AMC powders unveiled, new materials can easily be
developed, and individual properties can be tailored for each application. Though manipulation
of cryomilling and AM processing parameters, factors such as reinforcement weight fraction,
microstructural characteristics, and particle size distribution can all be modified to tailor strength,
thermal, and fatigue properties of AM-fabricated AMC components. New customizable materials
for AM will broaden its utility within biomedicine, aerospace, other commercial sectors.
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2. Literature Review
In this section, the status of metal materials for additive manufacturing and common
production methods are explained. To begin, an overview of the principles of metal additive
manufacturing is given. Following, current metal AM composites and production techniques are
discussed in depth.
2.1 Additive Manufacturing
Additive manufacturing (AM) is the fabrication of components through layer-based
material consolidation (ASTM, 2012). There are seven unique processes that comprise AM: binder
jetting, directed energy deposition, material extrusion, material jetting, powder bed fusion, sheet
lamination, and vat photopolymerization. Metals, polymers, ceramics, and composite materials
are all available for AM processes. Additive manufacturing offers certain advantages compared
to traditional subtractive or formative techniques:
1. Component complexity: AM produces complete parts or assemblies to near-net shape
directly from Computer Aided Design (CAD) software with minimal process
planning and tooling (Manfredi et al., 2014).
2. Tailored strength: Parametric control of the growth rate (R) and thermal gradients (G)
in AM allows custom microstructures and mechanical properties for various
applications (Dehoff, Kirka, List, Unocic, & Sames, 2015). A 2014 study at Oak Ridge
National Laboratory (Oak Ridge, TN, USA) produced custom microstructures using
PBF as seen in Figure B (Dehoff et al., 2014).
8
3. Material conservation: AM wastes less
raw material compared to subtractive
manufacturing, which is especially
important for highly expensive
aerospace materials (L. J. Kumar &
Krishnadas Nair, 2017). Moreover,
leftover material in AM processes may be sieved and reused or recycled after the
process is complete.
The lack of extensive tooling and wastages in AM makes it ideal for prototyping
applications. Further applications range from biomedical implants, complex engine blades and
vanes for jets, high-temperature fuel injectors for NASA rockets, and even complete rocket engine
components (L. J. Kumar & Krishnadas Nair, 2017; Manfredi et al., 2014; Sames et al., 2016).
2.1.1 Additive Manufacturing of Metal
Of the seven AM processes, three are commonly used to fabricate metal components:
binder jetting, direct energy deposition (DED), and powder bed fusion (PBF) (Manfredi et al.,
2014). Binder jetting uses an adhesive to selectively join metal powder, and post-processing heat
treatment is necessary to catalyze the strengthening secondary metallic phase. Both DED and PBF
utilize a directed heat source to join metal particles without the use of additional binding agents.
The most popular technologies are laser metal deposition (LMD) and wire arc additive
manufacturing (WAAM) for DED and laser beam melting (LBM) and electron beam melting
(EBM) for PBF. In LMD, wire or powder stock is expelled directly onto the build plate under a
high-temperature laser. In LBM and EBM systems, powder is evenly distributed across each layer
Figure B: Custom crystallographic orientation map for
nickel used in EBM (Dehoff et al., 2015).
9
by a blade or roller, and the beam selectively joins particles above their melting temperature.
LMD and LBM are typically performed in an inert environment to avoid oxidation and inhibit
flammability, while EBM is performed in a near-vacuum to minimize oxidation and unwanted
electron scattering.
Despite the benefits, metal AM retains several drawbacks. Most notable limitations
include slow speeds, poor surface finish, high residual stresses, high equipment and material
costs, and limited available materials (Manfredi et al., 2014). Optimal AM process parameters
such as power and scanning speed are still under development for numerous materials, and very
few lightweight alloys or composites are available for this reason. Future developments in AM
should focus on creating new, feasible feedstock powders and determining appropriate process
parameters for application-specific components.
2.1.2 Powder Feedstock
In powder LMD, LBM, and EBM systems, the feedstock is important for final component
characteristics. Sames et al. identified that feedstock morphology, flowability, and size
distribution directly influence porosity and material chemistry and indirectly influence
microstructural evolution and the resultant mechanical properties (Sames et al., 2016). Powder
sphericity is especially important in PBF systems, as the powder delivery methods are typically
more reliant on the preferable spreadability of spherical powders. Near-spherical gas atomized
powders distribute more evenly across build platforms, reducing the risk of voids or inclusions,
thus making them preferable for PBF (Tan, Wong, & Dalgarno, 2017b). Likewise, particle size
distributions between 10µm and 100µm are preferred to ensure efficient packing density without
inhibiting powder flowability. Currently, no standard quantifies powder morphology, but
10
qualitative analysis may be achieved by analyzing images and using light scattering techniques
(ASTM, 2014). Other characterization standards through ASTM exist for testing flow rate,
apparent density, particle size distribution, and chemical composition of metal powder. Pertinent
topics still under development for metal AM feedstock powder include powder recyclability,
melt pool dynamics, and metal vaporization (Lewandowski & Seifi, 2016; Ly et al., 2017).
2.1.3 Metallic Materials for Powder Bed Systems
Metals, alloys, and metal matrix composites for PBF remain under development. Complex
solidification dynamics in metal AM render only a small subset of the over 5,500 alloys available
for AM (Martin et al., 2017). For instance, only five metal powders have been validated by Arcam
(Arcam EBM, Sweden) for commercial use in their EBM system: Ti6Al4V, Ti6Al4V ELI, Titanium
Grade 2, Arcam ASTM F75 Cobalt-Chrome, and nickel-based Alloy 718. A brief list of metallic
materials studied thus far in PBF is shown in Table 1.
11
Table 1: Metals and alloys available for PBF-AM.
Process Base Material Alloy Reference
LBM Copper --- (Herzog et al., 2016; Manfredi et al., 2014) Gold --- (Manfredi et al., 2014)
Silver --- (Herzog et al., 2016)
Titanium (grade 2) --- (Herzog et al., 2016)
Aluminum AA6061 (Manfredi et al., 2014)
--- AlMg1SiCu (Herzog et al., 2016)
--- AlMg4.4Sc0.66MnZr (Herzog et al., 2016)
--- AlSi10Mg (Manfredi et al., 2014; Mower & Long, 2016)
--- AlSi12 (Manfredi et al., 2014)
--- AlSi12Mg (Manfredi et al., 2014)
Nickel Inconel 625 (Herzog et al., 2016; Manfredi et al., 2014)
--- Inconel 718 (Herzog et al., 2016; Manfredi et al., 2014)
Steel Hot-work steel (Manfredi et al., 2014)
--- Stainless steel 17-4PH (Herzog et al., 2016)
--- Stainless steel 304L (Herzog et al., 2016)
--- Stainless steel 316L (Herzog et al., 2016; Manfredi et al., 2014)
--- Martensitic steel (Manfredi et al., 2014)
--- Tool steel (Manfredi et al., 2014)
--- Maraging steel (Manfredi et al., 2014)
Titanium Ti6Al4V (Herzog et al., 2016; Manfredi et al., 2014)
--- Ti6Al7Nb (Manfredi et al., 2014)
EBM Titanium (grade 2) --- (Herzog et al., 2016)
Nickel Inconel 625 (Herzog et al., 2016)
--- Inconel 718 (Herzog et al., 2016)
Titanium Ti6Al4V (Herzog et al., 2016; Manfredi et al., 2014)
Nickel-based superalloys such as Inconel 718 have been studied in both LBM and EBM
for their use in high-temperature applications as turbine blades and vanes (Dehoff et al., 2015).
Titanium alloys like Ti-6Al-4V and Ti-6Al-7Nb are particularly popular in both LBM and EBM
for biomedical and aerospace applications due to high machining costs for titanium (Herzog et
al., 2016; Manfredi et al., 2014). Titanium alloys have been particularly interesting to study due to
the clear relationship between AM process parameters and the resultant microstructure. Large
thermal gradients, complex thermal cycles, and alloy composition influence final microstructures,
composition, and material properties. Aluminum is not as widely studied as titanium or steels,
however, largely due to complications during melting and solidification. Recently, an Al-Mg-Sc
12
alloy, termed Scalmalloy®, was developed specifically for LBM. Other aluminum alloys and
composites have yet to be studied for LBM. For EBM, aluminum alloys and composites remain
widely unavailable.
Aluminum alloys retain several limitations rendering them impractical in AM. For one,
simple aluminum parts are comparatively easy to machine, so AM may be unideal. Additionally,
many aluminum alloys deemed unweldable are also unable to be consolidated in AM. Hot
cracking renders 6000 and 7000 series aluminum alloy components unusable after manufacturing
(Martin et al., 2017). Likewise, volatile elements such as aluminum, magnesium, and zinc are
known to vaporize when subjected to high temperatures in a near-vacuum environment, causing
unwanted porosity, melt pool turbulence, and altered chemical composition (Herzog et al., 2016;
Ly et al., 2017; Matthews et al., 2016). Aluminum’s high reflectivity at laser wavelengths also
poses heat-absorption issues in LBM machines (Herzog et al., 2016). In short, development of
AM-feasible aluminum alloys for aerospace applications lags steel, titanium, and nickel alloys.
Many unweldable alloys suffer from solidification cracking caused by columnar grain
growth during solidification. After melting, the primary equilibrium phase solidifies first at a
different composition than the bulk liquid (Martin et al., 2017). Solutes enrich the bulk liquid,
causing both a change in the liquidus temperature and growth of unstable undercooled regions.
The variance in composition between the solid primary region and the bulk liquid region causes
poor wetting and columnar grain growth. Columnar grains are especially prone to thermal
contraction which promotes hot tearing and solidification cracking through intergranular cavities
and print layers leading to undesirable anisotropic strength properties.
13
To cope with poor solidification and absorption properties, current studies are aimed at
generating fine equiaxed microstructures in aluminum using nanoparticle refiners (Karthik,
Panikar, Ram, & Kottada, 2017; Martin et al., 2017; Sokoluk et al., 2019). Heat-absorbing
nanoparticles act as nucleation sites during solidification, reducing the required rate of
undercooling for equiaxed grain growth to occur (Martin et al., 2017). The equiaxed grain
structure is better at suppressing thermal contraction and retaining strength than columnar
microstructures. Manipulating AM process parameters to promote significant undercooling may
also be used to encourage equiaxed grain growth. Specifically, low thermal gradients (G) and
high growth rates (R) from slower beam travel encourage equiaxed grain growth in the processed
alloy (Kou, 2003).
In short, metal AM shows promising results for the future of design-based manufacturing.
Currently, major limitations include the lack of high-performance, lightweight alloys, optimal
process parameters, and porosity. Although steel, titanium, and nickel alloys have been reliably
produced for AM, aerospace-grade aluminum alloys are scarce for LBM and EBM processing.
Solidification constraints eliminate many aluminum alloys, and ones currently available for AM
maintain strength but lack ductility and corrosion resistance. Developing aluminum-based alloys
with reliable solidification and strength properties would advance metal AM’s applicability.
2.2 Metal Matrix Composites
Metal matrix composites (MMCs) are materials with at least two distinct phases: a metal
matrix which comprises the bulk of the material, and one or more reinforcement materials in
laminate, fibrous, or particulate form. Reinforcement materials may be metallic or ceramic and
14
often have stronger mechanical properties than the matrix material (Suryanarayana & Al-Aqeeli,
2013). MMCs offer several enhancements over unreinforced materials:
1. Mechanical properties: Reinforcement particles cause dislocation strengthening by
increasing the frequency of dislocations within the metal matrix (Chawla & Chawla, 2013).
Likewise, fine matrix grains lead to grain boundary strengthening governed by the Hall-
Petch principle.
2. Thermal capabilities: Reinforcement particles constrain matrix flow properties, increasing
thermal behavior and creep resistance (Chawla & Chawla, 2013). In practice, nickel-based
superalloys have been found suitable for gas turbine blades due to their superior high-
temperature creep resistance.
3. Fatigue behavior: High-stiffness ceramic reinforcement particles have been shown to
increase fatigue resistance of pure metals (Chawla & Chawla, 2013). It has also been
shown that increasing particle volume fraction and decreasing particle size and
agglomeration result in higher fatigue resistance in MMCs.
The microstructures within the MMCs govern the magnitude of these enhancements
(Jayalakshmi & Gupta, 2015). Despite the elevated properties, MMCs possess drawbacks:
1. Machinability: MMCs are often difficult to machine due to their low ductility and fracture
toughness.
2. Matrix and reinforcement interface: Discrepancies in elastic modulus and coefficient of
thermal expansion (CTE) between the matrix and reinforcement can cause undesirable
porosity and interfacial stresses (Chawla & Chawla, 2013). Additionally, poor wettability
and density mismatch between MMC constituents create low fracture toughness.
15
3. Low ductility: While MMCs possess higher strength than alloys, they often suffer from
extreme brittleness without post-processing heat treatment (Jayalakshmi & Gupta, 2015).
Ductility has also been inversely correlated to reinforcement particle size, i.e. ductility
increases with decreasing particle size (Suryanarayana & Al-Aqeeli, 2013).
4. Inhomogeneous reinforcement distribution: Large Van der Waals forces between
submicron particles can cause agglomeration of reinforcement particles. MMC properties
are directly influenced by reinforcement distribution with a homogeneous distribution
being preferable (Aversa et al., 2017; Jayalakshmi & Gupta, 2015).
Despite the drawbacks, the enhanced properties compared to traditional alloys make MMCs
suitable for high-performance applications such as Co-WC in cutting tools, Al-SiC in Boeing 777
engine components, and Ti-SiC in nozzle actuator controls in the F-16 aircraft (Chawla & Chawla,
2013). MMCs in the remainder of this work specifically refer to metal matrix composites with
discontinuous ceramic reinforcements.
2.2.1 Strengthening Mechanisms in Particulate MMCs
Several strengthening mechanisms may be used predict mechanical properties based on
MMC properties. The primary mechanisms are precipitation hardening, grain size strengthening,
and solid solution strengthening:
1. Precipitation hardening - The addition of precipitates into a matrix encourages the
Orowan mechanism which provides additional strength and reduces the tendency for hot tearing.
The Orowan mechanism is due to the “bowing” of dislocations around non-deforming
precipitates in a matrix and can be expressed with the governing equation:
16
𝜎𝑜𝑟 =𝐺 ∗ 𝑏
𝐿
where 𝐺 is the shear modulus, 𝑏 is the magnitude of the Burgers vector, and 𝐿 is the distance
between precipitates (Gladman, 1999). As seen in the above equation, stress is inversely
proportional to interparticle spacing. For a solution with a fixed volume of reinforcement
material, particle spacing 𝐿 decreases as reinforcement particle diameter decreases. In other
words, as reinforcement material is smaller and more evenly dispersed within the matrix, the
reinforcement particle spacing L decreases. Therefore, resistance of dislocation motion is
increased with finer precipitates in the matrix (Aversa et al., 2017). Coincidentally, decreasing
reinforcement particle size and increasing processing temperature have also been shown to
increase wettability between primary and secondary phases in MMCs (Chawla & Chawla, 2013).
2. Grain size strengthening - Several studies have aimed at increasing tensile strength
through refined grain sizes within MMCs (Aboulkhair, Tuck, Ashcroft, Maskery, & Everitt, 2015;
Aversa et al., 2017; Gupta et al., 2015; H. Wang & Gu, 2015; Xu et al., 2015). The primary
mechanism for grain size strengthening is given by the Hall-Petch relationship:
𝜎𝑔𝑏 =𝑘
√𝑑
where 𝜎𝑔𝑏 is grain boundary strengthening, 𝑘 is the material-specific Hall-Petch coefficient, and
𝑑 is the average grain size (Chawla & Chawla, 2013). The inverse relationship between average
grain size 𝑑 and strengthening indicate stronger components for decreasing grain sizes. It should
be noted that a critical limit exists in the Hall-Petch relationship. As grains become very fine (<10
nm), grain boundaries begin sliding past each other, eliminating strength enhancements from
small grain boundaries.
17
3. Solid solution strengthening – During MMC consolidation, high temperatures may
cause some of the reinforcement particles to coalesce into the matrix, forming a solid alloy phase
that restricts dislocation motion (Chawla & Chawla, 2013). Additional strength from solutes can
be calculated with the equation:
𝜎𝑠𝑠 = 𝐻𝐶𝛼
where 𝐻 and 𝛼 are constants and 𝐶 is the concentration of solute atoms in at.% (Gupta et al., 2015).
The choice of materials in MMCs and process parameters affect the type of solute formation and
the resultant strengthening effect (Meyers & Chawla,
2009). An example of solute strengthening in steel
alloys can be seen in Figure C.
Initial studies have indicated the preference for
smaller, homogeneously distributed reinforcement
particles, but further investigation should be completed
(Aversa et al., 2017). In addition to better understanding
the strengthening mechanisms, obtaining a
homogeneous distribution of reinforcement particles continues to challenge MMC production
methods (Cabeza et al., 2017). Particle clustering caused by extreme size differences between
matrix and reinforcement particles and high surface energy diminishes the strength and ductility
benefits from sub-micron reinforcements. Ultrasonic dispersion has achieved limited success
with homogeneous particle distribution, but poor wetting between matrix and reinforcement
continues to prevent desirable results.
Figure C: Increase in strength based on solute
content in steel (Meyers & Chawla, 2009). Solid
lines represent substitutional solutes and dashed
lines represent interstitial solutes.
18
2.2.2 Metal Matrix Nanocomposites
Typical reinforcement particles for MMCs are on the micrometric scale. MMCs with
nanometric reinforcements, or metal matrix nanocomposites (MMNCs), are also a topic of recent
interest for their increased tensile characteristics (Cabeza et al., 2017). By decreasing
reinforcement particles sizes in the metal matrix to less than 100 nm, several studies have been
able to add tensile strength and ductility to MMCs (Cabeza et al., 2017; Chawla & Chawla, 2013;
D. Gu, Wang, & Zhang, 2014; Jayalakshmi & Gupta, 2015; Mazahery, Shabani, Alizadeh, & Tofigh,
2013; Torabi & Ebrahimi-Kahrizsangi, 2012; H. Wang & Gu, 2015). It is proposed that enhanced
mechanical properties are due to the Orowan strengthening mechanism brought forth by the
addition of fine, hard particles into the matrix microstructure. Still, critical issues such as
inhomogeneous dispersion and poor wettability limit processing methods (Jayalakshmi & Gupta,
2015; Suryanarayana & Al-Aqeeli, 2013). It is suggested that solid-state methods are the most
appropriate form of production for metal matrix nanocomposites (Suryanarayana & Al-Aqeeli,
2013).
2.2.3 Additive Manufacturing of Metal Matrix Composites
Addition of ceramic grain refiners into crack-susceptible materials is not a new concept.
Ceramic grain refiners or inoculants have resolved cracking in as-cast aluminum alloys for over
half a century (Quested, 2004). Welding operations have also found success using metallic and
ceramic inoculants to promote heterogeneous nucleation and equiaxed grain growth (Kou, 2003).
Recently, aluminum matrix composite (AMC) filler rods containing nanoscale TiC were used to
resolve weld cracking in notoriously unweldable Al7075 (Sokoluk et al., 2019). Only within recent
19
years have nano-ceramic grain refiners been applied in AM to mitigate phase segregation and
cracking in aluminum alloys (Martin et al., 2017).
Metal matrix composites have found utility in PBF but add new complications to the
tailoring process as MMC properties must be considered. Particle shape and size distribution
affect flowability and densification, and the distribution of reinforcement particle influences
microstructural evolution. Molten liquid flow driven by Marangoni convection within the AM
melt pool as depicted in Figure D influences reinforcement segregation, grain refinement, and the
resulting crack resistance (Yuan, Gu, & Dai, 2015). Additionally, matrix and reinforcement
wettability can affect intermetallic formation and solid
solution strengthening. Producing MMC powders with
these desired properties can be challenging, and
subsequent sections discuss the benefits and drawbacks
of various powder production methods.
Because of the challenges within MMC powder
production and AM process parameters, few MMCs have been used in PBF. Examples of MMCs
manufactured using AM can be found in Table 2.
Figure D: Theorized melt pool dynamics in LBM
(Yuan, Gu, & Dai , 2015).
20
Table 2: MMC fabrication studies using AM.
Method Matrix Reinforcement Reference
LBM A357 Al2O3 (Aversa, 2017)
A357 TiB2 (Aversa, 2017)
AA6063 TiB2 (Aversa, 2017)
Al4.5Cu3Mg SiC (Manfredi, 2014)
Al4SiC4 SiC (Chang, 2015)
AlMg SiC (Kumar, 2010)
AlSi SiC (Kumar, 2010)
AlSi10Mg TiB2 (Aversa, 2017)
AlSi10Mg TiC (Wang, 2015)
AlSi10Mg MgAl2O4 (Aversa, 2017)
Co WC (Kumar, 2010)
Fe Graphite (Kumar, 2010)
Fe SiC (Kumar, 2010)
Ti Graphite/diamond (Kumar, 2010)
Ti SiC (Kumar, 2010)
EBM NiBSi WC (Peng, 2016)
MMC powder consolidation using metal AM is a relatively new field. Complex melting
and solidification dynamics in AM can cause unfavorable properties like anisotropy, porosity,
reinforcement aggregation, increased grain size, and altered chemical composition in composites
just as in metals or alloys (Yuan et al., 2015). Without homogeneous reinforcement distribution,
solidification cracking renders many MMCs unusable in AM. Elemental vaporization, especially
in EBM’s vacuum environment, eliminates desirable properties. These challenges must be
overcome via new processing routes or material combinations for MMCs to become mainstream
in AM.
In summary, MMCs are becoming desirable in AM processing due to their elevated
mechanical properties, enhanced thermal properties, and manufacturability in powder form.
Several MMCs have been developed for LBM processing, but reflectivity, solidification cracking,
poor ductility, and element vaporization remain obstacles. MMCs designed for PBF remain
significantly less studied than for other powder metallurgy techniques, and the mechanisms
21
behind solidification cracking and vaporization have yet to be mitigated. Recent studies support
nanometric reinforcement particles in MMCs to increase mechanical behavior, but
inhomogeneous reinforcement dispersion using traditional powder production methods persists.
Research should focus on the ability to produce MMCs with ideal particle morphology and size
distribution for AM while obtaining a homogeneous sub-micron reinforcement particle
distribution.
2.3 Current Methods for PBF Composite Powder Fabrication
Several production methods have been used to create MMC powders for PBF-AM, but
each method produces vastly different shapes and compositions of powder, yielding different
properties in final AM components. This section seeks to define the principle advantages and
limitations of each method.
2.3.1 Gas Atomization
Due to its ability to produce spreadable powders, gas atomization is typically chosen for
AM powder production. Molten metal, rapidly cooled by high pressure inert gas, forms near-
spherical powders that achieve the desired flowability and apparent density in PBF. Gas
atomization is also preferable when producing powder from volatile elements as the inert gas
stream prevents oxidation. The choice of inert gas has been shown to affect microstructure of both
the as-atomized powders and the AM-fabricated component (Herzog et al., 2016). Water and
plasma atomization are also popular alternatives for producing AM powder but are limited due
to oxidation and process cost, respectively.
22
Gas atomization is not normally preferable for creating MMC powders, however. To
envelop reinforcement particles within the matrix powders, ceramic particles must be injected
into the atomization spray stream in a process termed spray co-deposition. With accurate and
precise injection, composite homogeneity may be produced, but incorrect ceramic injection
results in reinforcement settling and composite inhomogeneity, especially with nanoscale
reinforcement particles (Chawla & Chawla, 2013). High equipment costs further limit spray co-
deposition’s applicability in producing homogeneous MMC powders.
2.3.2 Reaction Synthesis
In situ reinforcement formation, termed reaction synthesis, has been used to generate
MMCs and MMNCs via casting and other traditional manufacturing routes. With AM’s higher
control of heat input and solidification parameters, PBF is an ideal candidate to control
intermetallic and ceramic formation in alloys, thus generating MMCs.
Dadbakhsh and Hao achieved nanoscale intermetallic and oxide reinforcement formation
in LBM with mixtures of Al, AlMg1SiCu, and AlSi10Mg with Fe2O3 (Dadbakhsh & Hao, 2012). A
homogeneous dispersion state for particulate phases coupled with the high cooling rate in LBM
led to heterogeneous nucleation and a resultant increase in hardness for all alloy compositions.
The authors identified solid solution strengthening from Al-Fe intermetallics and aluminum
oxide (Al2O3) as a significant strengthening mechanism. However, poor wettability between
phases and hydrogen gas entrapment significantly contributed to undesirable part porosity.
The Al/SiC chemistry and reaction interface in LBM was also studied in terms of SiC
particle size by Chang et al. (Chang, Gu, Dai, & Yuan, 2015). Coarse (D50 = 50 μm) SiC particles
blended with AlSi10Mg powders were less prone to react and form a tertiary Al4SiC4 phase than
23
fine (D50 = 5 μm) SiC particles during LBM. Complete melting of the fine SiC particles contributed
to both Al4SiC4 synthesis, melt pool stability, and 97% relative density. Inhomogeneous heat
absorption with coarse SiC powders contributed to severe balling of the melt pool and an 85%
relative density. The finely reinforced AMCs exhibited a 72% increase in HV0.1 microhardness and
66% decrease in wear rate compared to the coarse reinforcements. Increased densification, Al4SiC4
formation, and wettability demonstrate the preference for decreased reinforcement particle sizes
for processing AMCs via PBF.
Not all in situ reactions are preferable, however. Ocelík et al. observed degradation of SiC
in an aluminum matrix during LBM (Ocelík, Vreeling, & De Hosson, 2001). Interfacial formation
of aluminum carbide (Al4C3) plates led to poor wetting, degradation of mechanical properties,
and crack initiation. Composites containing Al4C3 are also notorious for poor corrosion resistance,
further limiting its usefulness.
Cormier et al. also reported difficulty controlling phases while attempting to generate
high-strength γ-TiAl intermetallic compounds from titanium alloys via EBM (Cormier,
Harrysson, Mahale, & West, 2007). Mechanically alloyed titanium and aluminum powders
demonstrated controlled reaction synthesis of a less preferable TiAl3 phase. Reaction synthesis of
prealloyed Ti-Al-Nb-Cr-Fe powder generated net shapes with α2- and γ-TiAl phases, but
vaporization of aluminum decreased the amount of the desirable γ-TiAl phase. Further work in
this field suggests reduced vaporization and enhanced phase control may be achieved through
modification of LBM and EBM parameters (Biamino et al., 2011; Gussone et al., 2015; Murr et al.,
2010; Tang et al., 2015).
24
Reaction synthesis has proven to be an effective method for producing homogeneous
alloys in PBF, but consistent MMC production via reaction synthesis remains a challenge. AM
process parameters greatly affect reinforcement formation and wettability between phases.
Further work should elucidate the effects of energy density and feedstock powder parameters on
the resultant porosity, microstructure, and mechanical properties of AM-synthesized MMCs.
2.3.3 Surface Coating
Several researchers have successfully produced MMC composites in PBF by coating metal
particles with ceramic reinforcement particles, and various methods have been employed.
Mechanical mixing or blending has been used extensively but with limited success. Aversa et al.
produced AlSi10Mg composites with nano- and micro-TiB2 reinforcements using a ball mill
without grinding media (Aversa et al., 2017). AM-consolidated AlSi10Mg/nano-TiB2 composites
exhibited a correlation between LBM energy density and resulting densification. The MMCs did
not exhibit cracking behavior in LBM, however, Ti segregation significantly reduced the Orowan
bowing mechanism and Hall-Petch grain strengthening. Gu et al. were able to distribute nano-
TiC onto spherical Ti powders through mechanical
mixing in a planetary mill without grinding media as
shown in Figure E (D. Gu, Wang, & Zhang, 2014). While
the same correlation between LBM energy density and
resulting densification was found, these composites
exhibited undesirable grain coarsening as linear energy
density increased from 250 J/m to 1000 J/m. Several
studies have reached the same grain coarsening
Figure E: Mechanically mixed Ti/nano-TiC
composites (D. Gu, Wang, & Zhang, 2014).
25
conclusion for higher energy densities in laser AM processing of AlSi10Mg- and Ni-based,
surface-coated composites (D. Gu, Wang, Chang, et al., 2014; Hong et al., 2015; Marchese et al.,
2018). However, these authors also concluded that increasing energy density positively affected
the wettability and the dispersion state of reinforcement particles, likely due to increased
Marangoni flow at higher energy densities. These results indicate a critical range of AM energy
densities for obtaining homogeneous MMCs with nanocrystalline microstructures.
Surface inoculation has recently been applied to MMC powder production for AM.
Lennart et al. ultrasonically dispersed TiC nanoparticles on gas-atomized Al7075 particles in an
electrolytic solution as shown in Figure F (Tasche et al., 2019). Subsequent drying etched the TiC
nanoparticles into the matrix oxide surface leading to a
homogeneous TiC dispersion. The authors reported preferable
epitaxial grain growth during LBM consolidation, but further
microstructural analysis is needed to confirm surface
inoculation’s ability to achieve homogeneous, nanocrystalline
TiC dispersion after AM processing.
Martin et al. recently explored a method of nanofunctionalization via electrostatic
assembly to distribute tungsten carbide (WC) nanoparticles on gas-
atomized AlSi10Mg powder as seen in Figure G (Martin et al., 2018).
LBM processing yielded a uniform reinforcement distribution
leading to an increase in strength and wear resistance compared to
unreinforced AlSi10Mg structures. Residual porosity was found to
exist within both functionalized and non-functionalized materials in
Figure G: Nanofunctionalized
AlSi10Mg/nano-WC composites
(Martin et al., 2018).
Figure F: Al7075/nano-TiC composite
particle prepared by surface inoculation
(Tasche et al., 2019).
26
LBM processing but was attributed to increased volumetric energy absorption with increased
reinforcement volume. Few details are given on the applicability and scalability of the
electrostatic assembly method.
Directly coating matrix powder surfaces with reinforcement particles offers the benefit of
including multiple phases while retaining the spherical particle morphology that is preferable for
near-complete densification in PBF. Several authors have concluded that higher applied energy
densities in AM lead to decreased porosity, likely due to increased energy absorptivity from
ceramic particles. However, reinforcement coarsening and phase segregation at higher energy
densities significantly reduce the Orowan effect and Hall-Petch grain strengthening in surface-
coated MMCs. These results indicate the importance of PBF parameters on the resultant
properties. Recent methods for chemically attaching reinforcement particles to matrix powder
surfaces also show preferable dispersion and grain refinement while retaining ideal flowability
and apparent density in MMC feedstock powders.
2.3.4 Mechanical Alloying
To combat the poor reinforcement homogeneity issue posed in other powder production
methods, several researchers have applied mechanical alloying to uniformly incorporate
reinforcement particles within metal matrix powders. Matrix and reinforcement powders are
placed with hardened metal balls in a sealed milling vial. The vial is then agitated in a linear,
cyclical, or planetary motion to repeatedly induce particle fracture and cold welding. Parameters
such as ball-to-powder ratio, milling speed, milling duration, and process control agents may be
used to tailor dislocation density, microstructure, and wettability (Suryanarayana, 2001).
27
Aluminum AM alloys, particularly AlSi10Mg, are a popular choice for mechanical
alloying as their ductility allows for homogeneous reinforcement encapsulation. Wang et al.
studied the effect of milling duration in producing 95.0 wt.% AlSi10Mg and 5.0 wt.% nano-TiC
composites via planetary milling (H. Wang & Gu, 2015). Both particle and grain sizes were refined
after 15 hours of milling. Additionally, the composite powders exhibited a uniform nano-TiC
dispersion both before and after LBM processing, leading to a 32% increase in hardness and a
20% increase in tensile strength with no loss in ductility as compared to unreinforced AlSi10Mg
components. The dynamics of the melt pool for a similarly milled AlSi10Mg/TiC composite in
LBM were observed by Yuan et al. (Yuan et al., 2015). It was observed that low (<500 J/m) linear
energy densities in LBM retained the desirable nanocrystalline reinforcement microstructure but
induced reinforcement particle agglomeration. Conversely, excessively high (1000 J/m) linear
energy densities in LBM retained uniform reinforcement distribution but induced undesirable
reinforcement grain coarsening, similar to the surface-coated MMC powders studied by Gu et al.
and Marchese et al. Increased Marangoni convection driven by higher energy densities
recirculates reinforcement particles to a uniform distribution in both ball-milled and surface-
coated MMCs. These results demonstrate the interconnectedness between the reinforcement
distribution in powder MMC feedstock, AM power and scanning speed, and the eventual
microstructure and mechanical properties of AM-produced components.
AM part porosity is also an important measure, especially with non-spherical powders
produced using mechanical alloying. Wang et al. fabricated AlSi10Mg and carbon nanotube
(CNT) composites via ball milling and LBM and studied the resultant densification and
microstructure (L. Wang, Chen, & Wang, 2017). A critical volumetric energy density of 131 J/mm3
28
was identified to minimize porosity and maximize mechanical properties. Lower energy densities
did not provide the wettability needed to enhance densification, while higher energy densities
proliferated porosity through the common “balling” effect, or spherical droplet formation, in the
melt pool. Higher energy densities also contributed to coarsening of the grain structure, resulting
in a decrease in hardness. For the critical energy density, fine, equiaxed grains contributed to a
15.7% increase in tensile strength compared to unreinforced AlSi10Mg despite an undesirable
1.2% decrease in ductility.
Han et al. utilized planetary milling to fabricate Al-Al2O3 composites specifically designed
for PBF (Han, Setchi, & Evans, 2017). Powders milled up to 20 hours experienced a transformation
from flat, irregular particles to rounded particles with preferable flow properties as seen in Figure
H. The authors noted the predominance of particle fusion early in the process as particle size
increased from 15 μm to 50 μm after 10 hours of milling. Further milling decreased the particle
size to 35 μm and created semi-spherical particle morphology. X-ray diffraction (XRD) revealed
a reduction in grain size from 71 nm after 8 hours of milling to 48 nm after 20 hours of milling.
Figure H: Morphological evolution of Al-Al2O3 composites after (a) 0, (b) 4, (c) 8, (d) 12, (e) 16, and (f) 20
hours of planetary milling (Han, Setchi, & Evans, 2017).
(a)
(b)
(c)
(d)
(e)
(f)
29
Subsequent work consolidated the 20-hour milled powder in SLM at various laser power
intensities and scanning speeds (Han, Setchi, Lacan, et al., 2017). Despite retaining relatively
porous surfaces, the as-milled feedstock powder achieved > 99% relative density at an energy
density of 315 J/mm3 and scanning speed of 300 mm/s. At this scanning speed, the consolidated
composite experienced further reduction in grain size attributed to the high cooling rates in LBM
and homogeneously dispersed Al2O3. Microhardness and tensile strength were increased by 17%
and 36% compared to unreinforced aluminum samples fabricated using LBM at the same
parameters. Grain refinement strengthening was attributed as the primary strengthening
mechanism. More importantly, these studies achieved flowable, PBF-feasible AMC powder using
mechanical alloying.
Like surface-coated MMCs, PBF parameters greatly affect the densification and
mechanical performance of mechanically alloyed MMC powder. However, dislocation density
and microstructure may also be tailored through mechanical alloying parameters. Non-spherical
powder morphology in mechanically alloyed MMCs is expected to significantly hinder
flowability, but current research indicates good flowability in mechanically alloyed composites,
despite achieving only semi-spherical particle shape. Future work should elucidate the
relationship between particle morphology and flowability in AM components.
Despite achieving excellent composite homogeneity and semi-spherical particle shape, all
of these studies utilized small amounts (2.0 to 5.0 wt.%) of a process control agent (PCA) during
the mechanical alloying process (Groza et al., 2007; Suryanarayana, 2001). This additive, typically
stearic acid, acts as a lubricant to prevent the subject material from excessively welding to the
milling equipment and is imperative when grinding ductile materials such as aluminum.
30
Nevertheless, PCAs are incorporated into the subject material and difficult to remove after
milling. When the fabricated powder is used in AM, the PCA degrades under high temperature
and risks forming a hydrocarbon defect in the consolidated material. These defects can lead to
premature failure of components, and methods of mechanically alloying ductile materials
without the use of PCAs should be investigated.
2.3.5 Cryomilling
Cryomilling, or mechanical alloying at cryogenic temperatures, is another MMC
production method that allows dispersoids to be fully encapsulated within matrix particles with
several benefits. Cryomilling achieves higher microstructural control when synthesizing Al-
based nanocomposites than traditional mechanical alloying (Suryanarayana & Al-Aqeeli, 2013).
Cryogenic temperatures also achieve the desired nanocrystalline microstructure more quickly
than room temperature mechanical alloying (D. B. Witkin & Lavernia, 2006). Similarly,
cryomilling creates high dislocation densities with minimal welding to milling equipment, thus
eliminating the necessity for process control agents as in room temperature alloying.
Homogeneously distributed in situ and ex situ aluminum nanocomposite powders have been
fabricated using this method (Suryanarayana & Al-Aqeeli, 2013).
Cryomilling promotes mechanochemical reactions which may be used to tailor MMC
powder composition. Typical cryomilling involves milling constituent powders and liquid
nitrogen slurry in a Szegvari attritor as shown in Figure I (Suryanarayana, 2001). Direct contact
31
between liquid nitrogen and aluminum during slurry
cryomilling forms in situ nanoscale aluminum nitride (AlN) and
aluminum oxynitride (Al(O,N)) phases which promote
dispersion strengthening via the Orowan mechanism. When
homogeneously distributed, these phases also act as nucleation
sites, effectively increasing the grain refinement strengthening.
Other ceramic particles such as oxides and carbides may also be
added in the milling process to encourage in situ reinforcement
synthesis in cryomilling (Suryanarayana & Al-Aqeeli, 2013). To avoid in situ oxidation and
nitridation, cryomilling may be performed in a manner where liquid nitrogen does not contact
the milling media. Kumar and Biswas successfully created a vibratory cryomilling vial with liquid
nitrogen cooling channels to fabricate high purity metallic nanoparticles (N. Kumar & Biswas,
2015).
Spex SamplePrep (Metuchen, NJ, USA) also produces a line of Freezer/Mill® that
submerge milling vials in liquid nitrogen, avoiding nitrogen and oxygen contamination. Contrary
to traditional studies which utilize a rotary attritor and spherical impactors to induce shear
fracture in a liquid nitrogen slurry, the Spex cryomill utilizes linear pulverization from a single
cylindrical impactor inside a milling vial
submerged in liquid nitrogen as depicted in
Figure J. The kinetics of this process are similar
to room temperature high energy ball milling in
the Spex 8000, a traditional ball mill for MMC
Figure I: Szegvari ball mill
(Suryanarayana, 2001).
Figure J: Side view of cryomilling vial, powder, and impactor.
32
production, but at cryogenic temperatures (Suryanarayana, 2001). Compared to attritor
cryomilling, the Spex cryomill avoids nitradation, the development of AlN dispersoids, since the
milling material does not contact liquid nitrogen. Key parameters for this process are impactor-
to-powder mass ratio, milling time, process control agents, and grinding impactor geometry. This
type of mill has been used to fabricate an Al-CuO thermite composite but remains unexplored for
PBF feedstock production (Badiola, Schoenitz, Zhu, & Dreizin, 2009).
Consolidation of the cryomilled AMC powders is also a pertinent topic as the processing
parameters directly affect the microstructure and resultant properties. Hot isostatic pressing
(HIP) is one common method for achieving full densification. Hayes et al. noticed full
densification and the formation of Al3Ti phases and Al2O3 and Al4C3 dispersoids that significantly
contributed to the part strength in cryomilled Al-10Ti-2Cu powder fabricated via HIP (Hayes,
Rodriguez, & Lavernia, 2001). However, an inconsistent microstructure featuring 30 nm to 70 nm
and 300 nm to 500 nm grains was found to degrade tensile strength at higher temperatures. The
as-pressed composite also suffered from low ductility attributed to low porosity, interstitial
reinforcement dispersion, and a lack of work hardening. Witkin et al. observed the same grain
coarsening phenomenon for cryomilled Al-Mg alloys consolidated using HIP (D. Witkin, Han, &
Lavernia, 2006). While higher processing temperatures yielded more complete densification,
significant grain growth occurred at these temperatures which degraded tensile performance
compared to extruded composites. These results indicate a tradeoff between densification and
microstructural homogeneity.
To the best of the author’s knowledge, cryomilled AMC powders have yet to be
consolidated via AM. With various cryomilling parameters, particle morphology, dislocation
33
density, and phase homogeneity may be tailored to fabricate ideal PBF feedstock powder. Several
authors have utilized cryomilling to fabricate nanostructured aluminum matrix composites, but
their customizable potential in AM remains unexplored. To make AMC powders feasible,
however, the specific powder morphology and size distribution required for PBF must be met.
Previous work has successfully achieved these properties in AMCs through mechanical alloying
(Han, Setchi, & Evans, 2017; H. Wang & Gu, 2015). Further work needs to be performed to identify
the ideal cryomilling parameters for designing PBF-compatible, nanocrystalline, AMC powders.
2.4 Summary and Critique of the Chapter
Metal matrix composites are becoming an increasingly important material in PBF.
Elevated mechanical performance due to grain refinement and dislocation strengthening
designate aluminum matrix composites and nanocomposites as next-generation AM materials.
However, the optimal method for producing AMC powder intended for PBF remains under
investigation. Near-spherical particle shape and a critical size distribution should be achieved in
PBF feedstock powder, and AM-produced components should retain a homogeneous
microstructure for maximizing strength and discouraging hot cracking. Producing
homogeneous, lightweight AMCs using additive manufacturing requires careful parametric
control. PBF parameters such as power intensity and scanning speed affecting the thermal
gradient and solidification rate greatly influence microstructure, phase homogeneity, porosity,
and resultant mechanical properties.
Mechanically alloyed AMCs demonstrate similar solidification issues as surface-coated
AMCs when subjected to PBF processing. Higher energy densities result in grain coarsening and
34
porosity and a subsequent weakening of mechanical properties, despite retaining a homogeneous
reinforcement distribution. Conversely, lower energy densities retain the desired nanocrystalline
microstructure, but suffer from reinforcement particle agglomeration and porosity. Therefore, a
critical energy density, tailored for AMC absorptivity, is needed to minimize porosity and grain
size, retain uniform reinforcement distribution, and maximize the AMC strengthening
mechanisms.
It is interesting to note that, in these studies, particle morphology was not presented as a
limiting factor for mechanically alloyed powders in PBF. Powder flowability is a critical factor for
achieving full density components in PBF. Mechanically alloyed and cryomilled powders are
known to have irregularly shaped particles that are expected to hinder flowability and increase
porosity. However, to the extent of the author’s knowledge, no study has examined flowability
and PBF part densification as a function of the mode of MMC powder production.
Similarly, each MMC powder production method achieves various dislocation densities
and dispersion states of the reinforcement phase. For example, surface coating methods provide
a reinforcement distribution that relies on Marangoni convection during AM to achieve
uniformity, whereas mechanical alloying and cryomilling produce uniform reinforcement
distribution and high dislocation densities prior to AM. The author is not aware of any study that
quantifiably compares various MMC powder production methods to the resultant reinforcement
distribution, microstructure, and mechanical properties of AM-produced components. The
degree of dislocation strengthening and grain refinement strengthening depends on the size and
dispersion state of the reinforcement phase after AM, and future studies should examine the post-
AM microstructural and mechanical effects of various MMC powder feedstock production
35
methods. Similarly, the author is not aware of any study that utilizes cryomilling as a successful
mode of AMC powder production for PBF feedstock powder. Higher microstructural control and
uniform reinforcement distribution obtained via cryomilling is expected to allow tailored AMC
feedstock powder for PBF-AM.
36
3. Preliminary Work
Additive manufacturing (AM) has proven to be an effective method for fabricating
complex metal objects directly from 3D models, but few lightweight metal materials are currently
available for powder bed fusion (PBF). Aluminum matrix composites (AMCs) are one lightweight
alternative, but their customizable potential in PBF has yet to be explored. Developing critical
parameters for AMC powder production and PBF processing would allow for new, lightweight
materials to be customized for biomedicine, aerospace, and other applications. Before beginning
the investigation of the primary hypothesis, a pilot study was conducted to define AMC powder
production and characterization methods.
3.1 Cryomilling Al-TiC Study
Work presented at the TMS 2019 Annual Meeting and Exhibition
Jakob D. Hamilton1, Mouda Tung2, Ola L. A. Harrysson2, Shalabh Gupta3, Iris V. Rivero1, Christopher D. Rock2
To investigate the feasibility of cryomilling as a method of AMC production, an initial
study was conducted at Ames Laboratory (Ames, IA, USA) in May 2018. The primary objectives
of this work were to develop cryomilling parameters to fabricate AMC powder and assess the
feasibility of as-milled composite powder for PBF AM.
3.1.1 Composite Materials
Pure (99.5%) aluminum (Al) with maximum particle size of 125μm from Goodfellow
(Corapolis, PA, USA) was used for the matrix material. Titanium carbide (TiC) with maximum
1 Department of Industrial and Systems Engineering, Rochester Institute of Technology, Rochester, NY,
USA 2 Edward P. Fitts Department of Industrial and Systems Engineering, North Carolina State University,
Raleigh, NC, USA 3 Division of Materials Science & Engineering, Ames Laboratory, Ames, IA, USA
37
particle size of 200nm from Sigma Aldrich (St. Louis, MO,
USA) was used as the reinforcement material due to its high
abrasiveness and strength properties. Images of as-received
Al and TiC powders are shown in Figure K. As-received Al
exhibited non-spherical particle shape possessing poor
flowability. TiC powder formed micron-sized agglomerates
characteristic of nanoscale particles. Nanoscale reinforcement
particles were selected to mitigate crack-inducing stresses
within the matrix as demonstrated in Martin et al. (Martin et
al., 2017). All materials were handled in an inert (2ppm oxygen) glovebox to inhibit oxidation and
reduce flammability.
3.1.2 Cryomilling
Linear cryomilling using a Spex SamplePrep 6875D Freezer/Mill® was selected to
appropriately incorporate nanoscale TiC particles within the Al matrix as well as to avoid
nitrogen contamination. An image of this mill is shown in Figure L. Previous researchers in the
Interdisciplinary Manufacturing Engineering and Design Laboratory (iMED) have tasted success
in homogenizing polymer composites using this style of cryomill (Allaf & Rivero, 2011; Allaf,
Rivero, Abidi, & Ivanov, 2013; Ramesh, 2017). However, to the best of the author’s knowledge,
this style of cryomill has yet to be employed in the fabrication of AMC powder designed for AM.
A mixture of 95.2 wt.% Al and 4.8 wt.% TiC was cryomilled for 2, 4, and 6 hours. The total
powder charge in the milling vial was 5.25 grams, and the impactor-to-powder mass ratio (IPR)
was 28:1. To reduce contamination in as-milled powder, process control agents were excluded
Figure K: as-received (a) Al and (b) TiC
powder images taken using SEM.
(a)
(b)
(a)
38
from the milling environment. Vials were preliminarily cooled for 20 minutes in liquid nitrogen
at 77 Kelvin to encourage particle fracture. The cryomilling was performed in intervals of 10
minutes of cryomilling at 10 cycles per second followed by 1 minute of cool down. Small amounts
of powder were removed after every 12 intervals to characterize as-milled powder.
3.1.3 Powder Characterization
Resultant powders were characterized for particle morphology, particle size distribution,
and surface-level reinforcement distribution through scanning electron microscopy (SEM).
Reinforcement distribution, quantified by 𝑑 indices, was compared to milling time. The 𝑑 index
is defined as
𝑑 =𝑠
�̅�
where �̅� and 𝑠 are, respectively, the mean and standard deviation of the set of nearest neighbor
distances between reinforcement particles in a sample image (Chawla & Chawla, 2013). Lower 𝑑
indices indicate uniform dispersion.
Particle morphology and size distribution of as-milled, unsieved powders were
compared to as-received AlSi10Mg feedstock powder from LPW-Carpenter Additive (Imperial,
(a)
(b)
(c)
Figure L: Images of the (a) outside, (b) inside, and (c) milling vial of the SPEX 6875D Freezer/Mill®.
39
PA, USA). The gas-atomized LPW powder exhibited near-spherical morphology and a particle
size distribution observed to be D10 = 18.6μm, D50 = 33.5 μm, and D90 = 58.2 μm. These values were
set as the ideal size for LBM.
Cryomilled composite powders exhibited significant changes in particle morphology
throughout the milling process as seen in Figure M. The 2-hour cryomilled samples exhibited a
significant number of large (>200 µm), flat chips, indicating the predominance of particle fusion
early in the cryomilling process. Longer milling times of 4 and 6 hours transformed particles to
more spherical shapes with rough surfaces indicative of particle fracture. The 6-hour cryomilled
powder exhibited many small (< 10 µm), irregularly shaped particles in addition to medium (~50
µm) particles similar to the 4-hour powders. As milling time increased, particles underwent work
hardening making them more prone to fracture, especially with the inclusion of hard
reinforcement particles and cryogenic temperatures.
The particle size distribution evolution was quantified using ImageJ particle analysis.
Particle areas were measured, and equivalent diameters were calculated assuming circularity.
Figure M: SEM images of cryomilled AMC powder after: (a) and (d) 2 hours, (b) and (e) 4 hours, (c) and (f) 6 hours.
(a)
(b)
(a)
(c)
(a)
(d)
(a)
(e)
(a)
(f)
(a)
40
Particle size distributions at
each milling time and a
comparison of mean particle
sizes between milling times
can be seen in Figure N and
Figure O, respectively. The
work hardening phenomenon
can be observed as particle
size drastically decreases for
longer cryomilling durations.
Mean particle size shifted smaller with longer milling increments. By 6 hours, the mean particle
size shifted below the AlSi10Mg control sample. Standard deviation in particle size also decreased
throughout the milling process indicating the homogenization occurring due to steady-state
particle fracture. Both the mean particle size and breadth of the size distribution are clearly
tailorable with milling time. These
results indicate the ability to produce
AMC powder characteristically
similar to commercially available
aluminum AM feedstock powder
through cryomilling. Figure O: Mean particle size comparison between all cryomilled Al-TiC
samples and LPW AlSi10Mg powder.
Figure N: Particle size distributions for (a) 2-hour cryomilled Al-TiC, (b) 4-
hour cryomilled Al-TiC, (c) 6-hour cryomilled Al-TiC, and (d) LPW.
AlSi10Mg powders
41
Reinforcement distribution was determined through image contrast thresholding via
ImageJ, and an example image is shown in Figure P. Comparison of reinforcement distribution
as a function of milling time is shown in Figure Q. The results indicated that there is little, if any,
difference in surface-level reinforcement distribution between milling times.
There are two important caveats to the measurement techniques used to measure
reinforcement distribution in this study. The first being the type of images taken. In backscattered
electron detection SEM, regions with higher atomic numbers backscatter electrons more strongly.
The effect is that these regions appear as a much lighter contrast, as seen with Ti in the Al-TiC
composites. However, particle topography also is communicated through image contrast, so
topographical noise may influence the distribution measurement. Instead, energy-dispersive X-
ray spectroscopy (EDS) should be employed to identify individual elements by the characteristic
X-rays emitted during measurement. The second caveat is that all images were of particle
surfaces. Because cryomilling imparts severe plastic deformation on particles, reinforcement
particles are expected to be found within matrix particles. A more relevant measure of
Figure P: Example of the original and threshold image for
reinforcement distribution measurement.
642
1.6
1.5
1.4
1.3
1.2
1.1
1.0
0.9
Milling Time
D I
nd
ex
Interval Plot of D Index95% CI for the Mean
Individual standard deviations are used to calculate the intervals.
Figure Q: Interval plot of d indices of surface-level
reinforcement distribution for individual milling times.
42
reinforcement distribution should be gathered by sectioning AMC powder to view internal
features.
The use of d indices for composite homogeneity measurement may also be a point of
improvement. d indices offer the flexibility of quantifying the reinforcement dispersion
normalized using the mean particle spacing. However, d indices overcompensate for outliers in
a uniform distribution of particles. This measurement is also reliant on the reference frame.
Clusters of particles may exhibit the same d index as uniformly distributed powder, as the nearest
neighbor would be similar for particles in clusters. Ripley statistics are a more rigorous measure
of composite homogeneity as they measure the particle location probability as a function of
distance from the origin. However, this measurement does not naturally consider edge-effects in
particle images and requires fine-tuning to use appropriately.
3.1.4 Conclusions
The results from this study indicate linear cryomilling may be used to fabricate an Al-TiC
composite powder with minimal contamination from process control agents or aluminum
nitrides. Surface-level observations indicate TiC nanoparticles were incorporated successfully
onto the perimeter of Al matrix particles. Moving forward with developing AMC powder on a
more practical scale for custom PBF feedstock will require a more comprehensive understanding
of key process parameters for powder fabrication such as reinforcement loading, cryomilling
time, and powder mass per vial. Based on previous literature, these factors are all expected to
impact the resultant powder characteristics (Suryanarayana, 2001; Zhou, Nutt, Bampton, &
Lavernia, 2003). Therefore, another study should be crafted to glean a fuller understanding of
43
cryomilling process parameters. The subsequent chapter is formatted as a separate journal article
for publication.
44
4. Solid-state Synthesis of Aluminum Matrix Composites for
Powder Bed Fusion Additive Manufacturing
Jakob D. Hamiltona, Srikanthan Ramesha, Ola L.A. Harryssonb, Christopher D. Rockb, Iris V.
Riveroa
aDepartment of Industrial and Systems Engineering, Rochester Institute of Technology,
Rochester, New York 14623, USA.
bEdward P. Fitts Department of Industrial and Systems Engineering, North Carolina State
University, Raleigh, North Carolina 27607, USA.
Abstract
The spectrum of aluminum-based materials available for additive manufacturing (AM) remains
limited owing to complex phase transition characteristics and poor mechanical performance. To
overcome these problems, fillers are often added to aluminum matrices to create a class of
materials called AMCs that combine the ductility of aluminum with the stiffness of ceramic or
carbon-based reinforcements. Traditional powder production methods through liquid-phase
atomization are not suited to produce composite powders as particle segregation discourages
composite homogeneity. AM powder production through solid-state mechanical alloying has
been studied with limited success, primarily due poor powder spreadability and inclusion of
lubricants in the alloying process. Cryogenic mechanical alloying, termed cryomilling, enhances
homogeneity between matrix and reinforcement particles by recurrent fracture and cold welding
sans lubricants. In this study, aluminum matrix composites tailored for use in powder bed fusion
(PBF) AM were produced via cryomilling at varying compositions, powder charges, and milling
times. As-milled powders were characterized for particle size distribution, morphology, and
homogeneity. Resultant powder demonstrated varying characteristics correlated to milling
parameters. A representative spreading test was performed to replicate spreading performance
in PBF. Powder metallurgy samples were also fabricated to understand as-sintered reinforcement
distribution and the resultant strengthening. This research provides an indication of cryomilling
capabilities to become an effective method for custom alloy powder production for PBF-AM.
Keywords: aluminum matrix composites (AMCs), cryomilling, additive manufacturing
4.1 Introduction
Metal additive manufacturing (AM) is an attractive industrial manufacturing technology,
as complex, near-net geometries may be fabricated without fixturing from difficult to process
alloys such as titanium and nickel alloys (Herzog et al., 2016). Powder bed fusion (PBF) is a subset
45
of AM methods that utilizes an energy beam to selectively melt a thin layer of powder. High
solidification rates in PBF promote mechanically-superior parts, however, the complex melting
and solidification dynamics has greatly limited the number of materials currently available for
this process (Herzog et al., 2016; Manfredi et al., 2014). For aluminum alloys, excessive reflectivity
at laser wavelengths, elemental vaporization, and solidification cracking limit the available alloys
to casting alloy AlSi10Mg and Al-Si-Mg derivatives (Herzog et al., 2016; Manfredi et al., 2014;
Martin et al., 2017). Particulate-reinforced aluminum matrix composites (AMCs) are one
alternative that have been utilized previously in aerospace and automotive industries through
formative manufacturing processes and show potential for suppressing excessive thermal
contraction in PBF (Chawla & Chawla, 2013; Manfredi et al., 2014). AMCs consist of a bulk metal
matrix and a dispersed reinforcement phase that encourages crystallite nucleation in molten
metal (Herzog et al., 2016). While components featuring micron-sized reinforcement materials
have been fabricated successfully through PBF, the creation of parts with high strength has been
a problem due to the loss of ductility in as-fabricated components (Tjong, 2007). Several authors
have employed nanoparticle-reinforced aluminum composites in laser powder bed fusion that
exhibit superior specific strength and ductility to unreinforced and micron-reinforced aluminum
materials (D. Gu et al., 2015; H. Wang & Gu, 2015).
Fabricating nanocomposite feedstock powder is challenging as large Van der Waals forces
between nanoparticles drive agglomeration and result in weakened mechanical performance of
as-built components (Cabeza et al., 2017; H. Wang & Gu, 2015). While traditional powder
production through atomization is capable of generating spherical powder with excellent
spreading properties, dispersion of secondary phase particles in the liquid state is a significant
46
challenge due to large Van der Waals forces between sub-micron particles (Chawla & Chawla,
2013). Several authors have employed severe plastic deformation through mechanical alloying to
homogeneously incorporate reinforcement materials in PBF feedstock. Wang et al. produced
AlSi10Mg and nano-TiC composites via planetary milling for 15 hours (H. Wang & Gu, 2015).
The composite powders exhibited a uniform nano-TiC dispersion both before and after laser
beam melting, leading to a 32% increase in hardness and a 20% increase in tensile strength with
no loss in ductility as compared to unreinforced AlSi10Mg components. Han et al. also utilized
planetary milling to fabricate Al-Al2O3 composites specifically designed for PBF (Han, Setchi, &
Evans, 2017). Powders milled up to 20 hours experienced a transformation from flat, irregular
particles to rounded particles with preferable flow properties. Subsequent work consolidated the
20-hour milled powder in laser beam melting and observed a 17% and 36% increase in
microhardness and tensile strength compared to unreinforced Al for the same parameters (Han,
Setchi, Lacan, et al., 2017). These studies achieved flowable, PBF-feasible AMC powder using
mechanical alloying. Nevertheless, stearic acid, a common process control agent, was utilized to
prevent welding to the milling equipment in both studies. Ductile materials such as aluminum
are especially prone to cold welding so dry grinding without process control agents lengthens
the required milling time and can damage milling equipment (Suryanarayana, 2001). These
solvents catalyze hydrocarbon defects in AM-consolidated materials which degrades mechanical
performance of AM parts, the driving purpose for utilizing AMCs.
One method of reducing or omitting process control agents in AMC powder is mechanical
alloying at cryogenic temperatures or cryomilling (Groza et al., 2007; Suryanarayana, 2001).
Cryogenic temperatures inhibit undesirable material reactions during milling and prevent
47
excessive aluminum welding to the milling material. Additionally, rapid formation of
nanocrystalline grain structures through cryomilling favors particle fracture, decreasing the
required milling time for composite homogeneity (Suryanarayana, 2001). Therefore, we
hypothesized that the use of cryomilling can eliminate contamination in the fabrication of AMC
powder for PBF as well as promote desirable spreading properties in composite powder.
Herein, we present a systematic investigation of cryogenic milling as a fabrication method
for generation of AMC powders for PBF. Experiments were designed to understand the role of
process parameters such as milling time, powder mass, and reinforcement loading in influencing
particle fracture, composite homogeneity, and resultant morphology.
To the best of our knowledge, this is the first study to look at linear cryogenic milling as
an option to fabricate AMC powder for PBF. Linear cryomilling was used to fabricate an Al-CuO
thermite composite but has yet to be explored for PBF feedstock fabrication (Badiola et al., 2009).
Milling time, milling speed, and ball-to-powder ratio are all factors that can control the shape and
reinforcement distribution of fabricated powder (Suryanarayana, 2001). Attritor cryomilling has
only recently been explored to fabricate AMC feedstock powder for PBF (Kellogg, Kudzal,
Rogers, & McWilliams, 2019).
In order to expand the scope of lightweight composite materials for PBF, methods of
homogeneously dispersing reinforcement materials in PBF feedstock are required. Linear
cryomilling may be an appropriate method for fabricating custom, homogeneous nanocomposite
powder with adequate spreading properties. Therefore, the present works aims to understand
morphological evolution corresponding to cryomilling process parameters and successfully
fabricate AMC powder characteristically similar to PBF feedstock materials. The present work
48
utilizes a single-stage cryomilling process and ex situ characterization methods to address these
goals.
4.2 Materials and Methods
Commercially available, gas atomized AlSi10Mg powder (D10 = 18.2 μm, D50 = 33.0 μm, D90
= 56.7 μm) from LPW-Carpenter Additive (Imperial, PA, USA) was selected for the matrix
material. TiC powder from Sigma Aldrich (St. Louis, MO, USA) with maximum particle size of
200 nm was used as the reinforcement material due to its high specific strength. SEM images of
as-received AlSi10Mg and TiC powder can be seen in Figure R. All powder was handled within
an inert (<1000 ppm oxygen) glovebox to avoid severe oxidation and reduce flammability.
Figure R: As-received (a) AlSi10Mg powder from LPW-Carpenter Additive and (b) nanoscale TiC powder from Sigma-Aldrich.
A SPEX SamplePrep 6875D Freezer/Mill® was employed to cryomill the AMC powder.
Figure S depicts the selected cryomilling method. This mill submerges a sealed, powder-filled
milling vial in liquid nitrogen. A cylindrical impactor within the vial is driven by an oscillating
magnetic field. The small vials from SPEX SamplePrep used in this study have a 19.2 mm internal
diameter and 79.0 mm maximum length of travel between stainless steel end caps. The stainless
steel impactor paired with this vial features 9.5 mm diameter, 60.3 mm length, and mass of 32.6
49
g. Two impactor-to-powder mass ratios (IPR) were tested, and two levels were chosen for TiC
loading. The distinct levels for the full factorial design were based in literature and previous
experimentation and can be seen in Table 3 (Suryanarayana & Al-Aqeeli, 2013). Milling speed
was set at 10 cycles per second, and a 20 minute precool preceded each 1 hour milling segment.
A small amount (< 20 mg) of AMC powder was withdrawn at one hour intervals to study
morphological evolution and reinforcement distribution as a function of milling time. In total,
each sample was cryomilled for 6 hours. Oversize particles were removed from the bulk 3-hour
and 6-hour cryomilled samples using a 200 mesh (86 μm opening) sieve screen before spreading
and strength characterization.
Figure S: Depiction of the cryomilling vial, impactor, and direction of impactor travel. The vial is submerged in liquid nitrogen
during the milling process.
Table 3: Design of experiment for the AlSi10Mg-TiC cryomilling study. “L” and “H” represent low and high powder mass,
respectively.
Sample Identifier
IPR Powder Mass (g)
TiC Charge (wt.%)
L-1TiC 16:1 2.04 1.0
L-10TiC 16:1 2.04 10.0
H-1TiC 4:1 8.20 1.0
H-10TiC 4:1 8.20 10.0
X-ray powder diffraction (XRD) was performed on as-received and as-milled materials to
view the composition and crystalline evolution before and after the cryomilled samples. A Rigaku
50
Geigerflex XRD scanned materials using Cu-Kα x-rays (λ = 1.54056 Å) with 2θ from 35° to 140°.
Particle morphology was studied using a JEOL JSM-IT100 scanning electron microscope (SEM)
(JEOL Ltd., Akishima, Japan). The working distance was between 15-22 mm, accelerating voltage
was 16 kV, and probe current was between 45-75%. Particle size distributions were determined
through thresholding of SEM images of unmounted powder in ImageJ software. Ellipses were fit
on thresholded particles in the ImageJ, and the particle diameter was calculated assuming
circularity. Energy dispersive x-ray spectroscopy (EDS) was used to study the reinforcement
distribution in images of epoxy-mounted and cross-sectioned powder. EDS images were taken
using the same SEM with an accelerating voltage of 16 kV, probe current of 75%, and working
distance of 10mm in low-vacuum mode.
To replicate the PBF recoating process, a doctor blade with a slot size of 16.75mm width
and 0.07mm depth was assembled. This method was intended to be an assessment tool to gauge
the spreading pattern of custom powder for a typical PBF layer height. Powder for this test was
used after passing it through the sieve. Powder samples were evenly dragged across the block
surface and into the slot using a razor tilted at a 45 angle°. Powder metallurgy was used to find
the relative strength between as-received and sieved, cryomilled materials. One gram of sieved
powder was added to a 12.7 mm diameter cylindrical mold and compressed to 152 MPa using a
Carver 3851 hydraulic press. Green compacts were sintered in an Ar environment at 577 °C for
one hour. Sintered samples were mounted and cross sectioned to evaluate their internal features.
Vickers Microhardness testing was performed on sintered samples using a 0.025 kg load and 5
second dwell. Five hardness measurements were taken from random locations on each sample’s
polished surface.
51
4.3 Results
4.3.1 X-ray diffraction (XRD) analysis
The diffractograms of as-received AlSi10Mg powder, L-10TiC, and H-1TiC after 3 hours
of cryomilling are shown in Figure T. Distinct, sharp peaks corresponding to aluminum planes
were visible for all cryomilled samples with no deviation from the unmilled powder used as
reference. Peaks did not appear to shrink or elongate, indicating no significant phase
transformation occurred during the cryomilling process. Peaks for the (220) and (311) planes for
silicon were also visible due to the 9.0-11.0 wt.% silicon content in the LPW powder. Magnesium
peaks were not visible as the magnesium content is less than 0.5 wt.%, sufficiently small to be
detected using XRD. Peaks for TiC were apparent in L-10TiC, however no TiC peaks were visible
in the XRD spectrum for H-1TiC due to the 1 wt.% TiC loading. Neither iron nor chromium was
observed in the XRD patterns, a sign that contamination from the milling vial and impactor was
minimized.
52
Figure T: XRD diffractograms for as-received AlSi10Mg powder and L-10TiC and H-1TiC after 3 hours of cryomilling.
4.3.2 Morphological Characterization
SEM images of cryomilled samples throughout the milling process are shown in Figure
U. The 4:1 IPR samples retained predominantly spherical particles into the second hour of milling.
After 3 hours, several large (> 100 μm) semi-spherical welded particles remained, however
particles continued to decrease in size. Small (< 20 μm) flakes formed immediately in both high
powder mass samples were most prevalent after 6 hours. These flakes were observed earlier in
the milling process for H-10TiC than H-1TiC indicating a positive correlation between
reinforcement loading and fracture rate. By the 6 hour milling mark, most of the coarse particles
became rounded and appropriately sized for PBF, i.e. 15 μm to 75μm.
53
Figure U: SEM images of all samples throughout cryomilling. Red circles denote agglomerated fine particles.
The 16:1 IPR samples lost the initial size distribution and shape and became large
agglomerates immediately during the first hour of milling. As cryomilling continued, particle
fracture dominated the milling process and formed predominantly jagged flakes. By the 6 hour
milling mark, a majority of particles were less than 5 μm and agglomerated into porous structures
of fine particles evident in Figure V. The size distribution of these powders indicates a higher
level of particle refinement was achieved than was necessary for PBF feedstock powder.
54
Figure V: Porous agglomerates comprising L-1TiC after 6 hours of cryomilling.
Table 4 quantifies the particle size distributions for each treatment combination as
determined by image analysis. Distinct powder characteristics due to independent factors after 3
hours of cryomilling are shown in Figure W. In this study, particle size appeared to vary for all
three factors: powder charge, reinforcement loading, and milling time. Milling time, TiC loading,
and impactor to powder ratio appear inversely correlated with resulting particle size distribution.
These results are consistent with previous findings for mechanical alloying (Suryanarayana,
2001).
Table 4: Particle size distributions of the cryomilled samples at all milling times.
Sample Identifier
Cryomill Time (h)
D10 (μm) D50 (μm) D90 (μm)
L-1TiC 1 1.05 4.91 47.27 2 1.34 6.09 38.22
3 1.12 4.71 23.09
6 1.13 4.71 27.99
L-10TiC 1 1.36 9.84 39.52 2 3.32 15.41 53.07 3 2.83 9.33 31.02 6 1.13 4.30 23.05
H-1TiC 1 7.32 28.21 54.73 2 4.28 23.51 58.04 3 2.18 10.49 46.74 6 1.24 5.75 23.53
55
H-10TiC 1 3.14 12.92 46.66 2 2.78 10.07 44.05 3 1.43 5.04 23.52 6 1.25 5.57 33.02
Figure W: Comparison between particle size and aspect ratio for (a) cryomilling time, (b) IPR, and (c) TiC loading.
Aspect ratio defined in image analysis was used to describe particle morphology; aspect
ratios closer to 1 indicate higher sphericity. The effects of independent factors on the aspect ratio
of powder particles are also shown in Figure W. The standard deviations are high as all powder
from all parameters is considered in the calculation of mean and standard deviation. Particle
images and previous literature demonstrate that particles lose sphericity during initial milling.
Continued milling increases the rate of steady state fracture and slightly homogenizes particle
morphology. Similar to particle size distribution, increased TiC loading and powder charge were
found to decrease the aspect ratio of as-milled powder.
4.3.3 Metallographic Analysis
Cross sections of mounted, ground, and polished powder were observed in SEM to view
intra-particle porosity. Full density was achieved in the majority of the observed powder particles
in all samples. For particles severely welded, a thin gap was observed between particles where
incomplete adhesion formed a thin pocket. This was primarily observed in the 16:1 IPR samples
56
where flake-like particles were much more predominant. The spherical and semi-spherical
particles in the 4:1 samples showed complete densification. The contrast between L-10TiC and H-
10TiC is shown in Figure X.
Figure X: Powder cross sections of (a) L-10TiC and (b) H-10TiC after 3 hours of cryomilling. The red rectangles outline
incompletely welded particles.
Examining internal TiC distribution connected milling time and reinforcement
incorporation. EDS maps for the primary and secondary phase in L-10TiC and H-10TiC
throughout the cryomilling process are provided in Figure Y. Outlines of Ti are clearly seen on
the surface of aluminum particles after the first hour of milling. As deformation persists,
previously surface-level TiC becomes a line of reinforcement material within welded particles.
Figure Z shows a clear example of Ti striations in an agglomerated particle. Further deformation
of welded particles yielded uniform TiC dispersion. The 4:1 IPR samples maintained surface-level
TiC dispersion throughout the milling process up to 6 hours. Low particle welding in these
samples did not promote the same level of internal TiC dispersion as the 16:1 IPR samples,
however surface-level TiC dispersion appeared uniform.
57
Figure Y: EDS images of cryomilled AMC powder displaying TiC distribution for distinct treatment combinations.
Figure Z: SEM and elemental mapping of a large particle agglomerate from L-10TiC after 3 hours of cryomilling illustrating the
primary and second phase striations.
58
4.3.4 Powder Spreadability
To fully understand the connection between particle shape, size distribution, and the
resultant spreading properties for PBF, a representative spreading test was performed on sieved,
3-hour and 6-hour cryomilled samples and the as-received AlSi10Mg powder. Select powder
spread patterns are shown in Figure 11.
Figure AA: Spread patterns of (a) L-1TiC-3h, (b) L-1TiC-6h, (c) H-10TiC-3h, (d) H-10TiC-6h, and (e) as-received AlSi10Mg
powder.
The spherical AlSi10Mg powder produced an even layer of powder when spread across
the apparatus. The cryomilled powder from H-1TiC and H-10TiC produced similar albeit less
dense patterns than the AlSi10Mg powder. Spread patterns from low powder mass samples, i.e.
L-1TiC and L-10TiC, were much sparser than the AlSi10Mg control and the other cryomill
samples due to their excessive flake-like morphology. High powder mass produced similar
spreading patterns at each cryomill time. Linear voids parallel to the drag direction appear to be
more prevalent in H-1TiC and H-10TiC after 3 hours of cryomilling. Oversize particles will not
flow beneath the razor blade and are the primary contributors to these streaks. These streaks
were not apparent in the 6 hour cryomilled samples, as the particle size distribution shifted
59
downward with longer milling time. Both 6 hour cryomilled samples produced qualitative
spread patterns characteristically similar to the as-received, gas atomized AlSi10Mg powder.
There was no apparent difference due to TiC loading with the exception of the color of the
powder.
4.3.5 Microhardness
Vickers microhardness measurements of compressed and sintered samples were used to
draw preliminary strength results for the TiC loading and IPR factors. Figure BB quantifies
microhardness for the 3-hour milled samples. The cryomilled samples exhibited improved
microhardness than the as-received AlSi10Mg samples. The 1% TiC samples exhibited a
minimum of 75% improvement in microhardness, while the minimum improvement for the 10
wt.% samples was 150%. Cryomilled samples exhibited much higher standard deviation than the
AlSi10Mg samples, the largest being L-10TiC-3h which exhibited standard deviation six times as
high as AlSi10Mg.
Figure BB: Vickers microhardness measurements for the 3-hour cryomilled samples. Error bars denote standard deviation.
60
4.4 Discussion
The milling parameters in this study yielded vastly different particle shapes and sizes as
they have been shown to affect the rate of fracture and welding of particles (Groza et al., 2007;
Suryanarayana, 2001). In planetary milling, particles undergo a significant morphological
transformation from spherical to flake-like and back to semi-spherical, rounded particles.
Initially, particles with low dislocation densities deform under the stress of milling to form flakes.
With repeated stress, flakes continue to deform and weld with other flakes to become
agglomerated particles. As dislocations build up in deformed particles through work hardening,
particles become brittle, and fracture becomes the dominating mechanism of material
deformation. This state of uniform particle refinement is referred to steady-state fracture and is
preferred for refining and homogenizing AMC powder. Both Wang et al. and Han et al. achieved
this stage for milled AMC powder designed for AM (Han, Setchi, & Evans, 2017; H. Wang & Gu,
2015).
Slightly different deformation mechanics were observed using linear cryomilling and
appear to depend on the impactor-to-powder mass ratio. Most particles in the 16:1 IPR samples
appeared as flakes with jagged surface texture throughout the duration of milling. Initial milling
of L-1TiC and L-10TiC produced excessively large (> 500 µm) particles which did not appear in
subsequent hours of cryomilling. This indicates the particle welding phase of mechanical alloying
transitioned to steady state fracture between the first and second hour of milling. Heavily stressed
particles, apparent by the extreme loss of sphericity in the first hour of milling, experience
continued work hardening and quickly transition to a stage of brittle particle fracture. Particle
refinement continued and eventually reached a predominance of fine, angular particles after 6
61
hours of milling. Excessive pulverization using this powder mass encouraged particle welding
early in the process and excessive particle fracture in the later stages.
Contrarily, 4:1 IPR samples did not demonstrate a welding-dominated milling phase.
Instead, particles maintained roundness and transitioned between minor particle welding to
steady state fracture slowly after the first hour, eventually reaching a size distribution featuring
semi-spherical coarse particles and fine flakes. Still, the 4:1 samples did not reach the same level
of TiC distribution homogeneity, particularly in coarser particles, as the 16:1 samples. In the 4:1
IPR samples, the applied energy from the impactor was dispersed amongst a heavier powder
charge. Likewise, increased powder volume in these samples also limits the length of impactor
travel, further reducing the applied energy. Because of this, H-1TiC and H-10TiC avoided
extreme particle deformation and retained adequate particle roundness for spreading. The
apparent lack of particle welding in the 4:1 IPR samples promoted surface-level TiC dispersion.
On the other hand, the severe particle welding in the 16:1 IPR samples promoted uniform TiC
distribution with loss of particle sphericity as a trade-off.
The particle sizes observed in the cryomilled samples are generally finer than the as-
received AlSi10Mg feedstock powder and other traditional PBF feedstock powder (Tan, Wong, &
Dalgarno, 2017a). Despite the disparity, recent work has suggested distributions with fine
particles may achieve superior densification after spreading in PBF (McGeary, 1961; Olakanmi,
Dalgarno, & Cochrane, n.d.; Zhu, Fuh, & Lu, 2007). Specifically, bimodal or trimodal distributions
with narrow peaks for coarse, medium-coarse, and fine particles properly fill interstitial gaps in
the powder bed. Cryomilling produced fines very quickly in this milling process. In addition to
incorporating reinforcement particles into matrix materials, it would also be suitable for tailoring
62
size distributions for maximizing packing density. The spreadability results from this study
indicate a visibly similar level of packing density to aluminum PBF feedstock powder may be
achieved through cryomilling. Combinations of additional milling stages or various sizes of
cryomilled powder may be appropriate methods for tuning size distribution for maximum
packing density.
These results also indicate an apparent strengthening from cryomilling likely due to two
mechanisms: the Orowan bowing mechanism promoted by fine dispersoids within the material
and Hall-Petch strengthening due to the high dislocation densities promoted by cryomilling.
These mechanisms have been described previously in mechanical alloying and metal matrix
composite literature (Chawla & Chawla, 2013; Groza et al., 2007; Hahn & Hwang, 2009; Lavernia,
Han, & Schoenung, 2008; D. B. Witkin & Lavernia, 2006). Orowan bowing occurs when a
dislocation cannot penetrate through a hard precipitate and instead must “bow” around it. The
resultant strengthening, a function of reinforcement size and spacing, elevates microhardness
with increasing reinforcement content and is the dominant strengthening mechanism in the 10
wt.% TiC samples. Composite homogeneity is also a driving factor in the Orowan mechanism,
which explains the 34% increase in strengthening between L-10TiC and H-10TiC. From EDS
imaging, L-10TiC featured superior TiC dispersion than H-10TiC after 3 hours of cryomilling
which explains the strength increase. Hall-Petch strengthening is likely occurring due to the grain
refinement inherent to the cryomilling process, however, no difference was observed in the 1
wt.% TiC samples, indicating IPR may not directly control the grain refinement process. In
summary, the ratio of the powder mass to impactor mass has an effect on the resultant mechanical
properties when the reinforcement fraction is high. The low TiC fraction samples were similar,
63
indicating IPR may not be as important of a factor in maximizing strength with lower
reinforcement loading.
It is also important to note the compressed and sintered specimens featured small spheres
that diffused onto the surface of the specimens during sintering. These appeared to be more
prevalent on the 1 wt.% TiC and the as-received AlSi10Mg samples. A brief EDS analysis
determined these diffused artifacts to be comprised of Al and Si. The chosen sintering
temperature was at the eutectic temperature for the Al-Si alloy system, so it is apparent liquid Al-
Si diffused out of the compressed specimens during sintering. These specimens appeared fully
dense during microhardness testing, however, it is possible that the preferential Al-Si diffusion
and decreased density within the 1 wt.% TiC specimens drove the lowered microhardness
compared to the 10 wt.% TiC specimens.
4.5 Conclusions
Composite powder fabricated from Al and nano-TiC was fabricated using a novel method
of cryomilling. Reinforcement weight fraction, powder charge in the milling vial, and milling
time were all explored for the fabrication of PBF-feasible composite feedstock powder. Higher
powder mass during cryomilling was found to retain the semi-spherical particle shape of as-
received AlSi10Mg powder at shorter milling times, however the TiC distribution was limited to
the surface of powder. Longer cryomilling produced more homogeneous powder, despite
producing additional fine particles. The spreadability of as-received AlSi10Mg and semi-
spherical H-1TiC and H-10TiC powder appeared to be similar indicating the feasibility of
cryomilled AMCs in PBF.
64
Pressed and sintered samples demonstrated elevated strengthening with increasing TiC
reinforcement driven by the Orowan bowing mechanism. Superior composite homogeneity in
low powder mass samples elevates this strengthening as evident by the disparity in
microhardness between L-10TiC-3h and H-10TiC-3h. Lack of disparity between microhardness
of the 1 wt.% TiC samples eliminates Hall-Petch strengthening as the primary cause of the 10
wt.% strength disparity.
Finally, combining the spreadability results with the apparent strengthening seen in
powder metallurgy alludes H-10TiC as the strongest and most feasible AMC from the current
study. The mixture of fine, angular particles with coarse, rounded particles after 6 hours of
cryomilling demonstrated similar spread patterns to traditional AlSi10Mg feedstock powder.
Likewise, superior strength compared to 1 wt.% TiC samples presents cryomilling as a novel
method for homogeneous reinforcement incorporation in the fabrication of AMC feedstock
powder for PBF additive manufacturing.
Acknowledgements
The authors would like to acknowledge Dr. Surendra Gupta in the Mechanical Engineering
Department at Rochester Institute of Technology for the expertise regarding x-ray diffraction
characterization. Likewise, the authors would like to acknowledge Samantha Sorondo for
assisting in characterization.
65
5. General Conclusions
This work successfully fabricated aluminum matrix composite (AMC) powder designed
for PBF-AM through a novel cryomilling process. The section summarizes the results of this
specific work and outlines interferences with results from research works on this topic of interest.
5.1 Conclusions
Linear cryogenic grinding was utilized in the fabrication of AMC feedstock powder in
order to expand the scope of high-strength materials for PBF-AM. This technology has been used
previously to produce an Al-CuO thermite composite, but remains otherwise unexplored for
aluminum composites (Badiola et al., 2009). Recently, attritor cryomilling was explored to
fabricate Al5083 powder for PBF with success, despite the inclusion of stearic acid, a process
control agent (Kellogg et al., 2019). To the best of the author’s knowledge, this is the first study of
its kind to utilize linear cryomilling in the fabrication of PBF-feasible AMCs sans lubricants.
In this work, particle morphology and size distribution varied slightly from traditional
AlSi10Mg PBF feedstock powder, but the spreadability of the powder was shown to be similar.
Semi-spherical AlSi10Mg-TiC composite powder was produced through manipulation of
cryomilling powder mass and milling time. Energy-dispersive x-ray spectroscopy confirmed the
presence of TiC in the interior of coarse and fine particles, indicating a homogeneous composite
was formed. A representative spreading test confirmed the ability of cryomilling to
homogeneously incorporate TiC particles within an aluminum matrix while retaining the
morphology and size distribution needed for appropriate spreading. The mix of semi-spherical
66
coarse particles with angular fine particles achieved using high powder masses in cryomilling
promoted spreading patterns similar to gas atomized AlSi10Mg feedstock.
Vickers microhardness measurements on powder metallurgy samples unveiled the
strengthening benefits of composite homogeneity and work hardening, both brought forth by
cryomilling. High powder mass samples demonstrated comparable microhardness to low
powder mass samples in addition to displaying superior powder spreading characteristics.
Higher TiC loading demonstrated higher microhardness and the ability to tune mechanical
properties through compositional changes.
5.2 Review of Contributions
In this study, linear cryomilling was utilized to fabricate custom aluminum matrix
composite feedstock powder designed for powder bed fusion additive manufacturing.
Calibration of cryomilling time and the powder mass allowed for spreadable, semi-spherical,
homogeneous composite powder to be fabricated. Using these guidelines, new alloy and
composite systems may be fabricated to widen the current scope of additive manufacturing
materials.
5.3 Future Perspectives
The immediate next step for this work is validation of cryomilled feedstock powder in a
metal PBF machine. This study utilized particle morphology, size distribution, and a
representative doctor blade spreading test as ex situ characterization methods of AMC powder
performance in PBF relative to a commercially available, widely used feedstock material. Future
work should also validate powder performance through consolidation in additive
67
manufacturing. Comparison of resultant part densification after AM-consolidation of as-received
AlSi10Mg and cryomilled AlSi10Mg-TiC materials would be a clear metric of the powder
performance. Likewise, verification of adequate wetting between AlSi10Mg and TiC during
localized melting is required to verify cryomilling as a feasible option for producing this
composite system.
As powder-based metal additive manufacturing transitions to an active role in the
manufacturing supply chain, material development will become increasingly important for
developing the next generation of high strength materials. In this specific study, only three
materials were tested in the cryomilling process: pure Al, an AlSi10Mg alloy, and nanoscale TiC.
An endless number of matrix and reinforcement combinations can be investigated, and the
resultant mechanical, thermal, and fatigue properties will vary. The unique phase transition
dynamics in contemporary metal additive manufacturing add to the variability in as-built
components. AlSi10Mg is the mainstream aluminum alloy for AM and was selected as it is known
to consolidate in PBF. To truly test the success of cryomilled nanocomposites in AM, traditionally
non-weldable aluminum alloys such as Al6061 or Al7075 should be cryomilled with a
reinforcement material and subsequently consolidated in AM. AlSi10Mg is reported to have a
Brinell Hardness of 119 ± 5, according to the EOS powder specification sheet (EOS-GmbH, 2014).
Al6061-O and Al6061-T6 have Brinell Hardness (500 g load, 10 mm ball) of 60 and 95, respectively
(ASM International Handbook Committee, 1990). The lower hardness value both with and
without temper lead the author to believe they would deform easily during cryomilling and
would be subject to excellent reinforcement dispersion. Al7075-O and Al7075-T6 have Brinell
Hardness (500 g load, 10 mm ball) of 60 and 150, respectively (ASM International Handbook
68
Committee, 1990). While Al7075 without temper would likely perform well during cryomilling,
the T6 temper would likely make cryomilling Al7075-T6 more difficult than AlSi10Mg. The
excellent hardness of this alloy after tempering would inhibit plastic deformation more than
AlSi10Mg and would likely need to be cryomilled longer. However, a cryomilling study using
Al6061 or non-tempered Al7075 would need to be performed to validate these conclusions.
Moreover, the materials in this study began in powder form. Because cryomilling will
reduce the particle size of any starting material through plastic deformation, it is not necessary
for the constituent materials to be in powder form. Future studies should investigate machining
refuse or chips as a potential starting material for the cryomilling process. Lower processing costs
and the potential for recyclability make this idea an attractive solution for fabricating custom PBF
powder. Recently, a group from Colorado State University have utilized room temperature
planetary milling to fabricate PBF-feasible stainless steel powder from machining waste (B.
Fullenwider, Kiani, Schoenung, & Ma, 2019; B. P. Fullenwider, 2018). Similarly, another group
has begun utilizing attritor cryomilling to repurpose Al5083 machining chips into aluminum
powder (Kellogg et al., 2019). Secondary or tertiary phases could easily be added to the milling
process to fabricate high strength composite materials for metal additive manufacturing.
Likewise, the preference for particle fracture in cryomilling would likely reduce the necessary
milling time to reach steady-state particle fracture from 60 hours as published (B. Fullenwider et
al., 2019). Further work in this area is needed to prove or disprove this hypothesis.
69
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