Additive Manufacturing Materials: Fabrication of Aluminum

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12-2-2019

Additive Manufacturing Materials: Fabrication of Aluminum Additive Manufacturing Materials: Fabrication of Aluminum

Matrix Composites Matrix Composites

Jakob Hamilton [email protected]

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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


Dr. Denis Cormier

Earl W. Brinkman Professor of Industrial and Systems Engineering and AMPrint Center Director


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

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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

viii

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

ix

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

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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

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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).

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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.

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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).

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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.

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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.

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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.

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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

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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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Additive Manufacturing Materials: Fabrication of Aluminum