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Development and characterization of hierarchical cellular structures

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Development and characterization of hierarchical cellular structures

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DEVELOPMENT AND CHARACTERIZATION OF
HIERARCHICAL CELLULAR STRUCTURES

by

Theresa Juarez

A Dissertation Presented to the
FACULTY OF THE USC GRADUATE SCHOOL
UNIVERSITY OF SOUTHERN CALIFORNIA
In Partial Fulfillment of the
Requirements for the Degree
DOCTOR OF PHILOSOPHY
(MECHANICAL ENGINEERING)

December 2017

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Acknowledgments

First, I would like to thank my advisor Professor Andrea Hodge for the guidance,
mentorship, and opportunities she provided me during my graduate studies. I cannot thank her
enough for the time and effort she has invested in both my professional and personal growth.
Thank you for the opportunity to work in your research group.
I am also grateful to the staff and faculty that have helped me navigate graduate school
at USC, especially Kim Klotz, Samantha Graves, Jennifer Gerson, and the AME department. I
would also like to acknowledge the Center of Electron Microscopy and Microanalysis
(CEMMA) at USC, John Curruli, and especially Dr. Matthew Mecklenberg for always
answering my questions and giving advice. Thank you also to Professor Michael Kassner,
Professor Ivan Bermejo-Moreno, and Professor Satyandra K. Gupta for agreeing to serve on
my dissertation committee.
Funding for this project was provided by the Air Force Office of Scientific Research
grant FA9550-14-1-0352. I would also like to thank the Department of Defense (DoD) for
support through the National Defense Science & Engineering Graduate Fellowship (NDSEG)
Program and the USC Graduate School for the Rose Hills Graduate Fellowship.
This work would not have been possible without my distinguished collaborators
including Almut Schroer and Dr. Ruth Schwaiger at the Karlsruhe Institute of Technology,
Professor Lorenzo Valdevit at UC Irvine, Dr, James Oakdale and Dr. Juergen Biener at
Lawrence Livermore National Laboratory, and Professor Jörg Weissmüller at the Hamburg
University of Technology. Also, thank you to Dr. Chris Bruser at Droseran Inc. for the
microtome and TEM work.
To my Hodge Group family: Dr. I-Chung Chen, Dr. Tim Furnish, Dr. Mikhail Polyakov,
Dr. Leo Velasco, and Dr. Nathan Heckman for teaching me and reminding me to stay the
course. Sebastián Riaño, Joel Bahena, Chelsea Appleget, Andrew Lindo, and Alina Garcia-
Taormina for all the food and lab adventures. Dr. Kamia Smith and Dr. Thien Phan for
encouragement and lots of laughs. Thanks to all of you for never hesitating to help me. It was
an honor to work with you all and I sincerely wish you the best in all your future endeavors.
I would also like to thank all my family and friends for their love and support. Thank
you to Daniel and Anais for being there during both good and bad times, despite the distance
you never felt very far away. Thanks to my brothers, Tommy and Steven, for making me laugh
when I needed to smile. Thank you especially to my parents, Sylvia and Tommy Juarez, for
showing me the value of hard work and dedication, for never letting me give up, and for always
encouraging me to pursue my dreams.
Finally, thank you to my dearest husband Shane, for your unfailing love, patience,
encouragement, and support. You believed in me, kept me sane, and provided me with the
happiness I needed these past five years. You stuck by my side throughout this journey and I
couldn’t have asked for better company along the way.

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Table of Contents
List of Figures ......................................................................................................................... viii
List of Tables.............................................................................................................................xv
Abstract .................................................................................................................................. xvii
Chapter 1 : Introduction ..............................................................................................................1
Chapter 2 : Background ..............................................................................................................3
2.1. Cellular Materials ............................................................................................................3
2.2. Types of Cellular Materials .............................................................................................5
2.3. Mechanical Behavior of Cellular Solids .........................................................................6
2.4. Fundamental Fabrication Methods ..................................................................................9
2.5. Synthesis of Hierarchical materials ...............................................................................19
2.6. Mechanical Behavior of Architected Materials ............................................................27
2.7. Applications of Hierarchical Materials .........................................................................29
Chapter 3 : Experimental Methods and Materials.....................................................................35
3.1. Electrodeposition ...........................................................................................................35
3.2. Heat Treatments ............................................................................................................41
3.3. Magnetron Sputtering ....................................................................................................41
3.4. 3D Direct Laser Writing ................................................................................................46
3.5. Characterization Techniques .........................................................................................47
3.6. Mechanical Testing of Microcellular Materials ............................................................50
Chapter 4 : Review of Nanoporous Materials with Structural Hierarchy .................................55
4.1. Hierarchical np materials produced through templating ...............................................56
4.2. Dealloying methods to produce hierarchical np metals ................................................59
4.3. Composite Structures based on nanoporous metals ......................................................65
4.4. Conclusions ...................................................................................................................70
Chapter 5 : Synthesis of Hierarchical Nanoporous Gold Structures .........................................73
5.1. Experimental methods for np gold tubes .......................................................................74
5.2. Results ...........................................................................................................................76
5.3. Effect of electrodeposition parameters ..........................................................................78
5.4. Post-process heat treatments for ligament coarsening ..................................................80
5.5. Discussion .....................................................................................................................82
5.6. Summary .......................................................................................................................83
Chapter 6 : Fabrication and Mechanical Behavior of Metal Coated Cellular Structures..........86
6.1. Experimental Methods ..................................................................................................88
6.2. Results and Discussion ..................................................................................................93
6.3. Outlook and Conclusion ..............................................................................................108

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Chapter 7 : Tensile Behavior of Cellular Materials Produced via Direct Laser Writing ........110
7.1. Experimental Methods ................................................................................................111
7.2. Results and Discussion ................................................................................................117
Chapter 8 : Conclusions and Future Work ..............................................................................123
References ...............................................................................................................................129
Appendix A: Complex Nanoporous Structures.......................................................................151
Appendix B: Summary of Nanoporous Gold Tube Samples ..................................................158
Appendix C: Summary of Coated Tetrahedral Truss Samples ...............................................161
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List of Figures
Figure 1: Ashby plot of strength vs. density for common engineering materials and novel
hierarchical architectures such as microlattices or honeycombs. Reproduced from [14]. ..........4
Figure 2: Classifications of cellular materials as either stochastic (a-c) or periodic (d-f), with
cellular features from the macroscale down to the nanoscale for each type of cellular structure.
f) is from reference [22]. .............................................................................................................6
Figure 3: Schematic representations of stress-strain compressive loading curves for bending
dominated structures (a) and stretch dominated structures (b) and open-cell failure modes (c)
buckling, and (d) brittle failure. Adapted from [43]. ..................................................................9
Figure 4 : Chart of established methods to produce micro and nanoporous metallic foams.
Adapted from [44]. ....................................................................................................................11
Figure 5: SEM images of np-Au foam prepared by (a) free corrosion and (b) electrochemically
driven corrosion. SEM images of free corrosion samples after 2 hour heat treatments at (c) 400°
C and (d) 600° C. Adapted from [78]. ......................................................................................15
Figure 6: Additive manufacturing processes to generate complex architectures. From left, a)
SLM generated Ti truss spinal implants [96], b) MICA Freeform made metal structures [97],
and c) 3D-DLW built micro/nano sized epoxy based polymer trusses [98]. ............................17
Figure 7: SEM images of a negative Poisson ratio structure that contracts when compressed.
Structures are produced by two-photon lithography using the Nanoscribe. Adapted from [100].
...................................................................................................................................................18
Figure 8: (a) Mask UV lithography using a polymer waveguide was to generate a (b) polymer
lattice. Adapted from [101]. ......................................................................................................19
Figure 9: Alloy opal structures (a) are dealloyed in nitric acid to develop a porous opal structure
(b). The overall film structure (c) is approximately 2 µm thick. Adapted from [106]. ............21
Figure 10: (a) Nested np structure generated by depositing an alloying element onto a np metal
followed by diffusion and dealloying [107]. (b) Nested np metal by two-step dealloying of an
alloy with abnormally high LNE content [108]. .......................................................................22
Figure 11: Schematic of the hierarchical scales used in metallic microlattices [19]. ...............23
Figure 12: Schematic showing the multi-step templating process used to synthesize hollow
microlattices. (A) UV waveguides were used to fabricate templates. (B) The polymer templates
were coated with nickel and etched resulting in hollow lattices (C). Reproduced from [50]. ..24
Figure 13: Hollow lattice structures developed by coating polymer materials generated by DLW
and removal of the polymer material. From left to right, a) ceramic nanolattice [111], b) Au
nanolattice [112], and c) Cu mesolattice [113]. ........................................................................25
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Figure 14: Tilted and top view SEM images of microcellular architectures after NiB electroless
deposition. The scale bar in all the images is 10 µm [114]. ......................................................26
Figure 15: In-situ X-ray tomography of nickel microlattice under compression. Note that the
deformation of the structure is concentrated at the nodes. Adapted from [122]. ......................28
Figure 16: Electrolytic cell for the deposition of silver from an aqueous silver nitrate solution.
Silver ions are reduced on the surface of a cathode as a metal film. ........................................36
Figure 17: (a) Schematic of an idealized electric field connecting two electrodes in an
electrolyte solutions. (b) Shows the actual electric field for two parallel electrodes with field
line concentrations near the edges of the materials. In practice, the areas with higher electric
field line concentration are regions with a higher localized current density. Adapted from [164].
...................................................................................................................................................37
Figure 18: Schematics showing the experimental electrodeposition setup. The relative
placement of the electrodes are shown in (a). (b) shows an image of the entire system including
the anode, cathode, reference electrode, and filter. ...................................................................38
Figure 19: Schematic of an updated experimental electrodeposition setup with two semicircular
anodes, the cathode, and a filter. ...............................................................................................38
Figure 20: (a) Electrode mounting piece connected to the Au and Ag electrodes out of the
electrolyte. (b) The entire system placed in a 2 L beaker with 1 L of Caswell Ag plating solution.
...................................................................................................................................................39
Figure 21: (Top) Au tube connected to a Cu wire with Ag pain as solder and epoxy to mask the
joint. (Bottom) Au tube coated with a layer of Ag. ..................................................................40
Figure 22: (A) MTI-GSL1100X tube furnace used for homogenizations heat treatments. (B)
Vulcan 3-550 furnace used for np Au ligament coarsening treatments. ...................................41
Figure 23: Simplified schematic showing the set-up and process of magnetron sputtering: a
neutral gas is introduced and ionized by the negatively biased target; the gas ions strike the
target and cause the target atoms to be ejected then coat the substrate [165]. ..........................42
Figure 24: Schematic of the DC magnetron sputtering chambers used in this study. Multiple
sputtering sources can be used simultaneously in this configuration. The attached DC motor
can be used to rotate substrates during coating or for direct line-of-sight to a specific target.
Reproduced from [166]. ............................................................................................................43
Figure 25: Image showing three targets used for magnetron sputtering in this study. From the
image and entrenched “race track” it is apparent that the Inconel 600 target has been sputtered
more than the Al or Ti-6Al-4V targets. .....................................................................................44
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Figure 26: a) Shows the representative sputtering schematic for stationary sputtering. b) Shows
schematic for rotating sputtering set up with samples mounted in the center of the rotating
holder. Sample is rotated 54 times per minute. .........................................................................45
Figure 27: a) An image of the Nanoscribe GmbH Photonic Professional GT laser instrument
Reproduced from [167]. b) Illustrates laser beam interaction with resist to produce 3D
structures. ..................................................................................................................................47
Figure 28: Schematic of the electron beam and specimen atom interaction volume under the
specimen surface. Reproduced from [168]. ..............................................................................48
Figure 29: Photo of micromechanical test set-up with interchangeable testing stages and
attached camera for digital image correlation to track sample deformation. ............................51
Figure 30: a) Photo of Agilent G200 XP nanoindenter used to conduct uniaxial compression
tests of microtruss structures [172]. b) Schematic of a compression test using a diamond flat
punch tip moving at a constant displacement rate. ...................................................................52
Figure 31: Representative stress-strain curve of tetrahedral truss structures. The stiffness is
calculated from the elastic regime, while the compressive strength is the maximum stress. ...53
Figure 32: Schematic representing various hierarchical structures in single element np materials
[108, 135, 173-175] (A-F) and composite structures [116, 140, 176-178] (G-K) that are based
on np metal morphologies [73, 75, 78, 179] (L-O) produced by dealloying. ...........................56
Figure 33: Various nanowire structures ranging from short np wires (a) to np nanotubes (b) to
long nanowires fabricated on silicon patterned substrates (c & d). Reproduced from ref [104,
117, 182]. ..................................................................................................................................58
Figure 34: Cross-section of alloy foam before dealloying. (a) Large block shaped macropores
from salt removal. (b) Enlarged area showing micropores from gas expansion. Reproduced from
[197]. .........................................................................................................................................62
Figure 35: Bimodal Cu rods developed by a combination of the Gasar and dealloying processes
to produce nanopores (a) within longitudinal pores (b). Reproduced from [199]. ...................64
Figure 36: EBSD map showing the grains of a textured alloy after cold rolling. b) Low-
magnification SEM image of layered structure after dealloying. c) Higher magnification SEM
image showing the layered structure with embedded nanoporosity. Reproduced from [135]..65
Figure 37: Left; color enhanced SEM image of PANI coated np-Au. Right; TEM image of
PANI layer on the surface of np-Au. Reproduced from [213]. .................................................68
Figure 38: Schematic representation of the general process of depositing np metal films onto
cellular structures. First, an alloy is deposited or sputtered onto a porous template. The alloy is
then dealloyed so that the remaining material is a macroporous foam with a nanoporous coating.
Reproduced from [178]. ............................................................................................................69
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Figure 39: Examples of np metals mounted on cellular scaffolds. (a,b) np Ag on porous carbon
produced by sputtering of AlAg then dealloying [178]. (c,d) np-Au on commercial Ni foams
produced by electrodeposition of AuSn then dealloying [177]. ...............................................69
Figure 40: Visualized proof of concept. Hollow cylinders as lattice struts with an embedded
nanoporous structure. ................................................................................................................74
Figure 41: Schematic representation and cross sectional images for the processing of
nanoporous gold tubes. Fabrication begins with a base Au tube (a) that is coated with a layer of
Ag via electrodeposition (b). The layered tube is homogenized in a furnace (c) and then
dealloyed to produce a nanoporous Au tube (d). The cross-section images for processing at each
step are shown in (e-h). .............................................................................................................76
Figure 42: (a) Cross sectional micrograph of homogenized AuAg alloy. (b) EDS spectra
confirming that the alloy is within the target compositional range. ..........................................77
Figure 43: Electrodeposition with direct current and a current density of 21.6 Acm
-2
showing a
rough silver top surface (a). Homogenization and dealloying led to a bimodal tubular structure
with an inner nanoporous network (b) and an outer macroporous region (c) that also contained
a nanoporous network inside larger ligaments. Pulsed current with polarity reversal at a current
density of 10.8 mA cm
-2
promoted even plating (d). SEM images of an optimized nanoporous
Au tube with np morphology inside structure walls (f). ...........................................................78
Figure 44: Tunable ligament sizes for nanoporous Au tubes (a) as-prepared and after 2 hour
heat treatments at (b) 200 °C and (c) 400 °C. ...........................................................................81
Figure 45: Schematic showing the design and development of polymer-metal truss structures.
a)-b) A tetrahedral truss structure model is designed and subsequently fabricated using 3D
direct laser writing. c) After being written, the structures are coated with a metal film using
magnetron sputtering. Inset shows schematic of a strut cross section. (d) SEM images of the
actual structures after processing is complete. ..........................................................................89
Figure 46: Sputtering setup: Left shows SEM image of sample layout on the glass slide. Middle
shows the representative sputtering schematic for stationary sputtering. Right shows schematic
for rotating sputtering setup with samples mounted in the center of the rotating holder. Sample
is rotated 54 times per minute. ..................................................................................................91
Figure 47: a-c) SEM images of sputtered coatings on tetrahedral truss structures: Aluminum,
Inconel 600, and Ti 6Al-4V. Below each image is a corresponding XRD scan and EDS spectra
for each type of coating from a 1 micron thick film sputtered in the same conditions. Note that
the sputtered alloys have maintained stoichiometry. ................................................................94
Figure 48: a) FIB image of a tetrahedral truss structure milled in half to expose the coating cross
section. b) SEM image of a strut cross section exposed with FIB milling. The actual coating
material is obscured by material re-deposition from the part of the truss structure that was
eroded during FIB milling. ........................................................................................................96
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Figure 49: Microtome prepared cross sections of Al coated structures with either the stationary
(a) or rotating (b) configurations. Images shown are for structure size D with 12.5 µm strut
length. The red arrows for both images show the thicker top side coatings, while the white
arrows show the thicker outer coatings generated with the rotating sample. ............................99
Figure 50: TEM micrographs of strut cross-sections coated in the stationary setup from center
of the structure (a) and from the bottom of the structure (b), which was masked during
sputtering. ................................................................................................................................100
Figure 51: TEM images of strut cross sections of structure A rotated and coated with Al. The
zoomed in images of the struts are arranged according to the cross section and exhibit a
decreasing top surface thickness of the coating at the lower levels of the truss structure. .....103
Figure 52: Measurements of the top side of the struts and the corresponding area of the structure
for the rotated Al structure corresponding to level of structure where the measurement was
taken. In both samples, the thickness of the top level becomes smaller as the measurements are
taken at lower levels in the samples. .......................................................................................104
Figure 53: Uniaxial compression stress-strain curves showing the range of mechanical behavior
observed for tetrahedral truss structures of type A coated with different metals. The stress-strain
curves are shifted to account for misalignment between tip and sample surface or roughness
effects (gray region in the graph). ...........................................................................................106
Figure 54: Helium ion microscopy (HIM) images of compressed truss structures of type A
coated with a ~35 nm layer Inconel 600 in stationary configuration. a) Structure failure at a
truss node is shown and marked with circle. The arrows note failure along the ligaments. b)
Close up image of structure node where the coating is non-conformal and forms a webbed
structural feature......................................................................................................................107
Figure 55: Schematic of log-pile structures printed on a glass slide. a) Shows a layout of the
printed structures. b) Cross-section of the geometry showing the dimensions of the tensile
specimens. c) Dimetric view of the tensile specimen showing overall shape. .......................113
Figure 56: Schematics showing two proposed methods for gripping tensile specimens for testing
micropillar specimens. a) Shows an adhesive method where one end of the sample is attached
to a crosshead using an adhesive. b) Schematic shows a mechanical grip where specimens are
placed into a slot piece and pulled until the sample is in tension. ..........................................114
Figure 57: Curve showing the change in load as a tensile sample is attached to the tensile test
crosshead with a UV cure silica filled epoxy. The load increases as the adhesive cures and
continues to increase, eventually causing failure of the sample before a tensile test could be
completed. ...............................................................................................................................114
Figure 58: Schematic showing the alignment of tensile specimens with slot piece to grip for
tensile tests. a) First the tensile gage section is aligned with the slot and moved above the lock
piece. b) After alignment, the lock-piece is moved into position under the tensile sample. c) The
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sample is then lowered into the slot and d) the slot piece is moved into a position where the
sample is engaged in tension. ..................................................................................................115
Figure 59: Load cell assembly for testing microcellular polymer tensile samples. a) Schematic
profile view of load cell. b) To scale schematic of the load cell, tilted so that the samples are
visible. c) Photo of load cell with tensile samples printed on the glass slide. d) Photo of complete
assembly with both the load cell and tensile samples, as well as the slot piece slide. ............116
Figure 60: Load vs. displacement data from pillars tested in tension. Blue curves show
continously printed pillars and red curves shows samples printed with a rest period half way
through the printing process. ...................................................................................................117
Figure 61: Load vs. displacement data from pillars tested in tension. Green curves show samples
printed with a rest period and printing overlap. These samples showed the highest maximum
achieved load before failure. ...................................................................................................118
Figure 62: Photographs from camera mounted on top of the microtensile tester of each of the
three structures during testing. First column of images show samples prior to failure. The second
column shows images at the moment of failure and the final set of images show the samples
after the slot piece has been moved away from the specimen. ...............................................119
Figure 63: Polymer truss structures coated in Au by magnetron sputtering. (a) single strut (b) a
unit cell (c) lattice structure made of repeating cells. .............................................................151
Figure 64: Cross section of polymer template coated with Au by magnetron sputtering. ......152
Figure 65: Processing of complex np structures (a) fabrication begins with generating a gold
base structure (b) a layer of Ag is deposited on the gold structure (c) the metals are diffused
during homogenization distorting the structure. .....................................................................153
Figure 66: Selection of different hollow strut wall architectures. (a) Polymer infiltrated sample,
(b) tube wall with a single face sheet, (c) and sandwich structure with foam core. ...............154
Figure 67: Schematic representation of a complex hierarchical structure made of hollow struts
with a metal-foam sandwich wall (inset on the right). ............................................................155
Figure 68: a) Ta tube coated with Au via electrodeposition b) Ta tube coated with Ag via
electrodeposition. ....................................................................................................................156
Figure 69: Bi-layer Ag-Au tube produced by separately plating Au then Ag onto a Ta base tube.
.................................................................................................................................................156
Figure 70: SEM image of AgAu coated Ta tube after homogenization heat treatment. The image
shows that the AuAg layers diffused, however the AgAu alloy began to pull away from the Ta
tube base and was no longer a conformal coating...................................................................157

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List of Tables
Table 1. Electrochemical series of metals .................................................................................13
Table 2: Ligament sizes of heat treated nanoporous tubes .......................................................81
Table 3: Summary of polymer tetrahedral truss dimensions per type of structure. ..................90
Table 4: Summary of sputtering targets and deposition parameters. Note that all the structures
are coated simultaneously in each sputtering procedure. ..........................................................93
Table 5: Description of samples and printing parameters with rest periods and overlaps. .....112
Table 6: Maximum load and displacement of pillars just before sample fracture. Pillar C, which
exhibited the hightest failure load, is in bold. .........................................................................120

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Abstract
Lightweight hierarchical materials, which have tunable structural features on multiple
length scales, have been a topic of increased research interest due to advances in techniques for
fabricating these cellular solids with nano and microscale features. The study of these materials
extends to a variety of applications including catalysis, optics, and batteries. However,
processing and mechanical stability remain an integral part of expanding and scaling up the
overall size of hierarchical structures with nano and microscale features at the base level of
structural hierarchy.
In this study, fabrication methods, characterization, and mechanical testing of
hierarchical materials is explored on different length scales. The initial focus of this work
explores a process called dealloying to develop hierarchical nanoporous (np) metal structures
with nanopores on the smallest feature length scale and a complex shape on the macroscale.
Using a step-by-step fabrication method, a binary alloy tube was developed and subsequently
dealloyed to produce a np metal tube. These tubes were generated as a proof of concept exercise
so that similar step-by-step procedures could be applied to more complex shapes.
In addition to templating and dealloying as a method to produce hierarchical cellular
materials, microlattice metal-polymer composites were generated using a combination of 3D
direct laser writing (3D-DLW) and metal coating via magnetron sputtering. Using magnetron
sputtering, metal coatings of Aluminum, Inconel 600, and Ti 6Al-4V were deposited onto
microlattice structures. The cross section of the composite microlattice materials were evaluated
using microtome, a microscopy preparation method where the structure is embedded in epoxy
and thinly sliced, which allows for the assessment of cross sections of individual lattice struts.
The inclusion of alloy coatings is significant because traditional deposition methods for metals
xviii

are limited to a few select elements. Therefore, using sputtering as a deposition method has the
potential to expand the achievable properties of composite microstruss structures. Uniaxial
compression of the samples revealed that the compressive strength can be influenced by the
metal that is deposited. Furthermore, by applying metal coatings, the deformation mode of the
composite structures is shifted from node-dominated failure towards an increase in ligament
failure.
The final aspect of this dissertation discusses the scaling up of microcellular polymer
materials produced with 3D-DLW, which has a limited print volume. In order to increase the
size of the printed structures, multiple printed areas are “stitched” together to form a larger
structure. The influence of the stitching along this interface on the mechanical behavior of
cellular structures was explored. The tensile behavior of the samples was compared to
continuously printed samples with no stitch. A custom micromechanical setup was used to
complete tensile testing of the structures, and no pronounced tensile strength difference between
stitched and continuously printed samples was found, although stitched samples regularly failed
along the stitch interface.
Combined, these results have addressed three critical aspects of the emerging field of
hierarchical materials, including: 1) developing novel methods to incorporate nanoscale
features in hierarchical materials; 2) combining multiple techniques towards functionalizing
polymer architected materials in a way that affords a high degree of material flexibility; and 3)
exploring issues with scaling up architected materials with microscale features so that 3D
printing protocols can be improved to mitigate these effects. Overall the main goal of this work
is to reform and contribute novel processing methods for designing and fabricating materials
on multiple length scales.
1

Chapter 1 : Introduction
The concept of material hierarchy is ubiquitous, encompassing both deliberately
designed structures as well as naturally occurring materials such as wood or bone. Careful
organization of the structural features can introduce a range of tunable properties into the
material, such as density or surface area, which can be optimized for specific applications.
Examples of ordered hierarchy are most evident in biological systems, where pore networks in
trees and bones reduce the density of the structure while providing a path for nutrient and fluid
flow, enhancing mechanical properties while adding functionality [1]. Biological systems are
especially complex because they employ hierarchy from the nanometer up to the macroscale,
organized so that the structural elements themselves have structure [2]. Humans have attempted
to mimic advanced biological systems by designing frame structures like the Eiffel Tower, but
the complexity that can be found in natural materials, such as bamboo, has yet to be realized
[3].
There may be opportunities in reproducing these types of structures in manmade
materials, where hierarchy is found in the length scales occupied by structural features, such as
tiered porosity in single element materials, or the spatial distribution of constituents in
composites. However, this requires the development of new strategies for advanced materials
design, specifically feasible preparation methods and an understanding of the effects of
hierarchical structure on functional properties. Several approaches can be used to achieve this
feat. For example, advances in metal or ceramic deposition methods, templating, and dealloying
methods have made it possible to produce hierarchical structures in a variety of elements and
configurations.
2

A subset of hierarchical materials is architected cellular lattice structures, which are
made up of ligaments or rods configured in a unit cell that is repeated in space. Current additive
manufacturing technology has made it possible to synthesize architected structures with
features on the nano and microscale in a hierarchical configuration. This has allowed scientists
to study previously unknown phenomena, such as pronounced size effects and unexpected
mechanical behavior exhibited by structures with feature shape and size control not previously
available. These interesting behaviors include ductility in traditionally brittle base materials
such as ceramics. However, lattice or truss structures generally fail at the nodes, indicating poor
load transfer to the ligaments. Further optimization of hierarchical materials should be studied
to improve the deformation behavior and mechanical efficiency of these materials.
The present study will focus on using established synthesis techniques such as
electrodeposition, sputtering, and additive manufacturing to produce protocols for developing
novel architected metallic structures. Micromechanical testing methods will be adapted to
evaluate the developed structures and inform further optimization of these nanostructured
materials. Finally, initial work on scaling up microcellular additive materials will be presented.
Together these results will help us understand the relationship between the architecture,
microstructure, and mechanical properties that result from designing and fabricating materials
on multiple length scales.
3

Chapter 2 : Background
2.1. Cellular Materials
Porous metals with porosity on the order of nanometers to millimeters are used in a
variety of structural and functional applications including heat exchangers [4, 5], filters [6, 7],
energy absorbers [8, 9], and catalysts [10, 11] due to their high surface area, as well as their
high thermal and electrical conductivity [12, 13]. In particular, the automotive, construction,
and aerospace industries are interested in foams as lightweight structural components to
maximize strength-to-weight-ratios, impact absorption, and vibrational damping [14, 15].
Incorporating additional length scales or architectures can significantly improve the
performance of cellular materials. Furthermore, these materials provide design advantages due
to their high tailorability for specific applications, expanding the potential applications when
compared to fully dense materials. A prime example of this is the design of sandwich panels
and their core structure to enhance both mechanical properties and heat transfer properties when
a fluid flows through the panel [16-18]. Potential core structures include traditional foam core,
honeycombs, or truss structures with periodic architecture. The variety of available core designs
and material options allows for the optimization of these metal cellular structures for
applications in the aerospace and automotive industry, where the panels are typically loaded in
bending [16, 19].
Despite their varied use, a key requirement of cellular solids in most industries is weight
minimization without the sacrifice of mechanical performance. Foams can usually sustain large
compressive strains, absorbing large amounts of energy without high stresses. Figure 1 shows
an Ashby plot, or material selection plot, illustrating the strength-to-density performance of
several common engineering materials. A major objective of this study is to design light-weight
4

materials to fill the unoccupied region in the plot corresponding to high strength and low density
with novel cellular architectures. In general, engineered structures, such as the lattices,
correspond to stretch-dominated behavior and accomplish higher strength-to-weight ratios than
random cellular foams and solid materials. Commercial cellular materials, such as honeycombs
or technical foams, can be made from glass, polymers, ceramics, or metals. Cellular structures
are also found in naturally occurring materials.

Figure 1: Ashby plot of strength vs. density for common engineering materials and novel
hierarchical architectures such as microlattices or honeycombs. Reproduced from [14].

5

2.2. Types of Cellular Materials
Cellular solids can be classified into two major categories, stochastic and periodic, see
the examples in Figure 2. Stochastic materials (Figure 2a-c) have a random pore distribution,
whereas periodic structures (Figure 2d-f) have a base cell that is repeated in space, such as
honeycombs (2D) or trusses (3D) [14, 15, 20]. Stochastic materials are considered isotropic
cellular solids which are characterized by structure and (compressive) properties with no
directionality. The honeycombs in Figure 2f for example would have compressive anisotropy,
where compressive response strongly varies with loading direction. Each category is associated
with distinct mechanical properties as well as differences in relative density and production
cost. Both stochastic and periodic foams can be fabricated with a wide range of metals or alloys,
from aluminum to Inconel, and in configurations such as closed-cell and open-cell, which are
permeable to fluids and have a lower density [14, 21]. The closed-cell solids are non-permeable
by liquids or air and are often “inflated” with trapped gas. The non-permeable foams have a
higher density and are more rigid, making them useful for structural applications.

6

Figure 2: Classifications of cellular materials as either stochastic (a-c) or periodic (d-f), with
cellular features from the macroscale down to the nanoscale for each type of cellular structure.
f) is from reference [22].

2.3. Mechanical Behavior of Cellular Solids
The mechanical properties of stochastic and periodic structures have been well studied.
For example, one finds that for a given relative density, periodic structures have superior load-
bearing properties when compared to their stochastic counterparts [23]. This is because periodic
trusses or networks are typically stretch-dominated materials when loaded, whereas random
networks are bending-dominated [24]. Additionally, closed-cell materials typically have an
increased relative modulus for a given relative density, making them more desirable for load
7

bearing applications [8, 15, 23]. The main advantages of using open porous structures in
practice includes mass reduction and the large surface area which can be accessed by fluids
permeating the pores [14, 15]. This is crucial for applications in catalysis [25], where product
formation is a function of available surface area, or applications requiring cooling, since heat
transfer with a high amount of conductive surface area is enhanced [26, 27]. For the purposes
of this study, only open cell structures are explored, since their large surface area enhances the
potential for added functional applications, such as catalysis and energy storage.
The fundamental mechanical behavior of cellular solids is described in the seminal work
by Gibson and Ashby [28]. The Gibson-Ashby approach predicts that mechanical properties do
not depend on the cell size, except for a slight dependence of strength to the ratio of the cell
wall thickness to cell length [28]. The main parameters in determining the mechanical
properties of the foam are the base material properties and the foam’s relative density, which is
shown in Equation 1:
𝜌 ∗
𝜌 𝑠 Relative Density (1)
where (*) refers to the foam properties and (s) refers to the properties of the solid from which
the foam is made.
This relative density is used to determine the Young’s modulus (E) and the yield
strength (σ) of macroporous cellular foams, represented by struts connected at joints, with the
Gibson-Ashby foam scaling equations shown below:
𝐸 ∗
= 𝐶 1
𝐸 𝑠 (
𝜌 ∗
𝜌 𝑠 )
2
Young’s Modulus (2)
𝜎 ∗
= 𝐶 2
𝜎 𝑠 (
𝜌 ∗
𝜌 𝑠 )
3/2
Yield Strength (3)
8

where ρ is density, and C1 and C2 are geometrical constants.
It is therefore unexpected that the properties of the np metals discussed later in this work
do not scale according to the Gibson and Ashby equations [29-40]. The discrepancies are
highlighted when considering the apparent changes in strength with increasing or decreasing
ligament size in np foams with the same relative density [30, 31]. While studies on the
mechanical behavior of nanoporous metals are numerous and suggest interesting size effects,
the influence of structural hierarchy on strength has only been thoroughly addressed from the
nm scale up in studies on biological materials [41, 42].
Generally, researchers are concerned with the compressive deformation of foams due to
their favorable properties in compressive loading conditions. Figure 3a shows a typical
compressive stress-strain curve for a bending dominated lattice, which represents most existing
foams. When compressed the structure behaves elastically with modulus E
*
, which is
designated as Ē in the figure below, until it begins to deform plastically. The foam then
continues to deform at a nearly constant “plateau” stress and begins to densify when the void
space in the cell becomes increasingly occupied by cell material [43].
In contrast, stretch dominated structures like truss lattices, which will be discussed in a
later section, can achieve much higher strength. Figure 3b shows a representative loading curve
for stretch dominated materials. Like the bending dominated structures, the foam initially
deforms elastically with modulus E
*
, but after the onset of plasticity there is a region of material
softening when struts buckle or collapse [43]. Densification occurs similarly to bending
dominated materials. There are different yielding modes for cellular structures. However, the
main failure modes observed in this study are buckling behavior and brittle failure shown in
Figure 3c and 3d respectively.
9

Figure 3: Schematic representations of stress-strain compressive loading curves for bending
dominated structures (a) and stretch dominated structures (b) and open-cell failure modes (c)
buckling, and (d) brittle failure. Adapted from [43].

2.4. Fundamental Fabrication Methods
In this section, we discuss fundamental fabrication methods for cellular structures
pertinent to this study. First, an overview of traditional methods associated with commercial
and widely available metal foams is given. This is followed by detailed explanations of the
a) b)
c)
d)
10

synthesis of nanoporous metals and an overview of architected materials generated by additive
manufacturing or 3D printing methods.

2.4.1. Porous Metal Foams

The synthesis of cellular metal structures can be divided into three types; top-down,
bottom-up, and hybrid techniques. Top-down approaches, such as dealloying, begin with a bulk
material that is deformed or reduced until the desired final structure is achieved. In bottom-up
processing, materials are built in a specific configuration by adding material, usually in layers.
Hybrid processing uses a combination of removing and adding material, such as using a
sacrificial template that is coated with material and removed before processing is finalized.
Figure 4 summarizes established synthesis techniques reviewed by Davies and Zhen for metal
foams, most of which develop macrocellular structures [44].
The earliest commercial foams can be traced back to the 1920s, when French scientist
De Meller described a process of injecting gas into molten metals [45]. Casting or liquid-state
processing, where molten metal is either agitated or aerated with a foaming agent, is still the
most common method for producing metal foams. In one commercial system, Al-SiC foams
are made by introducing gas and SiC particles into molten aluminum. The bubbles produced by
the gas form the foam and the ceramic particles stabilize the system [46, 47].
11

Figure 4 : Chart of established methods to produce micro and nanoporous metallic foams.
Adapted from [44].

Alternative processing methods use powder sintering, where fine metal powders are
compacted with a metal hydride and heat treated near the melting temperature of the metal. This
releases gas from the hydride and forms pores in the soft metal [9]. Aluminum and steel foams,
which offer higher strengths and thermal stability compared to Al, can be made using this
method [48]. A simple approach towards enhancing the properties of metal foams includes
careful selection of the base metal alloy to produce strengthening intermetallics [49].
It is important to note that these foaming techniques form mm sized pores and are
generally limited to macroscale porosity (Figure 2a). Furthermore, these techniques offer
limited ability to impart additional levels of structural hierarchy other than material
characteristics such as grain size and dispersed particles. While some work has been done to
form microscale porous foams (1-100 microns) using metallic melts, the established method for
producing finer pores (~100s of microns) in metallic foams involves the deposition of metal
onto a polymer scaffold and subsequent etching or burning out of the polymer, resulting in a
hollow strut metal foam. Unlike the foaming methods, this approach can be tailored by using
engineered microlattice polymer scaffolds [50, 51]. and by controlling the amount of metal
12

deposited on the scaffold, allowing some influence on the wall thickness. Thus far, traditional
preparation methods for metal foams have not been able to achieve cell sizes below the micron
scale.

2.4.2. Nanoporous Metals

Recent fundamental work on the synthesis, properties, and applications of nanoporous
metals, with feature sizes in the tens of nanometers, has generated a completely new field and
class of porous metals. In particular, dealloying has been established as the main synthesis route
towards monolithic np metals by providing both compositional flexibility and a high level of
morphological control. Np metals made by dealloying lack the cell structure that is
characteristic of materials produced by foaming. Instead, the new class of materials is
characterized by an interconnected network of nano or microscale ligaments, interpenetrated
by continuous pore channels. The smaller featured networks exhibit high stiffness, like their
macroporous counterparts. However, a key distinction for np metal resides in their drastically
increased surface area, which scales inversely with feature size. The large surface area makes
them uniquely suited as catalysts [11, 52-55], sensors [11, 56, 57], and actuators [58, 59].
The mechanisms of dealloying to produce np metals have been explored since the early
20th century [60, 61], while the fundamental and more recent work towards synthesizing np
metals were done by Pickering and Swan [62, 63], Forty [64], as well as Newman and Sieradzki
[65, 66]. Essentially, dealloying is the selective electrolytic corrosion of binary or ternary alloys
made of a more noble element (MNE) and a less noble element (LNE), where the LNE is
selectively leached from the alloy. Removal of the LNE is achieved by either free-corrosion
dealloying through submersion of the alloy in a corrosive solution, or electrochemical
13

dealloying, where the electrode potential affords control of LNE removal [66, 67]. The process
requires a difference in the metal/metal ion potential of the two metals so that the LNE is soluble
when oxidized and the MNE can diffuse in the structure. Table 1 shows the standard electrode
potential of select metals relative to the standard hydrogen electrode. Typical MNEs and LNEs
are noted, where Ag is commonly used as both a MNE when paired with Al and a LNE element
when paired with Au. In corrosive environments, the atoms in the alloy undergo dissolution
(for the LNE) and surface diffusion (for the MNE), resulting in a monolithic bicontinuous
structure with a uniform, stochastic, 3D sponge-like morphology such as that shown in Figure
5. This present understanding of dealloying was established when Erlebacher et al. reproduced
key experimental observations in a simple atomistic simulation scenario [67-69]. More recently,
liquid metal dealloying has also been employed where one or more components of the starting
alloy selectively dissolve in molten metal [70].

Table 1. Electrochemical series of metals

Metals Half reaction Potential E° (V) Dealloying Role
Gold
Au
3+
+ 3𝑒 −
⇄ Au(s)
+1.53 MNE
Platinum
Pt
2+
+ 2𝑒 −
⇄ Pt(s)
+1.188 MNE
Silver
Ag
+
+ 𝑒 −
⇄ Ag(s)
+0.7996 MNE, LNE
Copper
Cu
3+
+ 2𝑒 −
⇄ Cu(s)
+0.340 MNE
Hydrogen
2H
+
+ 2𝑒 −
⇄ H2 (g)
0
Nickel
Ni
2+
+ 2𝑒 −
⇄ Ni(s)
-0.25 MNE
Zinc
Zn
2+
+ 2𝑒 −
⇄ Zn(s)
-0.7619 LNE
14

Manganese
Mn
2+
+ 2𝑒 −
⇄ Mn(s)
-1.185 LNE
Aluminum
Al
3+
+ 3𝑒 −
⇄ Al(s)
-1.66 LNE

Besides the metal/metal ion potential, other significant parameters in dealloying are the
parting limit and the critical potential for dealloying [71, 72]. For chemical dissolution to
proceed when an alloy is placed in a corrosive environment, the percentage of the LNE and
MNE must be within a parting limit. For example, studies suggest an upper parting limit of
approximately 45 at% Au for Ag-Au alloy [71]. The critical electrode potential, Ec, for
dealloying is composition dependent. For example, in Ag-Au alloys Ec increases with
increasing Au content. The dealloying process can be applied to other alloys for the synthesis
of np-Cu [73], Ni [74], Ag [75], Pd [76], and Pt [77] to name a few, but Au is often preferred
because of its chemical stability.
The resulting length-scale associated with as-dealloyed materials is about 40-50 nm, but
these dimensions can be altered by adjusting the electrochemical dealloying parameters to
reduce the ligament size to ~5 nm [78-80], or coarsening heat treatments to increase the
ligament size up to the micron scale [78]. The scanning electron microscope (SEM) image in
Figure 5a shows a np-Au foam synthesized by free corrosion in nitric acid while Figure 5b
shows the structure produced from the same alloy but using electrochemically driven dealloying
in 1 M HNO3 and 0.01 M of AgNO3. A 2 hour heat treatment at 400 °C can double the ligament
size (Figure 5c). Higher temperatures and longer heat treatments can yield substantially larger
ligaments such as those in Figure 5d after 2 hours at 600 °C.
15

Figure 5: SEM images of np-Au foam prepared by (a) free corrosion and (b) electrochemically
driven corrosion. SEM images of free corrosion samples after 2 hour heat treatments at (c) 400°
C and (d) 600° C. Adapted from [78].

The ease with which one can alter the ligament size is advantageous since it allows for
tuning of properties that depend on size or specific surface area, such as the mechanical strength
[29, 35, 40, 78]. Additionally, dealloyed porous materials can be designed for specific
applications by tuning the ligament size, generating a set of parameters that distinguish these
materials from other porous metals. Currently, np metals have also been studied for applications
in fields such as catalysis [11, 53-55, 81], sensing [11, 56, 57], optical switching [82-84],
electropumping for microfluidic devices [85], integrated circuit contacting [86], biological
implants [87-90], and as electrodes where a large surface area is needed [91-93].
Thus far, the field of np metals has grown to encompass several types of materials and
corresponding applications. For this reason, in Chapter 4, we focus on studies that have
expanded the achievable configurations and applications of np metals by adding structural or
material hierarchy.
16

2.4.3. Architected Materials

Additive manufacturing (AM) or 3D printing techniques have become commercially
available methods for synthesizing arbitrary structures. Generally, structures are first designed
or rendered in modeling software such as SolidWorks, then integrated with the desired printing
technique where they are built layer by layer. Depending on the desired part material, different
processes such as selective laser melting (SLM) for metals or photopolymerization for polymers
are used. A brief overview of techniques and examples of their achievable feature sizes is shown
in Figure 6.
For 3D metal printing, SLM and direct laser metal sintering (DLMS) are two processes
that use a high-powered laser to melt or fuse together metal powders in a layer-by-layer fashion.
Commercially available SLM systems can achieve a minimum feature size of 100-200 microns
and are used in a wide variety of industries including aerospace, where 3D printed parts are
used in rocket designs, or the medical industry to generate custom spinal implants like those
shown in Figure 6a. SLM printable metals include several commercially relevant alloys
including Ti-6Al-4V, SS, Al 2xxx, and Inconel. Methods for qualifying and characterizing SLM
3D printed parts for commercial use is an ongoing and increasingly active area of research.
For finer features, MICA Freeform, a process developed by a company called
Microfabrica, can be used to synthesize complex metal structures with minimum feature sizes
on the order of 10-30 microns (Figure 6b) [94, 95]. These precision materials are made by
repeatedly electrodepositing metal through photomasks, analogous to semiconductor
processing. Each layer is then planarized before the next layer photomask is deposited and the
17

process is repeated. Note that this process is limited to metal systems that can be
electrodeposited and excludes nearly all common alloys.

Figure 6: Additive manufacturing processes to generate complex architectures. From left, a)
SLM generated Ti truss spinal implants [96], b) MICA Freeform made metal structures [97],
and c) 3D-DLW built micro/nano sized epoxy based polymer trusses [98].

Overall, metal additive manufacturing is the fastest growing sector of AM due to the
possible commercial applications of 3D printed parts with alloy materials. Still micro/nano
sized 3D printed materials can only be generated from photopolymerization techniques. An
example of an epoxy based truss structure with feature sizes below one micron is shown in
Figure 6c.
This method of producing polymer structures is called 3D direct laser writing (3D-
DLW) or two-photon lithography via the commercial Nanoscribe Photonic Professional GT
system. Unlike UV wave guides, or even traditional lithography, this technique offers sub 200
nm resolution without masks or substantial optical systems [99]. This machine allows for more
arbitrary, complex, and micro/nano scale structures to be constructed. For example, Figure 7
18

shows a cellular solid composed of a cell geometry that gives the structure a negative Poisson’s
ratio, meaning the structure contracts when compressed [100]. This is just one example of the
many intricate architectures that can be achieved with this printing method.

Figure 7: SEM images of a negative Poisson ratio structure that contracts when compressed.
Structures are produced by two-photon lithography using the Nanoscribe. Adapted from [100].

3D-DLW represents state-of-the-art lithography in terms of feature size and structure
complexity. However, this technique is not capable of generating bulk samples like other 3D
printing methods and the overall process is time consuming. While the resolution provided by
these materials allows for the study of the deformation behavior of architected materials in the
smallest achievable dimensions, scaling up the overall size while maintaining minimum feature
control is the next step for this technology.
Another established, albeit less refined, method for polymer manufacturing is mask-
pattern lithography, which uses UV light exposure guided by a mask to produce a pattern in a
photomonomer that hardens upon exposure. Although mask UV lithography is generally used
for small or flat structures, Jacobsen et al. adapted this process by developing a self-propagating
photopolymer waveguide technique to fabricate three-dimensional polymer truss structures on
19

the order of millimeters (Figure 8) [101]. This process allows for control of the
microarchitecture by altering the cellular geometry to make large truss structures with small
features and known mechanical properties based on the material and cell geometry [102].

Figure 8: (a) Mask UV lithography using a polymer waveguide was to generate a (b) polymer
lattice. Adapted from [101].

2.5. Synthesis of Hierarchical materials
While the fundamental synthesis techniques discussed in the previous section are varied,
their commonalty, apart from the macroscale foaming techniques, lies in their ability to generate
porous cellular materials with a high degree of control over the morphology. In the present
section, an abbreviated overview of processing techniques that build upon fundamental
fabrication methods is given. Both hierarchical foams, using np materials as a base constituent,
and highly architected materials, using templating and additive manufacturing, are discussed.
A more detailed discussion on the recent advances for generating nanoporous materials with
structural hierarchy is given in Chapter 4.

20

2.5.1. Hierarchical foams

The first hierarchical bulk np-Au samples with two distinct levels of porosity were made
using a combination of templating and dealloying. Ag/Au polymer core-shell particles were
first prepared by electroless deposition of Ag and Au on polymer spheres. The spheres were
then heat treated to homogenize the Ag/Au layers and remove the polymer interior. Finally,
dealloying of the Ag-Au alloy shells resulted in hollow np shells [103].
Alternatively, templating can be achieved through the electrodeposition of metal into
unmasked areas. In this manner, nanowire arrays can be made by electroplating an alloy into
the nanopores of a suitable template. The alloy nanowires are then removed from the template,
dispersed on a substrate, and then dealloyed to form np wires [104, 105].
Aside from 2D nanowire arrays, it is possible to fabricate hierarchically structured 2D
np-Au films by electrodepositing a Au-Ag alloy into a 3D template. The final np structure is
obtained by etching the template, followed by free corrosion dealloying of the alloy in nitric
acid [106]. An example of this type of structure is shown in Figure 9.

21

Figure 9: Alloy opal structures (a) are dealloyed in nitric acid to develop a porous opal structure
(b). The overall film structure (c) is approximately 2 µm thick. Adapted from [106].

The first hierarchical structure based on the dealloying of np metals was achieved by
depositing Ag onto a coarsened np Au-Ag leaf. After Ag deposition and subsequent annealing
to form an alloy, another dealloying step is performed, generating the nested bimodal pore
structure in Figure 10a [107]. Another study demonstrated a two-step dealloying protocol for
high LNE alloys that yields a similar nested network in bulk materials (Figure 10b) [108].

22

Figure 10: (a) Nested np structure generated by depositing an alloying element onto a np
metal followed by diffusion and dealloying [107]. (b) Nested np metal by two-step dealloying
of an alloy with abnormally high LNE content [108].

Hierarchical pore structures can also be achieved by exploiting the difference in
electrochemical behavior between elements or phases in a material. For example, a nested np-
Ni structure can be made from a two-step dealloying process, where the precursor alloy was
first dealloyed, the ligaments coarsened via heat treatment, and then a second dealloying step
generates the smaller pore features [104].

2.5.2. Hierarchical Ordered Structures

The use of 3D-printed polymer structures as scaffolds for material deposition, has
facilitated the fabrication of highly controlled and complex architected structures [50]. Etching
of the polymer after coating extends the achievable properties of the overall structure, based on
the intrinsic properties of the base materials [109]. This has led to the development of several
cellular topologies such as hollow metallic microlattices, which are made through an additive
manufacturing process that controls structural elements on multiple length scales [50]. The
levels of achievable hierarchy are summarized in Figure 11, where the overall structure
23

occupies the macroscale, the lattice struts are on the mesoscale or microscale, and the grain size
of the base metal is on the nanoscale. Hierarchically designed materials, such as the lattice,
enable the synthesis of lightweight and strong structures by controlling deformation and
strength throughout the material.

Figure 11: Schematic of the hierarchical scales used in metallic microlattices [19].

These materials are made using the UV waveguide polymer lattices discussed in Section
2.4.3, where the lattices were used in a multi-step synthesis technique to develop hollow
hierarchical metallic structures. The polymer trusses were coated in nanocrystalline Ni through
electroless deposition and the polymer was subsequently etched leaving a hollow metal
structure with even lower density (Figure 12). This particular deposition method generates
conformal coatings. However, there are limitations with regards to material choice, since only
select materials will undergo the desired chemical reactions.
Other similar metamaterials were generated by using stereolithographic printing to
generate octet microlattices as polymer templates and coated with NiP and alumina via
electroless plating and ALD, respectively. For these materials, the polymer was removed via
thermal decomposition.

24

Figure 12: Schematic showing the multi-step templating process used to synthesize hollow
microlattices. (A) UV waveguides were used to fabricate templates. (B) The polymer templates
were coated with nickel and etched resulting in hollow lattices (C). Reproduced from [50].

Templating was also used to fabricate hollow ceramic nanolattices from Nanoscribe
printed templates [110, 111]. These structures were coated with alumina using ALD and the
polymer was removed with plasma etching to generate hollow ceramic structures, shown in
Figure 13a [111]. Atomic layer deposition is an effective technique for coating these materials
with a uniform coating thickness since it relies on two sequential, self-limiting surface reactions
directly on the surface of the template, which means that even the smallest surface features are
captured in the coating. For example, hollow np ceramic structures were made by coating np
Au with alumina via ALD followed by the dissolution of the Au template. However, the process
requires many ALD cycles and is limited to select material systems, which are generally metal
oxides or ceramic materials such as alumina.

25

Figure 13: Hollow lattice structures developed by coating polymer materials generated by
DLW and removal of the polymer material. From left to right, a) ceramic nanolattice [111], b)
Au nanolattice [112], and c) Cu mesolattice [113].

Metal coatings were used to generate hollow lattice structures using base Nanoscribe
structures. Au microlattices [111] (Figure 13b) and Cu mesolattices [113] (Figure 13c) were
made from the electrodeposition of Au and sputtering of Cu, respectively, onto templates
followed by template removal through plasma etching. These two deposition methods result in
variations in the coating thickness throughout the structures since they do not rely on direct
chemical reactions on a surface but rather electrochemistry, which relies on changing ion
concentrations, and physical vapor deposition, which is a momentum driven line-of-sight
process. However, where electrodeposition, like ALD and electroless deposition, is limited by
the materials that can be deposited, sputtering can be used to coat with any number of elements
including alloys. This critical advantage is leveraged later in this dissertation and explained in
further experimental detail in Chapter 3.

2.5.3. Hierarchical Composites

In this section, synthesis methods for incorporating secondary or tertiary materials for
auxiliary hierarchy are reviewed, particularly deposition methods that add an additional layer
a) b)
c)
26

of material onto existing complex structures. Metals have been used as coatings on 3D-DLW
structures, including the electroless deposition of amorphous NiB on a variety of template
architectures that are shown in Figure 14 [114]. The key difference between these materials and
those previously discussed is that they are composites, where the polymer material is retained
instead of removed.

Figure 14: Tilted and top view SEM images of microcellular architectures after NiB
electroless deposition. The scale bar in all the images is 10 µm [114].

Nanoporous composites can be generated by coating np metal ligament surfaces with
additional materials. The challenge of material deposition onto the ligaments of np metals is to
achieve uniform coatings on the fine ligaments. Ding et al. developed a gas-liquid interphase
electroless plating technique to uniformly coat np-Au leaf with Pt [115]. The same gas-liquid
(N2H4 - KMnO4) interphase electroless plating technique was later used by Lang et al. to
fabricate a np-Au leaf / MnO2 composite material [116]. This is an interesting electrode material
for electrochemical supercapacitors, as it combines the high electrical conductivity of np-Au
with the high charge storage capacitance of MnO2.
27

Another deposition method, underpotential deposition (upd), was used to deposit a thin,
conformal layer of Cu onto np-Au wire arrays [117, 118]. The Cu was replaced by Pt in a
spontaneous displacement reaction after submersion in an electrolyte with Pt ions, resulting in
a layered Au/Pt material [119, 120].
Protocols for developing np materials with structural hierarchy due to architecture or
composite materials, including the examples presented in this section (Section 2.5), are
discussed in further detail in Chapter 4.
2.6. Mechanical Behavior of Architected Materials

In general, the properties of architected materials are governed by both the base
materials and the configuration of the materials in space or the architecture. Shaedler et al. have
demonstrated that nanocrystalline nickel microlattices, introduced in section 2.5.2., with the
smallest feature being wall thicknesses of hundreds of nanometers, exhibit dimensional
recovery after 50% compression [50]. Despite the nickel being locally brittle because of its
small grain size, the designed configuration allowed the overall structure to deform with
ductility and recover after compression. The key factor for this recovery behavior was the
optimization of the wall thickness to strut radius ratio [50]. Similar dimensional recovery has
been achieved with hollow ceramic lattices at an even smaller scale with wall thicknesses below
100 nm [121]. In these same ceramic materials, a transition from global ductile behavior to
catastrophic failure was observed when the wall thickness reached a critical value.

28

Figure 15: In-situ X-ray tomography of nickel microlattice under compression. Note that the
deformation of the structure is concentrated at the nodes. Adapted from [122].

The main cause of failure for these materials is high stress concentration at the nodes,
shown in Figure 15, which leads to node cracking and buckling [122]. This failure is expected
since the nodes are only supported by hollow strut walls. To realize the full potential strength
of these ultralight materials, the critical factor is improving nodal strength. Numerical studies
indicate that significant strengthening of the overall structure can be achieved by modifying the
thickness of the cell walls or strengthening them [122]. Additional levels of structural hierarchy
to control deformation behavior may be the key to solving the nodal issue and implementing a
new parameter which may improve nodal strength. Thus, an opportunity remains to increase
strength by introducing new morphological features in these materials.
Nodal failure is similarly an issue for solid lattice structures including 3D printed
microlattice materials. Aside from as printed IP-Dip structures, these materials have a
substantial amount of nodal failure and are plagued by brittle fracture. For example,
29

unprocessed IP dip tetrahedral trusses fail by buckling in the ligaments while heat treated
samples exhibit both increased strength and a transition to brittle failure at the nodes [98].

2.7. Applications of Hierarchical Materials
Many promising opportunities for using highly ordered structures as functional
materials rely on exploiting the large specific surface area in the pore space of open cells. In
most instances the function relies on the transport of electric signals in the form of ions in a
fluid through the open pore space. Non-hierarchical nanoporosity, however, restricts mass
transport and thus response time [25]. It has been pointed out that many of the potential
applications, such as catalysts or batteries, entail conflicting requirements on the material
structure [123, 124]. Therefore, while smaller features translate into higher surface area and an
increased number of active sites for the desired function, they are detrimental to fluid transport.
Structural hierarchy is a promising approach towards reconciling these conflicting
requirements, where a simple hierarchical structure contains small pores for function and large
pores for transport. In this regard, the use of additively manufactured structures as base
materials would be beneficial since they have known geometry and thus more predictable mass
transport properties than stochastic hierarchical materials.
Further details on the fabrication of hierarchical materials based on nanoporous foams
can be found in Chapter 4. The following brief sections detail specific applications suited to np
metals with structural hierarchy and, separately, architected materials generated by additive
manufacturing or 3D printing methods.

30

2.7.1. Applications for hierarchical np metals
Catalysis
To date, the most studied and viable application of np metals is catalysis. Raney nickel,
first developed in 1926, is used for a variety of industrially important hydrogenation reactions,
and Raney copper is a widely-used catalyst for the water-gas shift reaction [125]. Other
examples of catalytic reactions possible with np metals include the electrooxidation of methanol
and ethanol [126-128], as well as the electroreduction of oxygen [129, 130], and carbon
monoxide [53, 131].
One notable catalytic study investigated hierarchical np-Pd, which has proven to be an
effective catalyst for alcohol oxidation [126, 132]. Nested bimodal np-Pd exhibited improved
reaction kinetics for ethanol oxidation when compared to a Pd plate electrode and
nonhierarchical np-Pd [126]. In the same study, two bimodal structures, one with larger pores
than the other, were compared for catalytic activity. Interestingly, the nested structure with
larger pore sizes, but with lower surface area, showed a higher reactivity for ethanol oxidation.
As pointed out before, hierarchical pore morphologies are ideally suited for applications such
as catalysis that benefit from the combination of high surface area (provided by nanopores) and
fast mass transport (provided by macropores).
Actuation
Due to the large surface area of np metals, the ligaments experience a large elastic strain,
which can be controlled through externally imposed electric [58, 133] or chemical [134]
potentials. This phenomenon functionalizes np metals so that that they become electrochemical
actuators. The strain amplitude may be amplified in microstructures with multiple length-
scales. Maximum strain amplitudes of up to 6% have been reported for layer structures of np-
31

Au [135]. Shi et al. have shown that nested hierarchical np-Pd, prepared by dealloying, is highly
deformable in compression without failure [76]. It exhibits hydrogen-enhanced actuation with
amplitudes up to 4% and stable actuation over > 1000 cycles.
Energy Systems
Recently, researchers developed a hierarchical np-Au electrode for use in a Li-O2
battery system. Compared to non-hierarchical as-prepared np metals, the hierarchical structure
increased the reversible capacity and led to longer cycling lifetimes at low overpotentials [136].

Other studies include Sn coated np metals for lithium battery electrodes [137, 138].

Oxides such as MnO2 [139] and TiO2 [140] have also been used to develop composite
np structures. These particular materials are effective electrodes due to the symbiotic
relationship between the np metal scaffold, which is a current collector, and the oxide coating,
which can act to stabilize the structure and store Li [140].

In addition to battery oriented studies,
np metal based supercapacitors have been generated as energy storage devices [116, 141, 142].

Sensors
Np metals are excellent substrates for surface enhanced Raman scattering (SERS)
sensing due to nanoscale surface roughness. Using added structural features, this sensing can
be enhanced. For example, Qian et al. decorated np-Au with Au nanoparticles by infiltrating
the nanopores, which significantly enhanced the Raman scattering when compared to np-Au
[143]. Zhang et al. demonstrated improved Raman scattering for wrinkled np-Au films, which
are produced by the thermal shrinkage of the film substrate [144]. Additionally, hierarchical
porous gold with aligned arrays of ~100 nm holes and evenly distributed mesopores exhibited
superior SERS activities compared to commercial materials [145].
32

Although not extensively studied, other hierarchical np materials have also been
developed for alternative sensing applications. One such study used np-Au nanowires for
sensing the adsorption of octadecanthiol [146]. Np-Au films mounted on commercial Ni foams
were shown to detect hydrazine gas and hydrogen peroxide [147]. Finally, bimodal np-Pt-Ti
have been explored for the sensing of ascorbic acid, dopamine, and uric acid [148].

2.7.2. Applications for architected materials

Structural materials
Incorporating additional length scales or architectures can significantly improve the
performance of materials. A prime example is the design of sandwich panels and their core
structure to improve mechanical properties [16-18]. Core architectures and materials vary
widely, however architected materials can provide novel designs compared to traditional foam
core or honeycombs. Hollow truss structures, in particular, have been studied as alternatives to
honeycombs [149].

The variety of available core designs and numerous material options allows
for the optimization of these metal cellular structures for applications in the aerospace and
automotive industry [16, 19].
Other traditional uses for cellular materials include enhanced impact or energy
absorption [150]. Microcellular materials [151] and microcellular [152] materials with
controlled periodic architecture have been explored in this capacity.
Thermal materials
In addition to structural optimization, sandwich panel structures with designed cores can
be tailored for thermal cooling when a fluid flows through an open-celled panel [26, 27]. A high
degree of control over the structural features and the availability of different build materials,
33

such as that afforded by additive manufacturing techniques, allows one to optimize both
mechanical properties such as stiffness and thermal properties such as heat transfer rates [19,
153].
Furthermore, thermal devices such as heat pipes [154] and heat exchangers [16, 155]
have been developed using lattice materials. One interesting study leverages hollow lattices to
develop dual fluid paths through the interior and exterior of the structure [156].
Photonics and Mechanical Phenomena
The earliest and most studied application for micro or nano cellular polymer materials
produced via 3D-DLW are for photonic applications [157, 158]. The precision and scale of
reproducible ordered structured materials affords control over the nanofeatured periodic
structures necessary to generate unique optical dispersion properties [159, 160].
The same precision that makes 3D-DLW ideal for photonic research is directly
beneficial to those exploring mechanical phenomena with micro/nano cellular materials.
Several auxetic materials, which have a negative Poisson’s ratio (i.e. they contract when
compressed), have been developed [100, 161, 162]. Additionally, pentamode materials, which
were previously theoretical materials that behave mechanically like a fluid (i.e. has nearly zero
shear modulus), were developed for the first time using 3D-DLW [163].

34

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35

Chapter 3 : Experimental Methods and Materials
This chapter presents synthesis methods and techniques used for the processing of
hierarchical materials with either multimodal porosity or highly controlled cellular architecture.
First, metal layer deposition methods, electrodeposition and magnetron sputtering, are
explained in detail relevant to their role in this study. Additionally, characterization tools to
study the foam morphology, cellular architecture, chemical composition, and film properties
are covered. This section concludes with the methods and materials used for the mechanical
characterization of microcellular metal/polymer composites and scaled up polymer cellular
materials.

3.1. Electrodeposition
Electrodeposition is a coating technique that uses an applied electrical current to reduce
metal ions in an electrolyte onto an electrode, forming a metal deposit. Figure 16 shows a
simplified electroplating system for the deposition of silver from an aqueous silver nitrate
solution (AgNO3) containing positively charged Ag cations. When an external current is
applied, Ag ions migrate toward the cathode and are reduced to metallic Ag at the interface
between the electrolyte and the cathode. At the anode, Ag is dissolved into the solution to
maintain bath equilibrium and replace the Ag that was removed from the solution.
36

Figure 16: Electrolytic cell for the deposition of silver from an aqueous silver nitrate solution.
Silver ions are reduced on the surface of a cathode as a metal film.

For the purposes of this study, Ag was electrodeposited onto tubular Au substrates using
a silver nitrate based commercial plating solution (Caswell Inc). Recommended bath
preparation and maintenance were followed throughout the study and use of the solutions. As
mentioned in Chapter 2, electrodeposition does not necessarily yield uniform coatings due to
both ion concentration variations near the deposition surface. However, the anode and cathode
positioning plays a role in coating uniformity. Consider the idealized electric field generated
between two anodes in Figure 17a. This parallel field is not feasible and, in reality, the electric
field lines have a higher concentration near the electrode edges and corners like in Figure 17b,
which leads to higher localized current densities in these regions.
37

Figure 17: (a) Schematic of an idealized electric field connecting two electrodes in an
electrolyte solutions. (b) Shows the actual electric field for two parallel electrodes with field
line concentrations near the edges of the materials. In practice, the areas with higher electric
field line concentration are regions with a higher localized current density. Adapted from [164].

Due to the cylindrical shape of the selected substrate, a custom electrochemical setup
was used to ensure annular coating of the tube structures. This was accomplished using a
cylindrical Ag anode surrounding the substrate so that any point on the outer surface of tube
has a similar distance to an anode surface. A schematic of the setup is shown in Figure 18a.
Note that the cathode is placed in the center with a slight opening in the anode to encourage
fluid flow. The fluid is both agitated and cleaned by a stir rod and a fish tank filter, with the
outlet aimed directly at the anode to support ion replenishment. This electrodeposition setup is
placed in a 1 L beaker with approximately 1 L of Caswell Silver Plating solution
a)
b)
38

Figure 18: Schematics showing the experimental electrodeposition setup. The relative
placement of the electrodes are shown in (a). (b) shows an image of the entire system
including the anode, cathode, reference electrode, and filter.

Figure 19: Schematic of an updated experimental electrodeposition setup with two
semicircular anodes, the cathode, and a filter.

a
b
39

An updated electrodeposition setup was developed using two semicircular Ag anodes
placed in a cylindrical configuration. To further standardize the electrochemical setup, a custom
mounting system, with attachment areas for the wiring and electrodes, was designed in
SolidWorks and generated with a commercial 3D printing system using ABS as the printing
material. The resulting structure is shown in Figure 20a with the Au and Ag electrodes attached
to the mounting piece in the intended regions and electrically connected with alligator clips
(green for the working electrode and white for the counter electrode). The entire setup is
designed to be inserted into a 2 L beaker with 1 L of Caswell Silver Plating Solution, shown in
Figure 20b.

Figure 20: (a) Electrode mounting piece connected to the Au and Ag electrodes out of the
electrolyte. (b) The entire system placed in a 2 L beaker with 1 L of Caswell Ag plating
solution.

Conductive silver paint (PELCO, Ted Pella) was used to make an electrical contact
between the Au tube and shielded Cu wire. The joint was electrically masked by with fast curing
epoxy (Loctite Pro), leaving 1 cm length of Au tube exposed (Figure 21). A small dab of epoxy
was used to mask the bottom of the tube to eliminate fluid access to the cross-sectional and
a b
40

inner tube surface area. Improper masking of cathode materials leads to an unknown exposed
surface area during plating which can inadvertently alter the applied current density. The wire
surfaces were cleaned in a 17.5 vol% sulfuric acid solution for 1 min, then thoroughly rinsed in
DI water prior to plating. The same acid dip was applied after the wire was plated with Ag. The
mass of the wires was measured before attachment to the Cu wire, with the wire/epoxy setup,
after plating while still attached to the Cu wire, and one final time with the total length of the
coated tube measuring 1 cm. By tracking the mass gain of the structure, samples with the desired
amount of coating were identified.

Figure 21: (Top) Au tube connected to a Cu wire with Ag pain as solder and epoxy to mask
the joint. (Bottom) Au tube coated with a layer of Ag.

The electrodeposition process was driven by a three-electrode electrochemical cell
controlled by a potentiostat (Gamry Reference 3000). The Au tubes were used as working
electrodes, Ag plates as the counter electrodes, and a Ag/AgCl reference electrode. All
41

deposition experiments were run with VFP600 software at a direct or pulsed current determined
by the chosen current density.

3.2. Heat Treatments
After electrodeposition, the bilayer tubes were heat treated in an argon environment for
24 hours at 900°C to diffuse the Ag and Au to form fully homogenized alloy tubes. These heat
treatments were carried out in a MTI-GSL1100X tube furnace shown in Figure 22a, where Ar
gas flowed through the furnace. The gas flow passes through an oil to prevent air from
backflowing into the furnace. Coarsening treatments for prepared np Au materials was
completed in the Vulcan 3-550 (Figure 22B) furnace with no controlled gas environment since
Au is inherently inert with no native oxide.

Figure 22: (A) MTI-GSL1100X tube furnace used for homogenizations heat treatments. (B)
Vulcan 3-550 furnace used for np Au ligament coarsening treatments.

3.3. Magnetron Sputtering
Magnetron sputtering is a physical vapor deposition (PVD) process that uses ionized
gas to bombard a target and eject atoms from the target to form a metal vapor, which will coat
any material exposed to the atoms. Figure 23 shows a schematic representation of the sputtering
set-up and process. First, a neutral gas, typically Ar, is introduced into a vacuum chamber and
42

a bias voltage is applied to the target causing the gas to be ionized and form plasma. The gas
ions are attracted to the negatively biased target and strike it. The momentum and impact of
the ions results in the ejection of atoms from the target surface, which travel across the vacuum
chamber and deposit onto the substrate.

Figure 23: Simplified schematic showing the set-up and process of magnetron sputtering: a
neutral gas is introduced and ionized by the negatively biased target; the gas ions strike the
target and cause the target atoms to be ejected then coat the substrate [165].

Typical magnetron sputtering chambers have multiple magnetron sources mounted
inside a high vacuum chamber with achievable pressures in the range of 10
-5
-10
-7
Torr. A model
of the specific sputtering chamber used in this study is shown in Figure 24. Before sputtering,
the chamber is pumped down for about 24 hours and high purity Argon (99.999%) pumped into
the chamber and used as the ionized gas. The shown configuration allows for multiple
sputtering sources in a variety of sizes to be mounted on each of the available source locations.
Ar ions
(plasma)
Negatively
biased
target
Ejected
target
atoms
43

A noteworthy attachment on this experimental set up is the installed DC motor on the chamber
door. This motor drives a rotating substrate stage used for substrate rotation during sputtering.

Figure 24: Schematic of the DC magnetron sputtering chambers used in this study. Multiple
sputtering sources can be used simultaneously in this configuration. The attached DC motor
can be used to rotate substrates during coating or for direct line-of-sight to a specific target.
Reproduced from [166].

Sputtering targets are attached to the source electrode using a magnetic keeper mounted
on the back of the targets. The targets used in this study are shown in Figure 25, where the
materials being used are pure Al, Inconel 600, and Ti-6Al-4V. Of these three targets, it is
apparent from the deeper ring or “race track” that the Inconel 600 target has been sputtered the
most. The trench is formed in regions where the ionization of the Ar gas and bombardment of
the target surface is at its maximum.
44

Figure 25: Image showing three targets used for magnetron sputtering in this study. From
the image and entrenched “race track” it is apparent that the Inconel 600 target has been
sputtered more than the Al or Ti-6Al-4V targets.

This particular PVD method is limited by “line-of-sight” from substrate to target, since
it is a physical momentum driven process. Therefore, areas that are shadowed from the source
material are not coated effectively and a variation of deposition layer thickness is generated.
For this reason, sputtering is most effective for the synthesis of films. However, 3D structures
can be coated by rotating or exposing additional areas of the structure to the metal vapor during
sputtering. This is accomplished by using alternative substrate holders and chamber
configurations designed with some consideration to minimize shadowing effects on the
substrate, dissipating the heat caused by the sputtering process, and the ability to move the
sample within the chamber. These criteria are crucial to the sputtering parameters in this study,
namely because it involves coating complex polymer structures that are susceptible to both
shadowing and melting. Furthermore, the complex morphologies require in-situ manipulation
of the sample to achieve coatings in otherwise obscure regions of the samples.
45

The two chamber configurations and substrate holders used in this study are shown in
Figure 26. A typical stationary setup includes a three-sample holder placed directly parallel to
the sputtering target, the specific holder used in this study is characterized by three separate
sample regions. Alternatively, a rotating configuration can be used by mounting a rotating
substrate holder onto the DC motor shown on the sputtering chamber illustrated in Figure 26.
Samples were mounted onto the substrate holders with vacuum tape. A substrate to target
distance of 10.2 cm was used for both chamber configurations. The rotating holder was rotated
at the maximum motor speed of 54 rotations per minute.

Figure 26: a) Shows the representative sputtering schematic for stationary sputtering. b)
Shows schematic for rotating sputtering set up with samples mounted in the center of the
rotating holder. Sample is rotated 54 times per minute.

b)
a)
46

The following general sputtering parameters were followed for each coated sample in
this study: After the desired base pressure was reached, the target is first cooled to
approximately 16º C. The desired Ar pressure (4 mTorr) was then set using a closed-loop
pressure controller. At this point in the procedure, the substrates are separated from the targets
by a physical shutter. The shutter is necessary to protect the substrates during pre-sputtering,
where the target is initially sputtered to remove any contaminates that are on its surface. The
power supply was set to the desired power (65 W) and the source was turned on to generate the
plasma and begin sputtering. Pre-sputtering was performed for about 90 s before the shutter
was opened to begin deposition on the substrates for the desired time. All sputtering rates were
predetermined by sputtering thin films on glass or Si wafers and measuring the thickness
divided by the sputtering time.
The main advantages of using sputtering as a coating method are the vast number of
metals that can be sputtered, including light elements like Al, refractory metals like Ta, and
even alloys. This affords engineers the freedom to explore multiple materials for their coating
properties.
3.4. 3D Direct Laser Writing
3D Direct laser writing (3D-DLW) was used to fabricate the microcellular structures
studied in this dissertation. This 3D printing process relies on two photon polymerization,
achieved by tightly focusing a laser inside a photosensitive polymer resist so that the laser
intensity is sufficiently high enough to initiate polymerization in this localized area or voxel.
The voxel size can be scaled using laser power or exposure time. 3D structures are achieved by
tracing out the desired architecture with the focused laser area. A schematic of the DLW process
is shown in Figure 27a, where the Nanoscribe software controls both the laser parameters and
47

the piezoelectric scanning stage. Figure 27b shows an illustration of the focused laser
interaction with the resist.

Figure 27: a) An image of the Nanoscribe GmbH Photonic Professional GT laser instrument
Reproduced from [167]. b) Illustrates laser beam interaction with resist to produce 3D
structures.

The structures used in this study were printed using a Nanoscribe GmbH Photonic
Professional GT laser lithography system with Nanoscribe’s epoxy based proprietary resin, IP-
Dip. The general preparation procedures are as follows: A drop of IP-Dip is placed on a glass
slide, which is then mounted up-side down on the piezoelectric stage. Structures are printed on
the surface of the inverted glass slide. The slide is then placed in a developing solution and later
submerged in isopropanol and separately acetone. After being placed in acetone, the slide is
then subjected to super critical CO2 drying.
3.5. Characterization Techniques
3.5.1. Scanning Electron Microscopy (SEM) and Electron Dispersive Spectroscopy (EDS)
Scanning electron microscopy (SEM) is the main imaging technique used in this study.
Images and data are generated from a focused electron beam that interacts with the surface of a
sample to generate secondary electrons, backscattered electrons, and characteristic X-rays.
a) b)
48

These signals are collected by detectors and used to examine topology, crystal orientation, and
elemental composition of a target specimen, respectively. This interaction is shown
schematically in Figure 28. All SEM images and EDS measurements in this study were
conducted on a JSM-7001F-LV field-emission SEM at the Center for Electron Microscopy and
Microanalysis at USC unless otherwise noted.

Figure 28: Schematic of the electron beam and specimen atom interaction volume under the
specimen surface. Reproduced from [168].
The energy of secondary electrons, used for imaging, is relatively small and are
therefore only emitted from the surface of the specimen. However, the characteristic X-rays
from a much deeper interaction volume are useful for another crucial characterization tool,
energy dispersive spectroscopy (EDS), which uses characteristic X-rays to determine the
elemental composition of the sample. Elemental analysis was conducted using EDAX Element
49

Silicon Drift Detector for characteristic X-ray detection and EDAX Genesis software for data
processing and elemental determination.
This technique was used in several capacities in this work. The initial function was to
determine the alloy composition of diffused bi-layer metal tubes and to determine the amount
of residual LNE after the alloy was dealloyed. EDS was used for the sputtering studies to
determine if alloy stoichiometry was maintained when sputtering alloys, particularly
commercial alloys such as Inconel 600 and Ti-6Al-4V.
3.5.2. Focused Ion Beam (FIB)

A JEOL JIB-4500 FIB was used in this study to explore the cross sections of micron
thick sputter deposited films as well as for the FIB milling of coated cellular structures to expose
their cross section. Unlike the SEM, the FIB system uses a beam of Ga
+
gas ions to bombard
the samples. A detector measures the intensity of the secondary electrons that are emitted due
to the bombardment. Since Ga
+
ions are significantly heavier than the electrons used in SEM
imaging the ions penetrate deeper into the material with a higher energy. FIB imaging produces
better channeling contrast, compared to SEM, allowing for different grain orientations to be
resolved. However, the higher energy and heavier gas ions can cause significant sample
damage.
In addition to imaging, a higher energy FIB beam can be used to mill or cut samples in
a manner analogous to sputtering, where ionized gas ions are used to erode a material. Using
the system controls, micromilling can be used to mill trenches in films so that their cross
sections can be exposed and evaluated. This method was used here in an attempt to evaluate the
cross sections of microlattice structures, however due to the eroding process of milling, a
50

substantial amount of material that was milled away was then redeposited inside the cellular
structure, thus obscuring the ability to make coating thickness measurements on these materials.
Due to redeposition caused by FIB milling, microtome was explored as an alternative
method to cross-section the microstructures. In microtome, samples are embedded in an epoxy
resin, heat polymerized, and sliced into thin (<100 nm) sections with a diamond knife. The
sliced samples were floated on water and picked up with a formvar-filmed slot grid for TEM
imaging.

3.5.3. X-Ray Diffraction (XRD)

XRD was used to explore the texture and microstructure of metal films generated by
magnetron sputtering. These measurements were completed in a Rigaku Ultima-IV X-Ray
diffraction machine. XRD was performed using Cu Kα radiation with the wavelength of 1.5 Å.
A series of 2θ scans from 30° to 100° with the rate of 3°/min were performed.

3.6. Mechanical Testing of Microcellular Materials
Mechanical testing of polymer and polymer/metal composite cellular structures was
completed using two micromechanical instruments. For the tensile testing of the polymer
structures, a custom set-up on a modular micromechanical testing machine, shown in Figure
29, was used. This instrument is similar to the machine used by Eberl et al. and Balk et al. for
the tensile testing of thin films with millimeter dimensions [169-171]. The machine is made of
a coarse actuator (ThorLabs DRV001), piezo actuator (ThorLabs PAZ015), two-dimensional
alignment stage, load cell, safety springs, and a camera. The mechanical testing stages and data
acquisition is controlled using a custom MATLAB script, which allows user regulation of the
51

actuators and automated testing protocols that measures and saves load, time and displacement
data. In addition to tracking the displacement and load, the set up includes a camera to track
points on the surface of the sample as it deforms. Using digital image correlation (DIC) software
obtained from Prof. Dr. Christoph Eberl at the University of Freiburg, stress-strain curves and
strain maps can be generated from tracked points. However, the sample dimensions used in this
study are too small for useful resolution.
Test stages on this instrument are interchangeable, with one side stationary and the other
side movable in the forward or reverse direction for compression or tensile loading. The
modular stages can be designed for specific testing configurations including tensile or
compression tests, and three or four-point bending.

Figure 29: Photo of micromechanical test set-up with interchangeable testing stages and
attached camera for digital image correlation to track sample deformation.
52

To evaluate the compressive response of polymer/metal composite lattice structures,
uniaxial compression tests were performed using an Agilent Technologies G200 XP
nanoindenter system, like that shown in Figure 30a. Typically, nanoindenters employ
Berkovick or Vicker’s indenter tips to evaluate the hardness of a wide range of materials
including metals, geological samples, biological materials, or even the np metals previously
introduced. Three main types of loading scenarios are possible with this instrument, which are
feedback control, constant load, or displacement control. These instruments can also perform
compression tests on cellular structures by using a flat punch indenter tip to uniaxially compress
the samples. In this work, compression of tetrahedral lattice truss structures was carried out
with a diamond flat punch (100 or 200 µm diameter) at a constant displacement rate of 100
nm/s, where the sample was loaded until failure. The structures in this study were not loaded
until densification, but to the onset of the sample fracture. A schematic of the test set-up is
shown in Figure 30b.

Figure 30: a) Photo of Agilent G200 XP nanoindenter used to conduct uniaxial compression
tests of microtruss structures [172]. b) Schematic of a compression test using a diamond flat
punch tip moving at a constant displacement rate.

The stress and strain of the samples were analyzed using the hexagonal cross-sectional
area of the trusses and the height of the sample, respectively. The compressive strength of the
b) a)
53

samples is designated as the maximum sustained stress at failure and the stiffness was calculated
from the elastic regime of the stress-strain curves. A typical stress strain curve for these
compressive tests is shown in Figure 31, with all major regions of the curve labeled.

Figure 31: Representative stress-strain curve of tetrahedral truss structures. The stiffness is
calculated from the elastic regime, while the compressive strength is the maximum stress.

54

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55

Chapter 4 : Review of Nanoporous Materials with Structural Hierarchy

The work presented in this chapter can be found in the publication titled Nanoporous
Metals with Structural Hierarchy: A Review in the journal Advanced Engineering Materials.

Materials with multiple levels of porosity or hierarchal elements can be achieved by
numerous synthesis methods, including the use of 3D printed polymer structures as scaffolds
for material deposition, generating more complex architected structures [50]. Etching (or
retention) of the polymer after coating extends the achievable properties of the overall structure,
based on the intrinsic properties of the base materials [109]. In a similar manner, using np
materials as a base constituent has its own advantages, namely the scalability and the
enhancement in surface area. Two main methods (dealloying and templating) for achieving np
materials with multimodal porosity are presented in this chapter. Specifically, we consider
materials in which the smallest and definitive length scale is occupied by np metals made by
dealloying. Examples of as-dealloyed np topologies are shown at the center of Figure 32.
Alternative processing methods that can produce additional levels of tunable structural features
in single-phase materials (left of Figure 32) or in composites (right of Figure 32) are also
reviewed.

56

Figure 32: Schematic representing various hierarchical structures in single element np
materials [108, 135, 173-175] (A-F) and composite structures [116, 140, 176-178] (G-K) that
are based on np metal morphologies [73, 75, 78, 179] (L-O) produced by dealloying.

4.1. Hierarchical np materials produced through templating

Nyce et al. developed the first 3D hierarchical bulk np-Au samples with two distinct
length scales of porosity using a combination of templating and dealloying [103]. Ag/Au
polystyrene (PS) core-shell particles were first prepared by electroless deposition of Ag and Au
on PS spheres followed by casting to obtain a monolithic 3D bulk sample, which was then heat
treated to remove the PS template while simultaneously alloying the Ag/Au layers. Finally,
57

dealloying transformed the hollow Ag-Au alloy shells of the structure into hollow np shells
showing the typical bicontinuous morphology of np Au. The 200 nm thick shell wall survived
dealloying without crack formation despite the low pre-dealloying Au content of only 15 at.%.
In general, hierarchical structures [108] and low-dimensional nanoscale objects [146] seem to
be more resistant to stress induced cracking than uniform bulk samples that tend to disintegrate
during dealloying of low Au content alloys.
Alternatively, templating can be achieved through electrodeposition of metal into
predetermined areas defined by masking. In this sense, Ag-Au nanowires can be made by
electroplating the alloy into the nanopores of a suitable template. The alloy nanowires are
stripped from the template and dispersed on a substrate, then dealloyed to form np wires (Figure
33a) [104, 105]. The same idea was applied to develop arrays of nanotubes with np walls, as
shown in Figure 33b, and porous Pt-Co nanowires [117, 180]. Still, the fabrication of these rods
through mesh templating limits the length of the nanowires to the thickness of the resist used
to mask the plating channels. The length of the potential nanoarrays was increased by a method
for the electrodeposition of Au-Cu onto a patterned silicon substrate. The substrate was made
conductive by a vapor-deposited Au seed layer [181]. The resulting patterned structure consists
of one-dimensional lines with np structure.
58

Figure 33: Various nanowire structures ranging from short np wires (a) to np nanotubes (b) to
long nanowires fabricated on silicon patterned substrates (c & d). Reproduced from ref [104,
117, 182].

More recent work has demonstrated centimeter long np nanowires developed by first
sputtering Au-Cu on a rotating silicon substrate with 120 nm wide nano-patterned grates and a
length of a few centimeters. Rotating the structure with the grates perpendicular to the target
throughout sputtering develops the circular cross section of the rod. The sputtering was
followed by electrochemical dealloying of the patterned nanowires resulting in the structures
shown in Figure 33c and 33d [182].
Aside from 2D nanowire arrays, it is also possible to fabricate hierarchically structured
2D np-Au opal films by electrodepositing a Au-Ag alloy into a silica inverse-opal template.
The final np-Au opal structure (Figure 9) is obtained by removing the silica template with an
HF etch, followed by free corrosion dealloying in nitric acid [106]. The overall structure of the
hierarchical film is displayed in Figure 9c; the smallest pore spacing inside the opals was
approximately 10 nm. The opals themselves were approximately 400 nm in diameter with an
59

interstitial spacing of 40-100 nm, where the larger gaps in the structure are a result of opal
shrinkage during dealloying. Overall, the film is about 2 μm thick and attached to a substrate.
Thus far, this type of material, like the shorter nanowires, has not been synthesized in bulk
quantities.
Templating and dealloying are methods that can be used in combination to incorporate
porosity and control of the overall structure of the final material, such as the 1D nanowires, 2D
opal structures and the 3D bulk materials discussed in this section. The use of templates,
however, increases the number of synthesis steps required to produce these intricate structures.
While templating is necessary when a specific shape is desired, multiple length scales can also
be implemented by other methods relying on electrochemistry or the initial microstructure of
the precursor alloy, which will be discussed in the next section.
4.2. Dealloying methods to produce hierarchical np metals

The earliest example of a hierarchical structure based on np metals, with distinct length
scales of self-similar porosity, is from Ding and Erlebacher [107]. The ligaments of this material
are composed of smaller ligaments with the same characteristic shape. This architecture was
realized by depositing Ag onto a previously dealloyed and subsequently annealed 1 micron
thick Au-Ag leaf, with ligament diameters in the range of 100 nm. After Ag deposition and
annealing to form an alloy, another dealloying step is performed, resulting in formation of the
bimodal pore size structure, which was first shown in Figure 10a [107]. The uniform deposition
of Ag inside the np-Au structure was achieved by the development of a new gas-liquid interface
electroless plating technique that takes advantage of the tendency of np-Au leafs to float on
aqueous solutions and therefore requires the film to be thin.
60

While the above concept is limited to the preparation of hierarchical structures in the
form of thin films, Qi and Weissmüller demonstrated a dealloying protocol that yields self-
similar nested networks on different length scales in the form of three-dimensional bodies [108].
They developed a method to electrochemically dealloy a high LNE alloy (Ag95Au5), while
retaining enough Ag content to approach the parting limit of the Au-Ag alloy and produce the
characteristic np-Au ligament morphology. Coarsening of the initial dealloyed materials
increased the ligament size from ~16 nm to ~200 nm, similar to the case of np-Au. The key
finding in this study is that the characteristic porous structure can be achieved without fully
leaching out the Ag, allowing for annealing and second dealloying step, producing a nested
structure with both a higher and lower level of porosity. This sequential dealloying process was
later applied to porous sheets of a AgAu alloy to generate tri-modal porous structures [183].
Qiu et al. also exploited the difference in electrochemical behavior between elements in
a study, where the difference in corrosive behaviors between Cu and Mn can be used to develop
a nested Ni pore structure from a Mn-Cu-Ni alloy [141]. Their study followed the same steps
as Qi and Weissmüller, where the precursor alloy was first dealloyed, the ligaments coarsened,
and then a final dealloying achieved the second level of porosity. However, the corrosive
environments were varied in each dealloying step so that only Mn was removed from the alloy
in the first step and Cu was leached in the second.
There are many publications where dealloying is applied to master alloy samples which
start out with a multiphase microstructure that is created during solidification. The dealloying
then works differently on the separate phases, and the coarser microstructure from solidification
later coexists with the finer microstructure created by dealloying. When carefully designed and
61

controlled, solidification strategies can result in highly ordered structures with well-defined
structural length scales.
For example, hierarchical np-Pd has been made by dealloying Pd-Al alloys produced by
spark plasma sintering. A bimodal porous structure is developed in regions where the
microstructure is heterogeneous due to the different chemical activities for Al and PdAl3 phases
[126]. In regions where both phases are present, a bimodal structure is formed, whereas in
regions with only PdAl3, a monolithic porous structure is formed. The regions rich in Pd produce
ligaments as small as 7 nm, and the regions rich in Al produce larger pores (~200 nm) with
embedded smaller pores due to the PdAl3 phase [126]. Note that doping a precursor alloy with
a more noble and low mobility element, such as Pt or Pd, can be used reduce the size of the np
ligaments in single or bimodal materials because of the slower surface diffusion of the more
noble element [79, 124, 184, 185]. Hierarchical np-Pd has also been generated by the
electrochemical dealloying of a Pd-Ni-P metallic glass [186].
An early study that showed clear evidence for the benefits of using two-phase starting
alloys was by Jin et al. [187]. In their study, the solidification of a Pt-Ag master alloy led to Pt-
rich dendrites, which resisted corrosion during later dealloying. The dendrites that remained
embedded in the nanoporous matrix acted as efficient crack stoppers, substantially enhancing
the toughness.
In addition to the studies discussed in detail above, several Al alloys, where the Al is
removed during dealloying, have been exploited to generate bimodal materials. These include
multiphase alloys, including Ni-Al [188], Al-Cu [189, 190], and Al-Au [175, 191], as well as
ternary alloys such as Al-Pd-Au [192] and Pt-Ti-Al [193] that generate binary alloy np
materials. In the two latter studies, Al was first dealloyed in NaOH or NaOH/HCl and the more
62

noble elements were removed with nitric acid. Further studies, using multiphase or multi-
element alloys for dealloying multimodal porosities can be found in the following references
[127, 194-196].
Hierarchical foams can also be made by dealloying preformed, multiphase grain
structures within percolation thresholds. Cox and Dunand formed macro/microporous Au-Ag
foam by mixing salt into a metal powder pack before compression and sintering [197]. The salt
was dissolved in water, which produced large block-shaped macropores, while micropores were
developed from the gas expansion of water vapor trapped in the powder compact. This structure
is shown in Figure 34. The sintering process alloyed the Au and Ag powders, resulting in a
Au19Ag81 at% alloy with graded porosity. The final tier of structure, nanoporosity, was added
by dealloying in nitric acid.

Figure 34: Cross-section of alloy foam before dealloying. (a) Large block shaped macropores
from salt removal. (b) Enlarged area showing micropores from gas expansion. Reproduced
from [197].

Preexisting commercial macroscopic foams can also be manipulated to exhibit
nanoporosity. Roughening the surface of commercial Cu foams using electrodeposition and
63

electrochemical dealloying develops a nanoporous layer on the foam surface [174]. This is
accomplished using a single aqueous electrolyte solution, where a Ni-Cu alloy is deposited onto
a macroporous Cu foam and the Cu is selectively etched during dealloying from both the alloy
coating and the pure Cu cell walls underneath. Undercutting of the Ni-Cu film results in the
removal of the Ni components, leaving a roughened np layer of Cu on the surface of the walls.
Both the number of deposition/dealloying cycles, as well as the time duration for the dealloying
pulse can control roughness [174].
In other studies, a combined Gasar and dealloying process fabricated np Cu rods. The
Gasar process, that consists of melting a material in a gas atmosphere to saturate it with
hydrogen followed by directional solidifying under strictly controlled thermodynamic and
kinetic conditions, generates rods with tubular pores oriented along the length of the cylinder
[198]. The method was used to develop Gasar Cu-34.5 wt.% Mn alloys that then were dealloyed
under free corrosion conditions. The resulting bulk samples exhibited both the longitudinal
pores from the Gasar process, but also lower level nanoporosity from dealloying of the
constituent alloy [199, 200]. The material is shown in Figure 35.

64

Figure 35: Bimodal Cu rods developed by a combination of the Gasar and dealloying processes
to produce nanopores (a) within longitudinal pores (b). Reproduced from [199].

The initial microstructure of the alloy before selective dissolution can also impact the
overall structure of the material. Channeling contrast in focused ion beam imaging [201] and
later electron backscatter diffraction imaging [40, 202], showed that nanoporous gold inherits
the polycrystalline microstructure of the master alloy, with perfectly coherent crystallites tens
of micron in size, in spite of the underlying nanoporosity. Depending on the processing of the
master alloy, its grain boundaries can either be completely dissolved to form a gap or can be
continuous across the boundary [203, 204]. Obviously, mechanical integrity and, even more so,
good mechanical behavior under load makes it mandatory to avoid weakening of the porous
structure at grain boundaries. Dealloying protocols for Ag-Au and Cu-Au can be tailored to
achieve this, as is evidenced by transgranular fracture, as opposed to intergranular fracture
along grain boundaries [204, 205].
In some instances, it was deemed advantageous to accept the extreme brittleness that
comes from preformed cracks at grain boundaries to achieve multiphase microstructures. Thus,
Detsi et al. found that a combination of rolling, from 5 mm to sub-millimeter material thickness,
65

and heat treatment of the master alloy can generate features in the microstructure (Figure 36a)
that lead to de-cohesion along planar faults during dealloying. Stacked microlayers of np-Au
resulted (Figure 36b), where a length scale is described by the layer thickness and another by
the ligament size [135].

Figure 36: EBSD map showing the grains of a textured alloy after cold rolling. b) Low-
magnification SEM image of layered structure after dealloying. c) Higher magnification SEM
image showing the layered structure with embedded nanoporosity. Reproduced from [135].
4.3. Composite Structures based on nanoporous metals

In this section, synthesis methods for incorporating secondary or tertiary materials for
auxiliary hierarchy are reviewed, particularly deposition methods that add an additional layer
of material onto a np metal surface. Synthesis techniques to develop functional np metal films
on larger structures, as well as hierarchical alloy materials, are discussed.
4.3.1. Film deposition on np ligaments

The challenge of material deposition onto the ligaments of np metals is to achieve
uniform coatings on these ultrahigh aspect ratio materials (where aspect ratio is defined as the
ratio between the pore length and the pore diameter). This is specifically true for bulk np metals
regardless if they have nonhierarchical or hierarchical porosities. Deposition requires mass
transport of some sort (for example via ion diffusion in electrodeposition) through the porous
network. Concurrent deposition leads to depletion of the species to be deposited, which results
66

in inhomogeneous coatings. Achieving truly homogeneous coatings in bulk np materials thus
requires a deposition technique that is self-limiting, i.e. after deposition of a certain amount of
material the surface becomes passivated against further deposition.
One technique that fulfills this requirement is atomic layer deposition (ALD), a
chemical vapor deposition process that relies on two sequential, self-limiting surface reactions.
The technique has been successfully used to deposit extremely thin (as thin as 0.5 nm) and
conformal layers of both metals and metal oxides on np bulk materials such as np metals. The
layer thickness is controlled by the number of ALD cycles. Biener et al. used this technique to
coat np-Au samples with either Al2O3, chosen for its thermal stability, or TiO2, targeted for
catalytic properties [206].
Liquid impregnation with a metal salt solution followed by freeze drying and reduction
has also been successfully applied to uniformly deposit sub-monolayer amounts of Ni on the
ligaments of bulk np-Cu [207]. Here, the amount of deposited metal/metal oxide is controlled
by the metal salt concentration of the solution used for impregnation; freeze drying then
preserves the uniformity of the metal ion distribution during drying [207]. Air drying of the
impregnated sample has also been used for applications where the uniformity of the deposited
material is not a critical issue.
Electroless deposition, underpotential deposition (upd), and galvanic
exchange/replacement are other frequently used techniques to deposit metals or metal oxides
on np metals. However, all of these techniques have limitations with respect to the sample
thickness that can still be uniformly coated (typically 100 nm up to a few microns). The
underlying causes of this limitation include mass transport limitations (electroless deposition),
voltage drop across the sample (upd), and non-local reactions (galvanic exchange/replacement).
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To overcome the diffusion limitation of conventional electroless-plating that results in non-
uniform coatings with the interior of the sample uncoated, Ding et al. developed a gas-liquid
interphase electro-less plating technique to uniformly coat np-Au leaf with Pt [115].
In another example, upd was used to deposit an ultra-thin and conformal Cu layer onto
the np-Au wire arrays discussed in Section 4.1 [117, 118]. The Cu coated arrays are
subsequently submerged in an aqueous solution containing Pt ions where the Cu is replaced by
Pt in a spontaneous displacement reaction, resulting in a Au/Pt bimetallic material [119]. The
same technique was used for Pt deposition on bulk np-Au [120]. Thicker coatings can be
achieved by repeating the upd and ion displacement reaction process, as demonstrated for the
deposition of Pd on np-Au [208-210]. Displacement reactions of Cu with Au and Pd/Pt,
respectively, have also been exploited to decorate np-Cu with Au or Pt metals to generate
inexpensive, core-shell nanostructures [211, 212].
Finally, polymer/np-Au leaf composites have been fabricated by electrochemical
oxidative polymerization of aniline in 0.5 M sulfuric acid [142, 213]. The resulting polyaniline
(PANI) layer shows excellent coverage and uniformity, and can be grown up to a thickness of
5 – 10 nm (Figure 37). The PANI /np-Au composite is an interesting material for electrical
energy storage as it combines the high conductivity of np-Au (high power density) with the
high electrical charge capacitance of PANI (high energy density) [142].
68

Figure 37: Left; color enhanced SEM image of PANI coated np-Au. Right; TEM image of
PANI layer on the surface of np-Au. Reproduced from [213].

4.3.2. Nanoporous metals as coatings

In contrast to composite materials where np metal ligaments have been coated by
supplemental elements, this section discusses np metals being used as a coating on preexisting
porous structures. A schematic representation for one way to generate this type of material is
shown in Figure 38: A template is coated by an alloy film, which is then dealloyed to remove
the less noble element and reveal a np coating. Hierarchical np-Au-C or Ag-C composite
structures have been produced in this manner, where Au-Ag or Al-Ag binary alloys are placed
on macroporous carbon (Figure 39a) [178]. After coating by sputtering or electrodeposition,
the film is dealloyed to generate the np film shown in Figure 39b. The bimodal porosity is
described by a pore diameter of ~800 nm, from the coated carbon, and the np film ligaments,
which are as small as 10 nm for np-Au and 70 nm for np-Ag.
69

Figure 38: Schematic representation of the general process of depositing np metal films onto
cellular structures. First, an alloy is deposited or sputtered onto a porous template. The alloy is
then dealloyed so that the remaining material is a macroporous foam with a nanoporous
coating. Reproduced from [178].

Figure 39: Examples of np metals mounted on cellular scaffolds. (a,b) np Ag on porous
carbon produced by sputtering of AlAg then dealloying [178]. (c,d) np-Au on commercial Ni
foams produced by electrodeposition of AuSn then dealloying [177].

70

The same principles were used to develop a bimetallic electrode of np-Au mounted on
a macroporous Ni foam (Figure 39c-d), which was made by electrodepositing a Au-Sn alloy
film followed by dealloying in 5M NaOH and 1 M H2O2. The pore size of the Ni foam is
between 100-200 µm, while the nanopores are between 30-90 nm [177]. A follow up study on
these materials involved an additional electrodeposition step to coat the surface with Pd for
enhanced H2O2 electroreduction [214].

4.4. Conclusions

In this chapter, several synthesis strategies for the development of np materials with
structural hierarchy have been presented. These include templating and dealloying methods, as
well as multistep approaches that combine more than one processing technique. Strategies have
also been demonstrated for a wide range of metals such as Au and more economical materials
like Cu or Ni. In particular, hierarchical porosities can be realized by selecting suitable multi-
element precursor alloys, alloy compositions, and multi-step dealloying protocols. For example,
the sequential removal of different LNEs from a multi-element precursor alloy in combination
with heat treatments can generate bimodal porosity.
Additionally, the fabrication of hierarchical composites, where the hierarchy comes
from the spatial distribution of the elements, were also included. Synthesis methods for these
hierarchical composite materials include coating of np metal surfaces with films composed of
metal oxides, polymers, or other metals to enhance their chemical, thermal, and mechanical
properties. Np metals can also act as functional coatings on other porous materials by depositing
precursor alloys on macroporous foams and dealloying the film.
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Much progress has been made towards developing enhanced np materials through
leveraging structure or symbiotic composites, and several applications are being explored. One
particular area that deserves attention is the use of np structures to reduce the need of precious
metals such as Pt or Pd for electrocatalysis. With the increased surface area and enhanced
transport offered by hierarchical materials, cheaper materials may be able replace precious
metal systems. Alternatively, alloying materials can also enhance the material properties of Pt,
Pd, or Au alloys.
Finally, scaling up and streamlining synthesis procedures must be on the forefront of
experimental work, such as the work that will be presented in Chapter 5. The usefulness of new
hierarchical structures in practice is linked to their size and straightforward production.
The potential practical applications of np metals with structural hierarchy are vast. The
combination of control over the pore feature sizes and surface/interface composition may be
the key to thoroughly understanding the influence of material structure and composition of
materials on electrochemical behavior, which has repercussions in both catalysis and energy
systems. It is the purpose of this literature review to consolidate recent work on hierarchical
structures and guide researchers towards newer unexplored synthesis procedures that may lead
to substantial insight on structure and functionality.

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73

Chapter 5 : Synthesis of Hierarchical Nanoporous Gold Structures
Hierarchical materials with tailored features provide promising structures for
applications requiring light-weight, yet strong, materials. The superior mechanical performance
of periodic materials has led to the development of novel cellular topologies such as the hollow
metallic microlattices discussed earlier [19, 50, 152]. Simulations on these materials indicate
that significant strengthening of the overall structure can be achieved by modifying the
thickness of the cell walls [122]. Thus, an opportunity remains to increase strength by
introducing new morphological features in these materials. In particular, nanosized features
could improve deformation behavior, several examples of this include ceramic lattices with nm
thick walls that display ductility when compressed [121], microlattices with nanocrystalline Ni
walls that recover after deformation [50], and collagen fiber, the nanosized building block of
bone, that provides toughness in the otherwise brittle bone composition [215].
In this study, a method for synthesizing nanoporous gold tubes as a hierarchical structure
is presented, where the ligaments and pores of the foam are the smallest structural constituents.
Np foams are used as the walls of tubular structures in order to implement another level of
hierarchy. The np Au network is generated through a multi-step process that uses a combination
of synthesis techniques such as electrodeposition and dealloying. The as-prepared ligament size
of the np foam can be adjusted, allowing an additional controllable parameter in the hierarchical
material. The main challenges are expanding the processing space from basic geometries such
as films and creating complex ordered nanoporous structures. Here we suggest one method to
synthesize tubular nanoporous gold structures using established techniques that can be scaled
up and applied to custom and more complex structures in future studies. In order to realize these
architectures in bulk samples we must first demonstrate the proof of concept of this study by
74

developing hollow cylinders, representatives of lattice struts, with architected walls as building
blocks for hollow lattices. This is illustrated in Figure 40. The work presented in this chapter
was published in Advanced Engineering Materials and can be found in ref [173].

Figure 40: Visualized proof of concept. Hollow cylinders as lattice struts with an embedded
nanoporous structure.

5.1. Experimental methods for np gold tubes
A 1 cm long gold tube (Goodfellow, 99.99%) with an outer diameter of 0.3 mm and inner
diameter of 0.1 mm was coated with a silver layer via electrodeposition in a silver nitrate
(AgNO3) based commercial plating solution (Caswell). The electrochemical setup was placed
in a beaker and used a conventional cell with a cylindrical Ag anode (Caswell, 99.97%), the
previously mentioned gold tube as the cathode, and a Ag/AgCl reference electrode. The
cylindrical anode completely surrounded the gold cathode and had a small slit along the vertical
length. The surface of the gold tube was prepared for plating by submersion in 17.5% sulfuric
acid and then rinsed in Deionized (DI) water. The electrolyte bath was maintained at room
temperature and a pH of 9.0. Electrolyte solutions were mechanically stirred with a magnetic
stir rod at 400 revolutions per minute and actively filtered during electroplating. Plating was
75

conducted with direct current and then later changed to a reverse pulsed current for subsequent
samples. The initial current density used was 21.6 mA cm
-2
. This was later reduced to 10.8 mA
cm
-2
for both DC and reversed pulse current conditions. Forward and reverse alternating current
of the same magnitude were controlled by applying a square waveform. Each cycle was a total
of 1 s with the forward (positive) current active for 0.85 s and the reverse (negative) current for
0.15 s. Total plating time was determined by the desired thickness based on the dimensions
and mass of the starting material. The electrolyte was replenished with a proprietary brightener
(Caswell) and the pH of the solution was readjusted to 9.0 with KOH before each plating cycle.
After the tubes were coated in Ag, they were placed in a furnace for 24 hours at 900 °C
under flowing argon. The outer diameter of the tube was polished to reduce surface roughness
and the tube was subsequently dealloyed by free corrosion in 70% nitric acid for 24 hours. After
the samples were removed from the acid they were rinsed in DI water and air dried. Additional
heat treatments were performed on select samples at 200 °C or 400 °C for 2 hours in air. A
schematic of the synthesis procedure and a representative cross section for each processing step
is given in Figure 41.
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Figure 41: Schematic representation and cross sectional images for the processing of
nanoporous gold tubes. Fabrication begins with a base Au tube (a) that is coated with a layer
of Ag via electrodeposition (b). The layered tube is homogenized in a furnace (c) and then
dealloyed to produce a nanoporous Au tube (d). The cross-section images for processing at
each step are shown in (e-h).

5.2. Results
In the present study, the wall thickness can be determined by the thickness of the initial
layer of gold. A gold tube (Figure 41a) was selected as the starting material because its hollow
cross section represents typical ligaments found in hierarchical structures such as lattices.
Figure 41b shows the Au tube plated with a Ag layer. Once the layers are diffused, the alloy
composition is approximately Au30 Ag70 at%. The AuAg dealloying system was used to
confirm the feasibility of this processing method because it is a well-established starting alloy
for the synthesis of nanoporous foams [216]. Other np metals can be synthesized from different
alloys such as np copper from CuZn [73] or np silver from AgAl [217]. However, the current
77

synthesis procedure is not feasible using these alloys since large differences in the melting
temperatures of the elements could cause sample distortion during diffusion of the metal layers.

Figure 42: (a) Cross sectional micrograph of homogenized AuAg alloy. (b) EDS spectra
confirming that the alloy is within the target compositional range.

Cross sectional micrographs of coated and homogenized tubes are presented in Figures
41f and 41g, respectively. Figure 41f indicates that the electrodeposited silver layer adheres
well to the gold tube. This is attributed to the inherently contaminant-free gold surface and the
similar atomic size and crystal structure of the two metals, which are both key factors in uniform
surface adhesion. The result of homogenization of the AuAg alloy is illustrated in Figure 42a,
where the distinction between the two metal layers no longer exists. The homogenized layer is
uniform and has slight Kirkendall porosity or pores caused by the diffusion of the metals. Alloy
composition determined by EDS, shown in Figure 42b, further confirmed the complete
diffusion of the layers. Homogenized tubes were subsequently dealloyed through a chemical
dissolution process. After dealloying, a structurally stable nanoporous Au tube was produced
(Figure 43e). The embedded nanoporous network can be seen in Figure 43f.
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Figure 43: Electrodeposition with direct current and a current density of 21.6 Acm
-2
showing
a rough silver top surface (a). Homogenization and dealloying led to a bimodal tubular
structure with an inner nanoporous network (b) and an outer macroporous region (c) that also
contained a nanoporous network inside larger ligaments. Pulsed current with polarity reversal
at a current density of 10.8 mA cm
-2
promoted even plating (d). SEM images of an optimized
nanoporous Au tube with np morphology inside structure walls (f).

5.3. Effect of electrodeposition parameters
The main challenge in developing the processing method was determining the critical
electrodeposition parameters. Figure 43 demonstrates the impact of plating conditions on the
morphology of the plated layer. For DC current plating at 21.6 mA cm
-2
, the silver layer grew
radially from the surface of the tube with round hills on the exposed surface of the silver layer
(Figure 43a). Voids are present in the less densely packed areas of the deposited silver.
Reducing the current density to 10.8 mA cm
-2
during plating did not improve the roughness of
the silver surface or the density of the deposit. Homogenization and dealloying were performed
79

on the DC plated samples plated at 21.6 mA cm
-2
. Dealloying of the tube successfully created
a nanoporous network (Figure 43c). Large voids were formed during the diffusion process in
the region where the silver layer was deposited, while the inner section is free of large voids.
To further improve the coating quality, a reverse pulse current was used during the
electrodeposition of silver. Using this type of current signal, silver deposition occurs in the
cathodic (forward) pulse while a small amount of the metal is stripped away in the shorter
anodic (reverse) pulse. Figure 43d shows a cross sectional image of a tube after reverse pulse
electrodeposition at 10.8 mA cm
-2
. Compared to the DC plated tube in Figure 43a, the surface
of the silver layer was significantly improved by use of the reverse pulse conditions. The two
important differences between the coatings in Figures 43a and 43d are the large roughness at
the top surface and the non-uniform thickness of the silver layer around the tube. Using the
reverse pulse plating, both of these issues were mitigated, shown in Figure 43d. During the DC
plating a negatively charged diffusion layer surrounds the cathode, making it difficult for metal
ions to reach the deposition surface. By reversing the polarity of the current, this layer can be
replenished, making it easier for silver ions to coat the cathode surface. This also reduces the
variance in the local current densities of the wire that tend to promote inconsistencies in the
layer thickness. Finally, silver was removed from the exposed surface during the reverse pulses,
limiting the continued growth of silver on protruding surface features and reducing the top
surface roughness, resulting in a much smoother deposit than previously observed.
Improvement in coating quality carried over to improvements in the subsequent
processing steps where the tubes were homogenized at elevated temperatures and dealloyed.
This was accomplished using the same conditions as were used for the DC plated samples.
Figure 43e illustrates the typical structure of tubes synthesized with optimized plating
80

conditions. These samples presented a single tubular region without the large voids that were
generated during homogenization in the DC plated samples. The average ligament size of the
dealloyed samples was ~45 nm (Figure 43f), similar to np Au films synthesized through free
corrosion in nitric acid [32]. EDS scans after dealloying revealed trace amounts of Ag (<4 at%)
in the samples.

5.4. Post-process heat treatments for ligament coarsening
The nanoporous tubes in this study were synthesized with three different ligament sizes
(Figure 44). Annealing treatments were used to coarsen the ligament diameter and morphology
of the np Au; thermal conditions and resulting ligament diameters are summarized in Table 1.
It has been shown that by adjusting the pore size of the foam, the material can be optimized for
applications in catalytic materials [55, 218], microfluidic devices [219], and as structures for
hybrid electrodes where a large surface area is needed [220]. The SEM image in Figure 44a
shows a typical as prepared sample with an average ligament diameter of 45 nm. SEM images
of the coarsened samples display ligaments that were grown to ~80 nm at 200 °C in Figure 44b
and ~450 nm at 400 °C in Figure 44c. These results agree well with other studies exploring np
foams with coarsened ligaments [78]. Overall, the samples maintained their tubular shape and
structural integrity after the heat treatments.
81

Figure 44: Tunable ligament sizes for nanoporous Au tubes (a) as-prepared and after 2 hour
heat treatments at (b) 200 °C and (c) 400 °C.

Table 2: Ligament sizes of heat treated nanoporous tubes

Sample Type Heat Treatment Ligament Size [nm]
Au
30
Ag
70
As-prepared 45
Au
30
Ag
70
200 °C, 2 Hrs 80
Au
30
Ag
70
400 °C, 2 Hrs 450

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5.5. Discussion
The key advancement towards improving the mechanical behavior of lattice structures
has been the optimization of wall thickness to strut radius ratio [50, 121], where both of these
parameters can be controlled through the method presented here. For example, Shaedler et al.
have demonstrated that nanocrystalline nickel microlattices, with the smallest feature being wall
thicknesses of hundreds of nanometers, exhibits dimensional recovery from 50% compression
[50]. Similar results have also been achieved with ceramic lattices at an even smaller scale with
wall thicknesses below 100 nm [121]. In both studies, the main cause of failure in the lattices
has been high stress concentration at the nodes, which leads to node cracking. Node failure is
expected since they are only supported by hollow strut walls. To realize the potential strength
of these ultralight materials, the critical factor is improving nodal strength [122]. Additional
levels of structural hierarchy may be the key to solving the nodal issue and implementing a new
parameter which may improve the nodal strength of the structures. Np structures can potentially
provide this added strength. Mechanical studies on np Au and np Cu demonstrate higher than
expected yield strengths [78, 221]. Np Au follows scaling laws developed for macroporous
foams, while the individual ligaments behave similarly to single nanowires [29, 35]. As a bulk
material, np metals exhibit densification under uniaxial compression which promotes
macroscopic work hardening [40, 73]. A study by Hodge et al. also showed that the strength of
np gold increases with decreasing ligament size [78]. Therefore, the strength of the tube can be
controlled by coarsening, although the strength of np Au is accompanied by limited ductility
and brittle failure modes of the material. Recent studies have demonstrated that adjustments to
dealloying conditions and impregnation of the pores with a polymer can significantly improve
83

the ductility of the foam [222]. Further hybridization of the np Au would impart ductility in
both tension and compression of the walls.
The step-by-step method described in the experimental section was designed for the
future application of this technique to custom template structures. New synthesis techniques
will facilitate the development of materials with high degrees of optimization for multiple
properties. For example, one can tune mechanical properties from the micro-scale architecture
and catalytic properties from the available surface area. The processing method presented is
similar to template techniques where a layer of material is placed on structures of a
predetermined shape [50]. However, those methods only allow for the wall thickness to be
modified. In this manuscript, the fabrication method involves further modification of the
material once it has been deposited on the template, allowing for nanopores to be implemented
in the overall structure. This process also introduces another parameter for structure
optimization. Furthermore, the tubular nanoporous structures are unique because they combine
both periodic and stochastic structures.
5.6. Summary
In summary, we have demonstrated a protocol for the synthesis of nanoporous Au tubes
using a combination of electrodeposition, homogenization, and free-corrosion dealloying. The
tubes maintained their structure throughout the processing sequence and yielded a nanoporous
network within the tube macrostructure. Heat treatments were used to further modify the
ligaments sizes and give another level of dimensional control. Future work will focus on
expanding the processing space, in order to synthesize more complex hierarchical structures
which can be used for a wide range of applications, including sensors and nano-devices.

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Preliminary work on applying the synthesis methods discussed in this chapter to more
complex configurations can be found in Appendix A.

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Chapter 6 : Fabrication and Mechanical Behavior of Metal Coated Cellular
Structures

Improvements in additive manufacturing (AM) techniques, including the reduction of
minimum achievable feature sizes and the expansion of the number of materials that can be
printed, have altered the structure optimization component of design. Rather than developing
materials or structures based on manufacturability, they can be designed based on function with
virtually no limits on configuration [223, 224]. Advances in metal printing, the fastest growing
sector of AM [225], have shown particular promise for aerospace applications, where metallic
structures can be built in configurations that are not amenable to traditional machining. This
means that components that were previously too costly or physically impossible for production
can now be produced in a manner that reduces the overall amount of material, yielding a
significant reduction in mass which is critical in payload oriented industries [226, 227].
However, to date, only 10 alloys can be reliably printed out of about 5500 commercial metal
alloys in existence [228]. Therefore, there is a critical need for improvements to metal additive
manufacturing or alternative methods in order to generate arbitrary metal shapes from any
desired alloy.
Compared to 3D polymer printing, where numerous printing systems exist, metal additive
technology still has several major areas for growth, such as improving surface finishes [229]
and reducing the internal stresses in printed parts [230]. One relevant area to this study is the
reduction of the achievable minimum feature size for metal materials with arbitrary structures.
Currently, metal microscale devices with feature sizes as small as 10 microns can be generated
by using an electrochemical fabrication (EFAB) process, which is essentially electrodeposition
of fine layers [231]. For other metal printing systems, such as selective laser melting (SLM),
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the minimum feature size is about 150 μm [14]. The ability to produce materials with fine
feature sizes is important when generating structures such as hierarchical materials which can
be used for a wider range of applications, from photonics to catalysis.
As an alternative, functional coatings on microscale polymer scaffolds can be used when
the preferred ceramic or metal structures cannot be directly printed in complex hierarchical
configurations. Two previously used deposition methods include atomic layer deposition
(ALD) and electroless deposition, which apply uniform and conformal coatings on the
structures [110, 114, 232]. However, these methods, are limited to materials that undergo
specific chemical reactions, such as in the case of ALD which produces mainly oxide films like
Al2O3 or TiO2 [206]. Electrodeposition of metals is likewise restricted to single elements or to
alloy systems such as CuNi and CuZn, where overlapping deposition voltage windows can be
achieved with the addition of cyanide as a complexing agent [233]. These constraints limit the
materials that can be used to coat lattice structures and eliminates commercially relevant
structural materials such as Al or Ti alloys, which cannot be plated through ALD or
electrodeposition.
In this work, composite microcellular structures were produced with 3D direct laser
writing (3D-DLW) and then coated with three different metals via magnetron sputtering, a
physical vapor deposition (PVD) process. By using this deposition method, nearly any metallic
target material can be sputtered, including commercial alloys. Pairing this material selection
independence with the technology to print polymer structures at microscale dimensions, is an
innovative approach to develop microcellular structures with both arbitrary shapes and metal
coatings. This allows researchers to move beyond the limited technology to print metals at the
microscale and achieve a previously unattainable range of material functionality. While other
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studies have explored sputtering to coat microlattice structures for mechanical or functional
purposes, this work focuses on an in-depth evaluation of coating quality and multiple types of
coatings on several sizes of structures. The metals used in this study, which are pure Al, Inconel
600, and Ti 6Al-4V, were purposefully selected to extend across a range of characteristic
material properties, such as strength, thermal stability, and crystalline structure, so that the
broad application and impact of this technique can be established. By leveraging the versatility
of magnetron sputtering and 3D-DLW to synthesize metal-polymer composite microcellular
truss structures, tailorable complex architected metallic structures at the microscale could be
fully realized.
6.1. Experimental Methods

The sequential protocol for developing metal-polymer tetrahedral trusses is shown in
Figure 45. The structures were designed and produced using 3D-DLW with a Nanoscribe
Photonic Professional as shown in Figure 45a and 45b. The proprietary, epoxy based resist IP-
Dip (Nanoscribe GmbH) was used as the base material for all of the structures used in this
study. Further details on 3D-DLW can be found in an article by Valdevit and Bauer [234].

89

Figure 45: Schematic showing the design and development of polymer-metal truss structures.
a)-b) A tetrahedral truss structure model is designed and subsequently fabricated using 3D
direct laser writing. c) After being written, the structures are coated with a metal film using
magnetron sputtering. Inset shows schematic of a strut cross section. (d) SEM images of the
actual structures after processing is complete.

Four different sizes of tetrahedral truss structures [98, 235] (A-D) were printed on glass
substrates in the configuration shown in Figure 46a. Sample dimensions are shown in Table 3,
where strut length refers to each individual lattice member that makes up the truss.

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Table 3: Summary of polymer tetrahedral truss dimensions per type of structure.

Structure
Nominal
Strut
Length
(μm)
Nominal
Height
(μm)
Strut Length
-to-
Equivalent
Diameter
Ratio
Relative Density
A 5.0 17.3 5.7 15.5 %
B 7.5 26.0 6.9 11.0 %
C 10 34.6 7.5 9.7 %
D 12.5 43.3 8.1 8.5 %

All samples were sputtered in a vacuum chamber evacuated to at least 4.5 x 10
-6
Torr.
Prior to sputtering the chamber was filled with Argon. Each metal coating was sputtered using
a working Ar pressure of 4 mTorr and a power of 65 W, only the sputtering times were adjusted
to compensate for differences in sputtering rates. The structures were placed 10.2 cm away from
a 7.6 cm diameter sputtering target of the selected source material. Three different target
materials were used in this study; Aluminum (99.95%, Plasma Materials), Ti-6Al-4V (Plasma
Materials), and Inconel 600 (prepared in-house). Two sputtering configurations were employed
and are shown in Figure 46b and 46c. In one configuration, samples were sputtered using a
stationary setup, where the sample holder is in direct line-of-sight of the target and actively
cooled from the backside using flowing water at 16 °C (Figure 46b). The second setup uses a
holder that rotates at 54 rpm throughout sputtering (Figure 46c) without cooling. Consequently,
the structures are not in line-of-sight of the target for half the sputtering time. Nominal thickness
is defined as the expected thickness of a film on a flat surface based on calculated sputtering
rates. An example of a coated microtruss structure is shown in the SEM images in Figure 45d.
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Additionally, 1 micron thick films were sputtered on Si wafers from each target for material
characterization.

Figure 46: Sputtering setup: Left shows SEM image of sample layout on the glass slide.
Middle shows the representative sputtering schematic for stationary sputtering. Right shows
schematic for rotating sputtering setup with samples mounted in the center of the rotating
holder. Sample is rotated 54 times per minute.

The metal films were characterized by scanning electron microscopy (SEM), electron
dispersive spectroscopy (EDS) and X-ray diffraction (XRD). XRD scans were acquired using
a Rigaku Ultima-IV X-Ray diffraction machine. A series of 2θ scans from 30° to 100° with the
rate of 3°/min were performed. EDS scans were performed using an EDAX Element Silicon
Drift Detector for characteristic X-ray detection and EDAX Genesis software for data
processing and elemental determination on both the 1 micron thick samples and directly on the
structures. For EDS on the structures, lower accelerating voltages were used to reduce the
elemental peaks from the epoxy based sample. SEM was performed using a JSM-7001F-LV
field-emission SEM, moreover, images of tested truss structures were taken with a helium ion
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microscope (ORION Nanofab, Zeiss). TEM images were taken on an FEI Morgagni 268 at an
accelerating voltage of 80 kV.
Film thicknesses were measured using aligned truss struts like that shown in Figure 50
below. Measurements were made using a MATLAB script that opens the images and makes
measurements by detecting contrast in the images between the dark film and light epoxy. A
hundred measurements were made along each side of the strut. Average thickness
measurements are shown in the Results section.
A JEOL JIB-4500 FIB was used in this study to explore the cross sections of micron
thick sputter deposited films as well as for the FIB milling of coated truss structures to expose
their cross section. The samples were also sectioned using microtome, where the samples are
embedded in an epoxy resin (Epon), heat polymerized, and removed from the glass slide using
thermal cycling in liquid nitrogen. The bottom of the sample, where the structures were
removed from the glass slide was then covered with the same epoxy. Once embedded in epoxy,
the samples were cut into 100 nm thin sections with a diamond knife. The sections were floated
on water and picked up using a formvar film slot grid. Structures A and D with Al and Inconel
600 coatings were sectioned using this method.
Uniaxial compression tests were performed using a G200 XP nanoindenter (Agilent
Technologies) with a diamond flat punch of 100 microns in diameter for structures A and B and
200 microns for C and D, respectively. Compression tests were carried out for truss structures
with 81nm Al, 205nm Al, 35nm Inconel 600, and 35nm Ti 6Al-4V coatings. The tests were
carried out at a constant displacement rate of 100 nm/s on 3 trusses of each type of sample.
Except for sample type A with 81nm Al coating and sample type D with 35nm Inconel 600
coating for which only 2 samples were tested. The structures in this study were not loaded until
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full densification, but to a fixed maximum displacement. The stress and strain of the samples
were analyzed using the nominal hexagonal cross-sectional area of the trusses and the nominal
height of the samples, respectively.

6.2. Results and Discussion

6.2.1. Film Synthesis

Table 4 shows a summary of the metals used to coat the structures and their
corresponding nominal thickness, which is defined as the thickness of the coating on a flat
surface for a given sputtering time. Figure 47 a-c shows images of the structures coated with
their respective metal films including respective XRD and EDS measurements. Visually, the
coatings for pure Al are significantly rougher than Inconel 600 or Ti 6Al-4V, which is typical
for PVD Al deposition.

Table 4: Summary of sputtering targets and deposition parameters. Note that all the structures
are coated simultaneously in each sputtering procedure.

Type of
Coating
Set-up
Sputtering
Time
Nominal
Thickness
Al Stationary 6 min 81 nm
Al Rotating 40 min 205 nm
Inconel 600 Stationary 3 min 35 nm
Inconel 600 Stationary 8.5 min 100 nm
Inconel 600 Rotating 80 min 305 nm
Ti 6Al-4V Stationary 6 min 35 nm

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Figure 47: a-c) SEM images of sputtered coatings on tetrahedral truss structures: Aluminum,
Inconel 600, and Ti 6Al-4V. Below each image is a corresponding XRD scan and EDS
spectra for each type of coating from a 1 micron thick film sputtered in the same conditions.
Note that the sputtered alloys have maintained stoichiometry.

For the Al coatings, EDS measurements showed that the coatings are deposited without any
contaminant elements, except for a small Si peak from the substrates, and the XRD data showed
that the deposited material is textured along the (111) growth direction. The textured
microstructure has been observed in other sputtered Al films [236] and was also present in the
Inconel 600. The elemental composition of the as-sputtered alloys, Inconel 600 and Ti 6Al-4V,
matches closely to the expected values with a slight deviation of about less than 2 wt% for some
a) b) c)
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elements. However, the deviations in composition is expected from the multielement detection
for qualitative EDS analysis [237].
XRD data of the Ti alloy reveals that sputtering of the commercial alloy results in a textured
β-phase Ti microstructure, where corresponding peaks were observed at 38.7° (110) and 82.9°
(220) with no α-phase peaks detected. The presence of a β-phase dominated microstructure is
unusual since the β transus of Ti 6Al-4V is about 980°C [238]. However, the formation of β-
phase in sputtered Ti alloys has been previously observed and attributed to the non-equilibrium
sputtering process [239]. Essentially, the energy of the condensing atoms is sufficiently high to
form the high-temperature phase, but the structure is “quenched” as the atoms arrive at the cool
substrate, thus forming the metastable phase [239]. The α-phase could form by changing
conditions such as temperature, as has been shown for other sputtered materials [240].
Compared to sputtering, the electrodeposition of Ti and Al, for example, is only
accomplished with non-aqueous solutions. Plating Ti can be done using a molten salt bath [241,
242], whereas the most common method for Al plating is with ionic liquid plating solutions
[243, 244], which require oxygen free environments for optimal performance.
Electrodeposition of any kind requires the optimization and balance of several influencing
factors. In contrast, the sputtering process presented here uses the same experimental setup for
three completely different metals, where only the source material or target is altered.
Furthermore, the sputtering rate of metals can be easily increased by increasing the power input
(target polarity) of the sputtering system [245, 246]. However, since magnetron sputtering
deposition process is momentum driven and line of sight sensitive, surface coverage of the
metal coatings needs to be addressed. The effect of sputtering on coating uniformity is
addressed in the next section, where the coatings are assessed with imaging techniques.
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6.2.2. Cross Section Evaluation

To closely evaluate the coatings, especially the film thickness, FIB cuts were made
across the center of the structure so that half of the sample was removed and the entire cross
section revealed (Figure 48a). This allowed for observation of the coating variation from the
top to the bottom of the structure as well as from the interior to the exterior. Closer evaluation
of the ligament cross sections with SEM, in Figure 48b, showed that FIB milling of the
tetrahedral trusses resulted in significant FIB redeposition of the eroded material inside the
structure. The redeposition layer becomes more pronounced towards the bottom of the structure,
which is the area nearest the glass surface. It is apparent that the ligament cross-sections and
the deposited layer of material is obscured by FIB damage and cannot be easily measured. The
redeposition layer is visible for all structure sizes and material combinations. However, the
layer is most prevalent in structure size A, which has the smallest strut length and highest
density.

Figure 48: a) FIB image of a tetrahedral truss structure milled in half to expose the coating
cross section. b) SEM image of a strut cross section exposed with FIB milling. The actual
coating material is obscured by material re-deposition from the part of the truss structure that
was eroded during FIB milling.

a) b)
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The effect of FIB redeposition on sample preparation has been observed for solid
samples, where redeposited material lands in the area to be sputtered [247]. However, the effect
is exacerbated for 3D materials because the material lands within the area that will be evaluated.
While several studies have used FIB milling as a method to section microlattice materials or to
expose the polymer core for etching [111, 113, 121, 248], to the author’s knowledge, only one
study has mentioned the possibility of FIB milling redeposition obscuring the cross-sectional
evaluation [249], although the issue was not further investigated. A recent study suggests that
the thicker films on the bottom side of the struts is the result of redeposition from atoms during
film growth and not from FIB milling [250]. Several factors may be contributing to FIB
redeposition, such as the FIB instrument quality or the structure density, which is an indicator
of how “open” the cellular material is to penetration from the metal atoms during sputtering.
Another factor may be the strut geometry; our study employs rectangular cross-sections while
others have circular or oval cross sections [111, 250]. However, our results indicate that FIB
redeposition from micromachining of 3D parts will result in some degree of redeposition that
is more prevalent at the bottom of the struts. Therefore, special care must be taken when
evaluating the coatings on 3D materials, particularly for methods that are known to yield non-
uniform coatings (i.e. electrodeposition and sputtering).
In order to mitigate artifacts from FIB and accurately measure the coating thickness,
microtome was used to embed the structures in epoxy so that the metal coating is preserved
during sectioning. Microtome has been shown to be a reliable method for sectioning biological
samples thin enough for electron microscopy [251, 252]. The sectioning method can be difficult
for hard samples because there is separation between the embedded epoxy from the sample as
well as damage incurred by the sectioning knife. However, since the metal layers in this study
98

are thin, and the core of the structure is epoxy based, the samples can be sectioned. An example
of the cross sections that can be obtained from this method are shown in Figure 49, where both
a structure coated with the stationary setup (Figure 49a) and rotating setup (Figure 49b) are
shown for Al coated structures with dimensions corresponding to structure size D. The nominal
Al coating corresponding to the stationary and rotating setup are 81 nm and 205 nm,
respectively. For both samples the coating thickness is highest on the top side of the structure
(red arrows), which directly face the sputtering target during the coating process. The rotated
sample also shows a dark and thicker coating on the outer edges of the samples (white arrows),
which are exposed to the sputtering target upon rotation.
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Figure 49: Microtome prepared cross sections of Al coated structures with either the
stationary (a) or rotating (b) configurations. Images shown are for structure size D with 12.5
µm strut length. The red arrows for both images show the thicker top side coatings, while the
white arrows show the thicker outer coatings generated with the rotating sample.

The coatings were further assessed by imaging the aligned truss ligaments at higher
magnification, like the TEM images in Figure 50, which was coated in a stationary
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configuration. A coating thickness variation is present along the metal/polymer interface when
comparing the left side of the strut in Figure 50a to the right side. This variation is not
uncommon with magnetron sputtering, a PVD process, which relies on atoms being ejected
from a source material as a result of bombardment from ionized gas particles. The process is
inherently momentum driven and therefore, the coating of complex geometries is challenging
due to a line-of-sight dependency [253-255].

Figure 50: TEM micrographs of strut cross-sections coated in the stationary setup from center
of the structure (a) and from the bottom of the structure (b), which was masked during
sputtering.

One method to improve the coating uniformity of the film is to rotate the substrate
during sputtering in order to increase the angles of exposure from the sputtering target to the
cellular structure. Rotation of the substrate in this study confirmed that the coating thickness
can be influenced; demonstrated by the increase in the outer edge coating thickness, which was
periodically exposed to the target upon rotation. In our study, the samples were rotated along
one axis during sputtering, as shown in Figure 46c. The coating variation could potentially be
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improved by rotating the cellular structures on multiple axes, by rotating the sample layout by
90°, or by mounting the glass slide perpendicular to the current experimental setup.
Aside from rotation, modified PVD processes have been previously used to coat macro-
cellular materials with metals, including coating carbon foams with titanium [256] or coating
polymer foams with a nickel-based super alloy [257]. Both these studies manipulated the flow
of the metal vapor to coat the structures more uniformly. Additional techniques to coat complex
shapes with PVD include the use of cylindrical sputtering targets, which completely surround
the substrate and help prevent material waste by localizing the sputtered material inside the
target [254]. The use of directed gas flow and pressure can also be used to mitigate uniformity
issues without substrate rotation [253, 255].
To investigate the line-of-sight coating issues for the bottom side of the struts, a
comparison was made between a strut from the interior of the structure (Figure 50a) and a strut
that was attached to the glass slide during sputtering. The strut attached to the glass slide is
completely masked during sputtering and cannot be coated (Figure 50b). Figure 50a shows that
the sample is coated on the bottom side of the structure facing away from the sputtering source,
while Figure 50b shows the embedded polymer-polymer interface. By comparing the two
images, it was concluded that the bottom sides of the samples are coated with a thin layer of
material. This was observed for all materials sectioned with microtome.
A collection of individual cross-section images for structure A from the rotated Al
sample is shown in Figure 51, where the images are placed relative to their locations within the
full structure cross section. From this visual perspective, it is apparent that the use of a rotating
substrate increases the coating thickness variation, as previously mentioned. If the center
images are compared across each level of the structure, the coating is thickest along the top side
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of the structure and the top side thickness decreases with each level of the structure, moving
from top to bottom. The exterior facing sides of the ligaments are also thicker than the interior
facing sides. This can be clearly seen in the right side image in the second level, where the outer
side has a coating that is considerably thicker. As the structures are rotated, the edges of the
material are directly in sight of the sputtering target, producing the thicker exterior coating. This
variation is more pronounced in structure size A, which is the denser structure. This suggests
that structure density or “openness” should be a factor in selecting materials for magnetron
sputtering of complex shaped materials.

103

Figure 51: TEM images of strut cross sections of structure A rotated and coated with Al. The
zoomed in images of the struts are arranged according to the cross section and exhibit a
decreasing top surface thickness of the coating at the lower levels of the truss structure.

In Figure 52, the thickness of the top side coating is shown relative to the level or depth
into the sample the measurements were taken. Values were measured on rotated structure sizes
A and D coated in Al. The average coating of the top surface is around 250 nm for both sizes
of the structures, which is expected since both are equally exposed to the sputtering target.
Measurements made on the 2nd level on the top side of the struts were ~100 nm less than the
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top layer for structure size D and almost half for structure size A. The coating thickness for both
structures is also reduced from the 2nd to 3rd level, but then remains about the same in the 4th
and bottom levels. The drop off in thickness is observed for the top side measurements for all
structures measured in this work; Al rotated ~205 nm, Al stationary ~81 nm, Inconel 600 rotated
~305 nm, and Inconel 600 stationary ~100 nm. This indicates that the sputtering process yields
a characteristic coating on the structures explored in this study.

Figure 52: Measurements of the top side of the struts and the corresponding area of the
structure for the rotated Al structure corresponding to level of structure where the
measurement was taken. In both samples, the thickness of the top level becomes smaller as
the measurements are taken at lower levels in the samples.

105

For the stationary coated samples, the interior facing strut side has a similar coating
thickness to the center region struts. However, for the rotated samples, the exterior facing side
of the struts has a thicker coating compared to the interior and exterior struts; consider the
bottom right image in Figure 51. This is explained by considering that the outer facing struts
are directly exposed to the sputtering target as the sample is rotated during coating. The
difference in interior and outer facing side thickness is not observed in the samples coated with
the stationary sputtering setup. Alternate rotating axes, for example, could be used to more
uniformly distribute the coating. Another idea, would be to sputter the structures with a low
vacuum chamber so that the mean free path of ejected metal atoms is reduced and scattering of
the film atoms is encouraged [258].

6.2.3. Mechanical Testing

After evaluating the coatings, uniaxial microcompression tests were used to determine
the impact of the metal coatings on mechanical properties. A collection of the tests for structure
size A is shown for select configurations of coating thicknesses in Figure 53. The fracture
strength for all coating materials and thicknesses is higher than the uncoated polymer structures
(Curve 1). Curves number 5 and 4 are shown for Inconel 600 and Ti 6Al-4V curves with the
same nominal thickness of ~35 nm. Theoretically, the only difference between these structures
is the coating material, since all other experimental parameters are identical. Therefore, it
follows that for a given nominal thickness the Inconel 600 coated structures have a higher
compressive strength than the Ti 6Al-4V coated structures. Compared to the uncoated scaffold,
Inconel has more than double the compressive strength with only minimal coating. For
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reference, the coating time for this structure was only 3 minutes. Overall, the highest failure
stress was achieved by a ~205 nm thick coating of Al in the rotating coating configuration.

Figure 53: Uniaxial compression stress-strain curves showing the range of mechanical behavior
observed for tetrahedral truss structures of type A coated with different metals. The stress-strain
curves are shifted to account for misalignment between tip and sample surface or roughness
effects (gray region in the graph).

Closer inspection of the Inconel coated structure in Figure 54b shows possible shape
optimization of the truss nodes due to the sputtered metal developing a webbed coating (white
circles). This coating morphology may be the source of a shift in fracture behavior from nodal
failure to fracture at both the nodes and along the struts. Figure 54a shows points of fracture
observed at the truss nodes, shown by the circles, and regions of fracture along the ligaments
are also present (noted by the arrows). This is significant since nodal failure, not ligament
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failure, is a common occurrence for microscale architected materials under compression [50,
98, 110, 152, 235, 259]. The nodal features added by the coating may be providing structural
features that enhance the mechanical efficiency of the trusses. A simulation by Valdevit et al.
showed that the strength of a hollow lattice structure can increase when the volume of material
around the node is thickened relative to the strut regions [122]. However, a comprehensive
mechanical study would be needed to determine the contribution of this coating technique to
mechanical performance.

Figure 54: Helium ion microscopy (HIM) images of compressed truss structures of type A
coated with a ~35 nm layer Inconel 600 in stationary configuration. a) Structure failure at a
truss node is shown and marked with circle. The arrows note failure along the ligaments. b)
Close up image of structure node where the coating is non-conformal and forms a webbed
structural feature.
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6.3. Outlook and Conclusion

The commercialization of 3D-DLW has sparked several studies in a wide range of research
areas including in-depth investigations into the relationship between hierarchical structure and
the mechanics of lattices. Novel microdevices are also being developed, such as sputter coated
microlattices for battery applications [260] and silver deposition on 3D structures for photonic
metamaterials [158, 232]. Advancements in coating capabilities, especially if line-of-sight
effects can be mitigated, would expand these applications so that tailormade structural features
could be accompanied by the unparalleled materials selection offered by magnetron sputtering.
This endeavor is explored in this study, where microlattice metal-polymer composites were
generated using a combination of 3D direct laser writing (3D-DLW) and metal coating via
magnetron sputtering. Using magnetron sputtering, a variety of metal coatings, including
Aluminum, Inconel 600, and Ti 6Al-4V, were deposited onto microlattice structures. The
inclusion of alloy coatings is significant because traditional deposition methods for metals are
limited to a few systems. Therefore, using sputtering as a metal deposition method has the
potential to expand the achievable properties of composite microstruss structures to include
unique combinations of strength, conductivity, and chemical activity. Uniaxial compression of
the samples revealed that, by applying metal coatings, the deformation mode of the composite
structures can be shifted from node-dominated failure towards a ligament dominated failure.
This realization may lead to improved lattice structure and nodal design for future microlattice
structures.

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110

Chapter 7 : Tensile Behavior of Cellular Materials Produced via Direct Laser
Writing

The use of advanced stereolithography techniques, such as the two-photon 3D-DLW
employed by this dissertation, allows scientists to generate topologies of nearly any desired
shape. However, the amount of material or overall size of the structures that can be produced is
limited by the maximum allowed stage travel of a printing system. For the Nanoscribe
instrument used in this work, two motors are employed, including one for the motorized stage
(coarse motor) and a piezo motor, which is used to achieve the fine dimensions for the
microscale structures presented in the study. The print volume for the piezo is limited to 300 x
300 x 300 μm for piezo scan mode and 200 x 200 μm in galvo scan mode, which prints in layers
[261]. Therefore, to increase the size of the structures, multiple piezo printed areas are stitched
together to form a larger structure by moving the motorized stage. Leveraging both motors is
necessary to achieve truly hierarchical structures, where features must be achieved on both the
microscale, determined by minimum feature size, and the macroscale, determined by the overall
print volume. The scaling up of these structures is also critical for implementation and
applications, such as batteries or catalytic surfaces, where a specific volume may be required to
achieve a certain power density or product formation, respectively.
Two fundamental issues arise with stitching multiple areas together. First, with the
movement of two motors, where there is a potential for a shift or offset between the intended
secondary print area and the actual new printing position. This would result in a physical
misalignment in the stitched volumes. The second issue is related to the local polymerization
process that forms the polymer structures. During printing, oxygen atoms can attach to the
growing polymer chain and delay chain growth. As printing progresses the oxygen content
111

decreases, favoring the polymerization. However, during stitching, the process is paused,
allowing the oxygen content to be replenished. When printing is reinitiated, the local oxygen
content of the print volume is different than the oxygen content when the initial print volume
surface was polymerized. This may produce subtle and potentially significant interface issues
that need to be mechanically characterized. Note that a certain amount of oxygen is necessary
to prevent the entire amount of resist to polymerize.
In this chapter, the influence of the stitched interface on the mechanical behavior of
microcellular structures is addressed. Log-pile cellular materials were produced via 3D-DLW
and mechanically tested in tension using a modular microtensile tester with custom grips. Both
testing methodology, implementation, and preliminary mechanical results are presented below.
7.1. Experimental Methods

Log-pile cellular structures were prepared via 3D-DLW using a Nanoscribe GmbH
Photonic Professional DLW in galvo scan mode. The proprietary resist IP-Dip (Nanoscribe
GmBH), which is an epoxy based polymer, was used as the base material for all of the structures
used in this study. To fabricate the structure, resist was placed on top of a glass slide and
mounted upside down in the printing system. Prior to printing, the glass slides were
functionalized to improve adhesion with the structure using a siloxane called TSPMA (3-
[Tris(trimethylsiloxy)silyl]propyl methacrylate) (Sigma Aldrich) followed by acid washing
(HCl/MeOH, H2SO4). The structures were printed on the glass surface in three varying
conditions detailed in Table 5: (A) one group of samples were printed continuously, (B) a
second set were printed with a five minute rest in the middle of the structure, and (C) the last
set was printed with a five minute rest plus a 2 µm overlap in the structure. A schematic of the
112

structure layout on the glass slide is shown in Figure 55a, while the geometry of the structures
is shown in Figure 55b. Note that the center or gage section is comprised of a log-pile
architecture, while the two wider attachment regions are made of full density (i.e. non-porous)
polymer material. For the structures with a rest period, the lower half of the structure is printed
and printing is paused for five minutes after which, the upper half is printed. The overlap is
defined as beginning the printing process 2 μm below the already polymerized structure,
whereas no overlap is defined as restarting the printing process directly on the structure surface.
Further details on generating the structures can be found in the study by Oakdale et al. for IP-
Dip structures [262].

Table 5: Description of samples and printing parameters with rest periods and overlaps.

Pillar Rest Period Overlap/offset
A 0 min 0 µm overlap
B 5 min 0 µm overlap
C 5 min 2 µm overlap

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Figure 55: Schematic of log-pile structures printed on a glass slide. a) Shows a layout of the
printed structures. b) Cross-section of the geometry showing the dimensions of the tensile
specimens. c) Dimetric view of the tensile specimen showing overall shape.

Two options were proposed for gripping and testing the structures in tension; an
adhesive attachment and a slot and key approach (Figure 56). Several adhesives were explored,
with the most promising being an ultra-low shrinkage (<0.07 %), UV cure silica filled epoxy
(EMI-UV 3410). The adhesive was cured using an OmniCure S2000 with a dual waveguide set
at 7 W/cm
2
for 90 seconds. However, despite the low shrinkage of the epoxy, a small volume
change upon curing was sufficient to pre-load and fracture the sample. A load versus time curve
showing the load change during curing and eventual failure of one specimen is shown in Figure
57.

114

Figure 56: Schematics showing two proposed methods for gripping tensile specimens for testing
micropillar specimens. a) Shows an adhesive method where one end of the sample is attached
to a crosshead using an adhesive. b) Schematic shows a mechanical grip where specimens are
placed into a slot piece and pulled until the sample is in tension.

Figure 57: Curve showing the change in load as a tensile sample is attached to the tensile test
crosshead with a UV cure silica filled epoxy. The load increases as the adhesive cures and
continues to increase, eventually causing failure of the sample before a tensile test could be
completed.

115

Subsequent tensile tests were carried out using the testing schematic shown in Figure
56b, a slot and key grip. The grip pieces were generated using DLW to match and accommodate
the tensile sample dimensions. Using a camera mounted on the testing apparatus (see Figure 29
in Chapter 3), each individual sample was aligned with the complimentary slot piece (Figure
58a). After the sample was aligned, the tensile specimen was raised above the slot piece and
the grip section was moved under the specimen (Figure 58b). The tensile specimen is then
lowered into the slot as shown in Figure 58c. After the specimen is in place, the grip section is
moved until the sample is engaged in tension.

Figure 58: Schematic showing the alignment of tensile specimens with slot piece to grip for
tensile tests. a) First the tensile gage section is aligned with the slot and moved above the lock
piece. b) After alignment, the lock-piece is moved into position under the tensile sample. c) The
sample is then lowered into the slot and d) the slot piece is moved into a position where the
sample is engaged in tension.

A custom S-shaped load cell was developed using a binocular micro load cell (Phidgets
CZl639HD) that is capable of measuring a maximum load of 1N with an instrument sensitivity
of about 0.5 % of the maximum load. A schematic of the assembly is shown in Figure 59a and
59b, where the glass slide with the printed tensile samples is attached using double sided carbon
116

tape (Ted Pella 16073). Figure 59c shows the actual assembly with the samples visible as small
dots on the glass slide. The overall tensile setup is shown in Figure 59d, where the light source
for the camera and the slot piece to the right hand side are shown. Once the samples were
slotted and engaged as shown in Figure 58, the tensile tests were performed moving the cross
head or slot piece side at a speed of 0.12 μm/s (the slowest speed of the coarse motor) or at a
1.2 x 10
-3
strain rate.

Figure 59: Load cell assembly for testing microcellular polymer tensile samples. a) Schematic
profile view of load cell. b) To scale schematic of the load cell, tilted so that the samples are
visible. c) Photo of load cell with tensile samples printed on the glass slide. d) Photo of complete
assembly with both the load cell and tensile samples, as well as the slot piece slide.
117

7.2. Results and Discussion

Figure 60 shows load-displacement curves for structure type A, which are continously
printed (shown in blue hues), and structure type B, which were printed with a rest period (shown
in red hues). Overall, the structures exhibited similar failure loads, which indicates that there is
not a significant difference in strength imparted on the structures due to the alternate printing
methods. However, the failure regions between the two types of structures are distinct. The
structures printed with a rest period fractured in the stitch region at the center of the gage
section, while each of the continually printed samples failed on the grip end of the tensile
specimen.

Figure 60: Load vs. displacement data from pillars tested in tension. Blue curves show
continously printed pillars and red curves shows samples printed with a rest period half way
through the printing process.

118

Figure 61 includes green colored test curves for structure C, which has a rest period and
printing overlap. Of the three printing protocols, structures printed with an overlap achieved the
highest failure load. Additionally, structure type C broke in the grip section similar to structure
A. Figure 62 shows representative images of each of the structures showing the failure regions
in each type of sample. The maximum achieved loads and displacements are shown in Table 6.

Figure 61: Load vs. displacement data from pillars tested in tension. Green curves show samples
printed with a rest period and printing overlap. These samples showed the highest maximum
achieved load before failure.

119

Figure 62: Photographs from camera mounted on top of the microtensile tester of each of the
three structures during testing. First column of images show samples prior to failure. The second
column shows images at the moment of failure and the final set of images show the samples
after the slot piece has been moved away from the specimen.

120

Table 6: Maximum load and displacement of pillars just before sample fracture. Pillar C, which
exhibited the hightest failure load, is in bold.

The results suggest that the printing overlap of 2 μm prevented failure in the stitch region
and slightly enhanced the strength of the structure. This behavior may be due to the additional
laser raster in the 2 μm overlap region, which could induce supplemental cross-linking in the
polymer. Studies have shown that as-printed DLW produced materials can be mechanically
enhanced with simple heat treatments or UV curing to facilitate further cross-linking in the
printed material [98, 262]. The additional crosslinking in this study could be related to the
secondary exposure to the polymerizing laser, improving the strength of the interface layer.
While the above work has revealed a potential method for mitigating sample failure at
the printing interfaces, the use of this testing method for measuring tensile loads on the
specimens can be further improved. Several studies have used similar slot key approaches to
testing micro/nanoscale samples or otherwise non-standard tensile specimens with success
[263-265]. However, one potential improvement to this testing method includes the design of a
more standard tensile specimen. The current design includes sharp corners in the grips that are
potential stress concentration areas causing failure of the tensile specimens in the grip. An
angled transition between the gage section and the grip region, which can easily be added by
modifying the sample design for DLW, would be an intriguing approach towards improving
testing reliability.
Sample
Type
Max Load
(mN)
Max Displacement
(μm)
A 45 10.6
B 44 11.4
C 48 12.6
121

This effort to measure the impact of a 3D printed interface on the structural and
mechanical integrity of materials produced with DLW directly addresses issues with scaling up
architected materials with microscale features. With additional testing, more definitive
suggestions, such as an ideal overlap depth or laser speed, could be proposed so that 3D printing
protocols can be improved to mitigate adverse effects from printing large structures with
attachment interfaces. This would increase the achievable length scales in hierarchical
structures, expanding the potential designs and applications for these materials.
122

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123

Chapter 8 : Conclusions and Future Work
As the methods to generate materials with increasing structural complexity continue to
be developed, the field of hierarchical materials with micro and nanoscale features has grown
to incorporate novel applications, strength-enhancing size effects, and a variety of materials
from ceramics to metals. In this work, contributions towards the fabrication and characterization
of metallic open-cell hierarchical materials were made. Metal structures were explored in
particular because they offer both enhanced strength over polymer materials and hold potential
for electrochemical and catalytic applications due to their reactivity and conductivity.
Initial efforts focused on expanding synthesis techniques for producing nanoporous
metals with complex and multimodal configurations. In Chapter 4, it was shown that numerous
methods could be employed to generate np metals with structural hierarchy, including
exploiting the differences in electrochemical behavior of different metals or alloy phases.
Another method included templating alloyed configurations on scaffolds with subsequent
dealloying to induce nanoporosity as the base level of structural hierarchy. The combination of
length scales and structural features that can be achieved with this method is beneficial for
several applications that require both high surface areas for reactivity and large enough pore
sizes for diffusion limited processes.
This work explored one particular templating and dealloying method, where np Au
tubes were fabricated with a step-by-step procedure using the established processes of
electrodeposition, diffusion, and dealloying. First, Au tubes were coated with Ag using a
commercial electroplating bath. It was shown that a more uniform and dense coating layer could
be achieved by employing a reverse pulse plating signal rather than direct current. The Au and
124

Ag layers were diffused in an Ar atmosphere at 900° C and then dealloyed via free corrosion in
concentrated nitric acid. The resulting material was a fully dealloyed np metal tube with an
average as-dealloyed ligament diameter of 45 nm. It was also demonstrated that the ligament
size could be coarsened to 80 nm, or 450 nm after 2 hour heat treatments at 200° C and 400° C,
respectively, which demonstrates the tailorability of these materials. In principle, similar step-
by-step procedures could be applied to more complex geometries. Overall, this study developed
a novel method for incorporating nanoscale features in hierarchical materials.
This dissertation also explored improving the properties of polymer 3D printed
microcellular materials by coating printed structures with metal films. The microcellular truss
structures in this work were designed with a tetrahedral unit cell and fabricated in four different
sizes via 3D direct laser writing (3D-DLW) with a Nanoscribe Photonic Professional (GmbH).
The structures were coated via magnetron sputtering using two sputtering configurations:
stationary, where both the structure and sputtering target are at constant distance and relative
position to each other, or rotating, where the structures are continually rotated during coating.
Using two sputtering configurations revealed that rotating the substrate is a method to control
the coating thickness variation. Three distinct materials were used as coatings: aluminum,
Inconel 600, and Ti-6Al-4V. These materials were selected for their distinctive characteristics,
such as melting temperature and crystal structure. The coatings were characterized using SEM,
XRD, and EDS, revealing that the coatings were all textured, but more importantly, the
elemental compositions of the deposited alloys confirm that the stoichiometry of the sputtering
target was preserved. By employing a variety of materials in this study with minimal
experimental changes, we have demonstrated the adaptability and strength of using sputtering
as a deposition method for coating 3D microstructures.
125

The coated microtrusses were further evaluated by sectioning in order to measure the
coating thickness. Initial attempts to section the structures using FIB milling were inhibited by
significant redeposition of the milled material in the open space of the structure. To evaluate
the unaffected cross section and coating of the material, microtome was used as the sectioning
method. Evaluation of the microtome cross sections revealed that the entire structure was coated
with a layer of material and that the coating thickness of the metal, particularly for the top side
of the struts facing the sputtering target, decreases with each layer as measurements were taken
deeper in the structure. Employing microtome as a sectioning method was a novel approach
towards accurately characterizing and understanding the deposition of 3D microstructures using
magnetron sputtering. In addition, uniaxial compression tests on the structures showed that
structures coated by this method failed at both the nodal region, which is typical, but also along
the truss ligaments, which is atypical. It has been proposed that structural features generated by
the non-conformal sputtered coatings may be shifting the failure modes for these materials. In
summary, the composite structures fabricated by combining DLW and sputtering is identified
as a method to functionalize polymer architected materials in a manner that affords a high
degree of material selection flexibility.
The final topic of this dissertation was dedicated towards testing the mechanical
properties of scaled up microcellular structures, which is hindered by the stage travel limit using
a piezo motor. To yield larger structures, multiple areas are stitched together, which introduces
an interface or potential weakness in the structure. In order to investigate the mechanical
properties of this interface, a mechanical testing protocol and testing stage was designed and
implemented. Log-pile modified dog-bone samples were printed using DLW, where one set of
samples was printed continuously and another set printed with a stitch interface in the center of
126

the gage section. Several adhesives were explored for gripping the dog-bones, which were
printed on glass slides. Ultimately a slot and key method was employed, where the dog-bone
samples were gripped by slotting the sample into a complimentary opening and pulling the
sample in tension. While no pronounced tensile strength differences were observed between
stitched and continuously printed samples, the stitched samples regularly failed along the stitch
interface, and continuously printed samples at the grip. This work explores issues with scaling
up architected materials with microscale features so that 3D printing protocols can be improved
to mitigate stitching effects.
The results presented in this dissertation have examined several topics in the field of
microcellular materials and the continuation of these studies have multiple paths forward. One
intriguing investigation would merge the knowledge gained from coating microcellular
materials and developing np tubes. While the development of np metal structures was explored
and realized for tubular structures, generating more complex np materials is still the ultimate
goal. Since we have shown that microtruss structures can be coated with alloy materials while
maintaining stoichiometry, one interesting study would be to coat cellular structures with a Ag-
Au alloy in atomic concentrations for dealloying. The main objectives for this study would be
to optimize the truss structure coating thickness to accommodate enough np ligaments in the
cross section for structural stability. Since the coating will be deposited in alloy form, a
homogenization step is not necessary, allowing for retention of the polymer scaffold and
structural enhancement. Furthermore, AgAu alloys were originally used as proof-of-principle
materials because the alloy has a simple phase diagram and a well-studied dealloying process.
With sputtering, other more economic and/or lower density np materials could be used to
expand the properties of np materials.
127

Furthermore, improvements to the coating thickness distribution in sputter coated
structures is also a primary concern. As suggested in Chapter 6, there are numerous methods to
mitigate and improve PVD coatings on 3-dimensional substrates such as foams. The most
promising method would be to employ a cylindrical or 3D sputtering target, where the entire
substrate is surrounded and coated from all directions. Additional cellular geometries or unit
cells should also be explored, for example a more open structure with a lower relative density
or larger structures with a higher overall size. Additionally, structures could be fabricated with
strategic printing so that the geometry is not confined by attachment to a substrate.
Finally, mechanical testing of microcellular materials should also be explored in more
detail. Particularly, the influence of sputtered coatings on mechanical properties. The addition
of surface features at the nodes was an interesting result of the sputtered films and could
potentially be a clue towards solving the issue of nodal failure in ordered microcellular
materials. An ideal study would focus on the influence of sputtered coating thickness and
compressive failure. For mechanical assessment of scaled-up structures, the main objective
would be to complete mechanical testing with a dog-bone specimen that more closely resembles
a standard tensile sample. The geometry of the tested structures is not ideal for tensile testing
as the sharp angles at the grip sections provide areas of stress concentration. Steps should be
taken to perform a more standard test so that samples break within the gage section and not at
the grip. This would further inform our understanding of the issues related to scaling up 3D
printed materials and aid in the development of interface mitigating printing protocols.
Overall, this study has contributed towards knowledge and techniques for generating
new hierarchical materials and enhancing material selection flexibility. The processing and
128

testing methods discussed here have shown potential for further contributions towards novel
methods for designing and fabricating materials on multiple length scales and should be studied
further to assess and influence the prospective engineering applications of hierarchical
materials.

129

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151

Appendix A: Complex Nanoporous Structures
The procedure for np Au tubes described in Chapter 5 was modified to generate np
lattice structures. A gold tube was the starting material in the previous fabrication method. Here,
initial structures were developed by coating sacrificial polymer templates, made from the UV
polymerization process described in Chapter 2.4.3, with a layer of Au. This was accomplished
using magnetron sputtering, a physical vapor deposition (PVD) process described in Chapter
3.3. The samples were rotated during the deposition of Au to ensure full coverage of the
structure. Figure 63 shows a sample of the different structures made by this process including
a single strut, a unit cell, and a lattice made of repeated cells.

Figure 63: Polymer truss structures coated in Au by magnetron sputtering. (a) single strut (b)
a unit cell (c) lattice structure made of repeating cells.

After sputtering, the samples were evaluated to assess the coating quality. Cross
sectioning of the lattice revealed that the entire structure was coated with Au. The coating
thickness at the center of the structure is expected to be smaller than that at the edges. Ligaments
in a single unit cell were mounted in epoxy and polished. Figure 64 shows a cross section image
152

of one polymer ligament coated with Au. Measurements determined that an average layer
thickness of ~4.3 μm was achieved.

Figure 64: Cross section of polymer template coated with Au by magnetron sputtering.

After the Au templates were coated, the unit cell was plated with Ag via
electrodeposition as described in Section 3.1. The amount of Ag deposited on the unit cell was
determined by the surface area and average thickness of the Au layer achieved by sputtering.
The polymer interior was etched using 1M NaOH and the sample was subsequently
homogenized in a tube furnace under flowing Argon at 900 °C for 2 hours. Figure 65 shows the
structure after each processing step. The unit cell was structurally stable after sputtering,
electrodeposition, and etching. However, the structure was significantly distorted after the heat
treatment (Figure 65c), which led to the collapse and structural failure during dealloying.
Polymer
template
Mounting epoxy
153

Figure 65: Processing of complex np structures (a) fabrication begins with generating a gold
base structure (b) a layer of Ag is deposited on the gold structure (c) the metals are diffused
during homogenization distorting the structure.

Following the collapse of the structure, new wall architectures and alternate processing
methods were explored to provide thermal stability of the complex structures during heat
treatments and also to improve the overall strength of the np struts. These specific architectures
include those shown in Figure 66, which shows three potential wall designs in tubular
configurations; a polymer infiltrated sample (Figure 66b), a np tube with one face sheet (Figure
66b), and a sandwich design with np Au as the foam core (Figure 66c).
a) b) c)
154

Figure 66: Selection of different hollow strut wall architectures. (a) Polymer infiltrated
sample, (b) tube wall with a single face sheet, (c) and sandwich structure with foam core.

The benefits of hybridizing np Au have been demonstrated by Wang and Weissmuller who
improved the ductility of np Au by infiltrating it with a polymer [222] The polymer infiltrated
design would improve the ductility of the structure, but preventing polymer from filling the
open channel of the tube during fabrication is a challenge. Further, the polymer would prevent
fluid or gas to access the np Au surface, eliminating the possibility to functionalize these
materials for catalysis or other applications that would require an open-cell foam. The sandwich
foam design was selected as the target wall architecture because the additional face sheets can
be implemented with current coating capabilities. The desired final hierarchical structure for
this study is illustrated in the schematic in Figure 61.

155

Figure 67: Schematic representation of a complex hierarchical structure made of hollow struts
with a metal-foam sandwich wall (inset on the right).

The first step in realizing these structures was to select a material for the face sheets and
modify the processing technique. Tantalum was chosen as the face sheet material for its high
strength and thermal stability. Ta is also conductive, which means it can be coated via
electrodeposition, and in later processing steps, it will not significantly diffuse into Au or Ag
during heat treatments. Furthermore, it is possible to sputter Ta with magnetron sputtering so
that template Ta structures can be produced like the Au structures shown in Figure 63.
Acknowledging the potential adhesion issues, purchased Ta tubes were plated with a
thin layer of Ag using the same parameters and set-up using previous processing parameters
(Section 3.1). Another set of Ta tubes were coated with gold via DC electrodeposition with a
commercial plating solution (Caswell Gold) using a flat gold foil as a sacrificial anode. The as-
plated tubes are shown in Figure 68.

156

Figure 68: a) Ta tube coated with Au via electrodeposition b) Ta tube coated with Ag via
electrodeposition.

After successfully coating the Ta tubes, the second layer of material was added for both
cases; a Ag layer was applied to the Au coated tube and Au was added to the Ag coated tube.
While it was possible to achieve a bi-layer coating for both combinations, the Ta-Au-Ag
combination showed the best adherence and uniformity. An example of the bi-layer tube is
shown in Figure 69.

Figure 69: Bi-layer Ag-Au tube produced by separately plating Au then Ag onto a Ta base
tube.

157

The bi-layer tube was homogenized in a flowing Ar atmosphere for 4 hours at 900 °C.
While the two layers began to diffuse, post heat treatment SEM imaging revealed that the Ag
and Au layers did not maintain the conformal shape of a tube layer and began to expose the Ta
tube as the metals diffused (see Figure 70). Furthermore, the Ta tube became brittle after the
heat treatment.

Figure 70: SEM image of AgAu coated Ta tube after homogenization heat treatment. The
image shows that the AuAg layers diffused, however the AgAu alloy began to pull away from
the Ta tube base and was no longer a conformal coating.
158

Appendix B: Summary of Nanoporous Gold Tube Samples

Sample
Name
Base
Material
Wire
Outer
Diameter
Current
Density
(A/cm^2)
Duty
Cycle
Forward
Current
Reverse
Current
Time Solution Visual Result
TJ-012 Au 500 μm 0.02 DC -3.14 mA - N/A
1 M
AgNO3
Anode
passivated
TJ-012 Au 500 μm 0.02 on/rev/off -3.14 mA
0.785
mA
1 hour
1 M
AgNO3
Poor adhesion
TJ-013 Au 500 μm 0.02 DC -3.14 mA - 30 min
1 M
AgNO3
Thin coating
TJ-014 Au 500 μm 0.02 DC -3.14 mA - 30 min
1 M
AgNO3
Thin coating
TJ-015 Au 500 μm 0.02 DC -3.14 mA - 30 min
1 M
AgNO3
Thin coating
TJ-016 Au 500 μm 0.02 DC -3.14 mA - 30 min
1 M
AgNO3
Thin coating
TJ-020 Au 500 μm 0.02 DC -3.14 mA - 30 min
1 M
AgNO3
Thin Coating
TJ-
020-2
Au 500 μm 0.02 DC -3.14 mA - 30 min
1 M
AgNO3
Denstritic
Coating
TJ-021 Au 500 μm 0.02 DC -3.14 mA - 1 hour
1 M
AgNO3
Dendritic
Coating
TJ-022 Au 500 μm 0.02 DC -3.14 mA - 130 min
1 M
AgNO3
Dendritic
Coating
TJ-023 Au 500 μm 0.02 DC -3.14 mA - 90 min
1 M
AgNO3
Dendritic
Coating
TJ-024 Au 500 μm 0.02 DC -3.14 mA - 60 min
1 M
AgNO3
Dendritic
Coating
TJ-025 Au 500 μm 0.02 DC -3.14 mA - 60 min
1 M
AgNO3
Dendritic
Coating
TJ-026 Au 50 μm 0.02 DC -3.14 mA - 14 min
1 M
AgNO3
Poor adhesion
TJ-027 Au 50 μm 0.01 DC -1.57 mA - 60 min
1 M
AgNO3
Dentritic
Coating
TJ-028 Au 50 μm 0.01 DC -1.57 mA - 30 min
1 M
AgNO3
Dentritic
Coating
TJ-029 Au 50 μm 0.005 DC
-0.785
mA
- 30 min
1 M
AgNO3
Dentritic
Coating
TJ-058 Au 300 um 0.0216 DC
-2.036
mA
- 100 min
Caswell
Silver
Dark Coating
TJ-060 Au 300 um 0.0216 DC
-2.036
mA
- 240 min
Caswell
Silver
Acceptable
TJ-061 Au 300 um 0.0216 DC
-2.036
mA
- 195 min
Caswell
Silver
Acceptable
TJ-062 Au 300 um 0.0216 DC
-2.036
mA
- 192 min
Caswell
Silver
Acceptable
TJ-063 Au 300 um 0.0216 DC
-2.036
mA
- 192 min
Caswell
Silver
Acceptable
TJ-068 Au 300 um 0.0216 DC
-2.036
mA
- 240 min
Caswell
Silver
Acceptable
TJ-069 Au 300 um 0.0216 DC
-2.036
mA
- 240 min
Caswell
Silver
Acceptable
TJ-078 Au 300um 0.0108 DC -2.44 mA - 240 min
Caswell
Silver
Dark Coating
TJ-079 Au 300um 0.0108 DC
-1.832
mA
- 240 min
Caswell
Silver
Dark Coating
159

TJ-080 Au 300um 0.0108 DC
-1.832
mA
- 240 min
Caswell
Silver
Dark Coating
TJ-081 Au 300um 0.0108 DC
-2.024
mA
- 240 min
Caswell
Silver
Dark Coating
TJ-082 Au 300um 0.0108 DC
-2.024
mA
- 240 min
Caswell
Silver
Dark Coating
TJ-083 Au 300um 0.0216 DC
-2.024
mA
- 240 min
Caswell
Silver
Acceptable
TJ-084 Au 300um 0.0216 DC
-2.036
mA
- 240 min
Caswell
Silver
Acceptable
TJ-085 Au 300um 0.0216 DC
-2.036
mA
- 240 min
Caswell
Silver
Acceptable
TJ-086 Au 300um 0.0216 DC
-2.036
mA
- 240 min
Caswell
Silver
Acceptable
TJ-087 Au 300um
Voltage
Control
VC 750 mV - 240 min
Caswell
Silver
Poor adhesion
TJ-088 Au 300um 0.0216 DC
-2.036
mA
- 194 min
Caswell
Silver
Acceptable
TJ-089 Au 300um 0.0108 DC
-1.013
mA
- 195 min
Caswell
Silver
Acceptable
TJ-090 Au 300um 0.0108 DC
-0.900
mA
- 192 min
Caswell
Silver
Acceptable
TJ-091 Au 300um 0.0108 DC
-0.672
mA
- 195 min
Caswell
Silver
Acceptable
TJ-092 Au 300um 0.0108 DC
-1.182
mA
- 180 min
Caswell
Silver
Acceptable
TJ-093 Au 300um 0.0108 DC
-1.018
mA
- 195 min
Caswell
Silver
Acceptable
TJ-094 Au 300um 0.0108 DC
-1.271
mA
- 195 min
Caswell
Silver
Acceptable
TJ-095 Au 300um 0.0108 DC -1.12 mA - 195 min
Caswell
Silver
Destroyed
TJ-096 Au 300um 0.0108 DC -1.12 mA - 195 min
Caswell
Silver
Acceptable
TJ-097 Au 300um 0.0108 DC -1.73 mA - 195 min
Caswell
Silver
Acceptable
TJ-098 Au 300um 0.0108 DC
-1.112
mA
- 210 min
Caswell
Silver
Acceptable
TJ-099 Au 300um 0.0108 DC
-1.112
mA
- 195 min
Caswell
Silver
Acceptable
TJ-100 Au 300um 0.0108 DC
-1.112
mA
- 195 min
Caswell
Silver
Acceptable
TJ-101 Au 300um 0.0108 DC
-1.112
mA
- 195 min
Caswell
Silver
Acceptable
TJ-102 Au 300um 0.0108 DC
-1.112
mA
- 195 min
Caswell
Silver
Acceptable
TJ-103 Au 300um 0.0108 90%
-1.112
mA
1.112
mA
243 min
Caswell
Silver
Acceptable
TJ-104 Au 300um 0.0108 DC -1.12 mA - 195 min
Caswell
Silver
Acceptable
TJ-105 Au 300um 0.0108 DC -1.12 mA - 195 min
Caswell
Silver
Acceptable
TJ-106 Au 300um 0.0108 85% -1.12 mA 1.12 mA
253.5
min
Caswell
Silver
Acceptable
TJ-107 Au 300um 0.0108 85% -1.12 mA 1.12 mA
274.5
min
Caswell
Silver
Acceptable
TJ-108 Au 300um 0.0108 85% -1.12 mA 1.12 mA
274.5
min
Caswell
Silver
Acceptable
160

Note: Current densities reflected in this section and Chapter 5 are updated from the
publication in Advanced Engineering Materials.
TJ-109 Au 300um 0.0108 85% -1.2 mA 1.2 mA
249.6
min
Caswell
Silver
Acceptable
TJ-110 Au 300 um 0.0108 85% -1.12 mA 1.12 mA
249.6
min
Caswell
Silver
Acceptable
TJ-111 Au 300 um 0.0108 85% -1.17 mA 1.17 mA
249.6
min
Caswell
Silver
Acceptable
TJ-112 Au 300 um 0.0108 85%
-1.068
mA
1.068
mA
249.6
min
Caswell
Silver
Acceptable
TJ-113 Au 300 um 0.0108 85%
-1.068
mA
1.068
mA
249.6
min
Caswell
Silver
Acceptable
TJ-114 Au 300 um 0.0108 85%
-0.864
mA
0.864
mA
249.6
min
Caswell
Silver
Acceptable
TJ-115 Au 300 um 0.0108 85%
-1.068
mA
1.068
mA
249.6
min
Caswell
Silver
Acceptable
TJ-116 Au 300 um 0.0108 85%
-1.068
mA
1.068
mA
249.6
min
Caswell
Silver
Acceptable
TJ-117 Au 300 um 0.0108 85%
-1.018
mA
1.018
mA
249.6
min
Caswell
Silver
Acceptable
TJ-118 Au 300 um 0.0108 85%
-1.094
mA
1.094
mA
249.6
min
Caswell
Silver
Acceptable
TJ-119 Au 300 um 0.0108 85%
-1.017
mA
1.017
mA
249.6
min
Caswell
Silver
Acceptable
TJ-120 Au 300 um 0.0108 85%
-0.9665
mA
0.9665
mA
246 min
Caswell
Silver
Acceptable
161

Appendix C: Summary of Coated Tetrahedral Truss Samples

Aluminum Samples
Sample
Name Substrate
Sputtering
Target Power
Working
Pressure Time Set-up Sputtering Rate
Al #11 1" glass Al-#1, 3" 65 W 4 mTorr 15 min
4" Distance, Stationary,
Rotational Holder 9.058 nm/min
Al #12 1" glass Al-#1, 3" 65 W 4 mTorr 12 min
4" Distance, Stationary,
Rotational Holder 8.9 nm/min
Al #13, Al S1 Structures Al-#1, 3" 65 W 4 mTorr 12 min
4" Distance, Stationary,
Rotational Holder 8.9 nm/min
Al #14, Al S2 Structures Al-#1, 3" 65 W 4 mTorr 12 min
4" Distance, Stationary,
Rotational Holder 8.9 nm/min
Al #15, Al S3 Structures Al-#1, 3" 65 W 4 mTorr 12 min
4" Distance, Stationary,
Cooling Stage, pos 1 8.9 nm/min
Al #16, Al S4 Structures Al-#1, 3" 65 W 4 mTorr 12 min
4" Distance, Stationary,
Cooling Stage, pos 2 8.9 nm/min
Al #17, Al S5 Structures Al-#1, 3" 65 W 4 mTorr 12 min
4" Distance, Stationary,
Cooling Stage, pos 1 8.9 nm/min
Al #18, Al S6 Structures Al-#1, 3" 65 W 4 mTorr 12 min
4" Distance, Stationary,
Cooling Stage, pos 2 8.9 nm/min
Al #19 Si (100) wafer Al-#1, 3" 65 W 4 mTorr 25 min
4" Distance, Stationary,
Cooling Stage, pos 1 14.08 nm/min
Al #20 Si (100) wafer Al-#1, 3" 65 W 4 mTorr 25 min
4" Distance, Stationary,
Cooling Stage, pos 2 14.08 nm/min
Al #21 Si (100) wafer Al-#1, 3" 65 W 4 mTorr 90 min
4" Distance, Stationary,
Cooling Stage, pos 1 13.5 nm/min
Al #22 Si (100) wafer Al-#1, 3" 65 W 4 mTorr 90 min
4" Distance, Stationary,
Cooling Stage, pos 2 13.5 nm/min
Al S7 Structures Al-#1, 3" 65 W 4 mTorr 155 s
4" Distance, Stationary,
Cooling Stage, pos 1 13.5 nm/min
Al S7-pos 2 1" glass Al-#1, 3" 65 W 4 mTorr 155 s
4" Distance, Stationary,
Cooling Stage, pos 2 13.5 nm/min
Al S7 Structures Al-#1, 3" 65 W 4 mTorr 155 s
4" Distance, Stationary,
Cooling Stage, pos 1 13.5 nm/min
Al S7-pos 2 1" glass Al-#1, 3" 65 W 4 mTorr 155 s
4" Distance, Stationary,
Cooling Stage, pos 2 13.5 nm/min
Al #23 1" glass Al-#1, 3" 65 W 4 mTorr 90 min
4" Distance, Stationary,
Cooling Stage, pos 1 13.5 nm/min
Al #24 1" glass Al-#1, 3" 65 W 4 mTorr 90 min
4" Distance, Stationary,
Cooling Stage, pos 2 13.5 nm/min
Al S8 Structures Al-#1, 3" 65 W 4 mTorr 5 min
4" Distance, Stationary,
Cooling Stage, pos 1 13.5 nm/min
Al S8-pos 2 1" glass Al-#1, 3" 65 W 4 mTorr 5 min
4" Distance, Stationary,
Cooling Stage, pos 2 13.5 nm/min
Al #25 1" glass Al-#1, 3" 65 W 4 mTorr 40 min
4" Distance, Rotating,
pos 1 (top) 4.25 nm/min
Al #26 1" glass Al-#1, 3" 65 W 4 mTorr 40 min
4" Distance, Rotating,
pos 2 (bottom) 4.25 nm/min
Al S9 structures Al-#1, 3" 65 W 4 mTorr 40 min
4" Distance, Rotating,
pos 1 (top) 4.25 nm/min
Al S9-pos 2 1" glass Al-#1, 3" 65 W 4 mTorr 40 min
4" Distance, Rotating,
pos 2 (bottom) 4.25 nm/min
Al S10 Structures Al-#1, 3" 65 W 4 mTorr 40 min
4" Distance, Rotating,
pos 1 (top) 4.25 nm/min
Al S10-pos 2 1" glass Al-#1, 3" 65 W 4 mTorr 40 min
4" Distance, Rotating,
pos 2 (bottom) 4.25 nm/min
162

Inconel 600 Samples

Sample Name Substrate
Sputtering
Target Power
Working
Pressure Time Set-up
Sputtering
Rate
Inconel #7 pos 1 Si (100) wafer
Inconel 600-
3" #1 65 W 4 mtorr 10 min
4" Distance, Stationary,
Cooling Stage, pos 1

Inconel #7 pos 2 Si (100) wafer
Inconel 600-
3" #1 65 W 4 mtorr 10 min
4" Distance, Stationary,
Cooling Stage, pos 2

Inconel #8 pos 1 Si (100) wafer
Inconel 600-
3" #1 65 W 4 mtorr 10 min
4" Distance, Stationary,
Cooling Stage, pos 1 11.4 nm/min
Inconel #8 pos 2 Si (100) wafer
Inconel 600-
3" #1 65 W 4 mtorr 10 min
4" Distance, Stationary,
Cooling Stage, pos 2 14.8 nm/min
Inconel #9 pos 1,
Inconel S1 Structures
Inconel 600-
3" #1 65 W 4 mtorr 3 min
4" Distance, Stationary,
Cooling Stage, pos 1 11.4 nm/min
Inconel #9 pos 2 1" glass
Inconel 600-
3" #1 65 W 4 mtorr 3 min
4" Distance, Stationary,
Cooling Stage, pos 2 11.4 nm/min
Inconel #10 Si (100) wafer
Inconel 600-
3" #1 65 W 4 mtorr 80 min
4" Distance, Stationary,
Cooling Stage, pos 1 11.75 nm/min
Inconel #11 Si (100) wafer
Inconel 600-
3" #1 65 W 4 mtorr 80 min
4" Distance, Stationary,
Cooling Stage, pos 2 11.75 nm/min
Inconel S2 Structures
Inconel 600-
3" #1 65 W 4 mtorr 3 min
4" Distance, Stationary,
Cooling Stage, pos 1 11.75 nm/min
Inconel S2 pos 2 1" glass
Inconel 600-
3" #1 65 W 4 mtorr 3 min
4" Distance, Stationary,
Cooling Stage, pos 2 11.75 nm/min
Inconel #12 1" glass
Inconel 600-
3" #1 65 W 4 mtorr 77 min
4" Distance, Rotating,
pos 1 (top) 3.8 nm/min
Inconel #13 1" glass
Inconel 600-
3" #1 65 W 4 mtorr 77 min
4" Distance, Rotating,
pos 2 (bottom) 3.8 nm/min
Inconel S3 Structures
Inconel 600-
3" #1 65 W 4 mtorr 80 min
4" Distance, Rotating,
pos 1 (top) 3.8 nm/min
Inconel S3 pos 2 1" glass
Inconel 600-
3" #1 65 W 4 mtorr 80 min
4" Distance, Rotating,
pos 2 (bottom) 3.8 nm/min
Inconel S4 Structures
Inconel 600-
3" #1 65 W 4 mtorr 80 min
4" Distance, Rotating,
pos 1 (top) 3.8 nm/min
Inconel S4 pos 2 1" glass
Inconel 600-
3" #1 65 W 4 mtorr 80 min
4" Distance, Rotating,
pos 2 (bottom) 3.8 nm/min
Inconel S5 Structures
Inconel 600-
3" #1 65 W 4 mtorr 8.5 min
4" Distance, Stationary,
Cooling Stage, pos 1 11 nm/min
Inconel S5 pos 2 1" glass
Inconel 600-
3" #1 65 W 4 mtorr 8.5 min
4" Distance, Stationary,
Cooling Stage, pos 2 11 nm/min
Inconel S6 Structures
Inconel 600-
3" #1 65 W 4 mtorr 8.5 min
4" Distance, Stationary,
Cooling Stage, pos 1 11 nm/min
Inconel S6 pos 2 1" glass
Inconel 600-
3" #1 65 W 4 mtorr 8.5 min
4" Distance, Stationary,
Cooling Stage, pos 2 11 nm/min

163

Ti 6Al-4V Samples

Sample Name Substrate
Sputtering
Target Power
Working
Pressure Time Set-up
Sputtering
Rate
Ti #1 Si (100) wafer Ti-6Al-4V, 3" #1 65 W 4 mtorr 10 min
4" Distance, Stationary,
Cooling Stage, pos 1
Ti #2 Si (100) wafer Ti-6Al-4V, 3" #1 65 W 4 mtorr 10 min
4" Distance, Stationary,
Cooling Stage, pos 2
Ti #3 Si (100) wafer Ti-6Al-4V, 3" #1 65 W 4 mtorr 40 min
4" Distance, Stationary,
Cooling Stage, pos 1 5.125 nm/min
Ti #4 Si (100) wafer Ti-6Al-4V, 3" #1 65 W 4 mtorr 40 min
4" Distance, Stationary,
Cooling Stage, pos 2 5.125 nm/min
Ti #5 Si (100) wafer Ti-6Al-4V, 3" #1 65 W 4 mtorr 230 min
4" Distance, Stationary,
Cooling Stage, pos 1 5.125 nm/min
Ti #6 Si (100) wafer Ti-6Al-4V, 3" #1 65 W 4 mtorr 230 min
4" Distance, Stationary,
Cooling Stage, pos 2 5.125 nm/min
Ti # 7 Si (100) wafer Ti-6Al-4V, 3" #1 65 W 4 mtorr 235 min
4" Distance, Stationary,
Cooling Stage, pos 1 5.125 nm/min
Ti #8 Si (100) wafer Ti-6Al-4V, 3" #1 65 W 4 mtorr 235 min
4" Distance, Stationary,
Cooling Stage, pos 2 5.125 nm/min
Ti-S1 Structures Ti-6Al-4V, 3" #1 65 W 4 mtorr 6 min
4" Distance, Stationary,
Cooling Stage, pos 1 5.125 nm/min
Ti-S1-Pos 2 1" Glass Ti-6Al-4V, 3" #1 65 W 4 mtorr 6 min
4" Distance, Stationary,
Cooling Stage, pos 2 5.125 nm/min
Ti-S2 Structures Ti-6Al-4V, 3" #1 65 W 4 mtorr 6 min
4" Distance, Stationary,
Cooling Stage, pos 1 5.125 nm/min
Ti-S3 1" Glass Ti-6Al-4V, 3" #1 65 W 4 mtorr 6 min
4" Distance, Stationary,
Cooling Stage, pos 2 5.125 nm/min

Abstract (if available)

Abstract Lightweight hierarchical materials, which have tunable structural features on multiple length scales, have been a topic of increased research interest due to advances in techniques for fabricating these cellular solids with nano and microscale features. The study of these materials extends to a variety of applications including catalysis, optics, and batteries. However, processing and mechanical stability remain an integral part of expanding and scaling up the overall size of hierarchical structures with nano and microscale features at the base level of structural hierarchy. ❧ In this study, fabrication methods, characterization, and mechanical testing of hierarchical materials is explored on different length scales. The initial focus of this work explores a process called dealloying to develop hierarchical nanoporous (np) metal structures with nanopores on the smallest feature length scale and a complex shape on the macroscale. Using a step-by-step fabrication method, a binary alloy tube was developed and subsequently dealloyed to produce a np metal tube. These tubes were generated as a proof of concept exercise so that similar step-by-step procedures could be applied to more complex shapes. ❧ In addition to templating and dealloying as a method to produce hierarchical cellular materials, microlattice metal-polymer composites were generated using a combination of 3D direct laser writing (3D-DLW) and metal coating via magnetron sputtering. Using magnetron sputtering, metal coatings of Aluminum, Inconel 600, and Ti 6Al-4V were deposited onto microlattice structures. The cross section of the composite microlattice materials were evaluated using microtome, a microscopy preparation method where the structure is embedded in epoxy and thinly sliced, which allows for the assessment of cross sections of individual lattice struts. The inclusion of alloy coatings is significant because traditional deposition methods for metals are limited to a few select elements. Therefore, using sputtering as a deposition method has the potential to expand the achievable properties of composite microstruss structures. Uniaxial compression of the samples revealed that the compressive strength can be influenced by the metal that is deposited. Furthermore, by applying metal coatings, the deformation mode of the composite structures is shifted from node-dominated failure towards an increase in ligament failure. ❧ The final aspect of this dissertation discusses the scaling up of microcellular polymer materials produced with 3D-DLW, which has a limited print volume. In order to increase the size of the printed structures, multiple printed areas are “stitched” together to form a larger structure. The influence of the stitching along this interface on the mechanical behavior of cellular structures was explored. The tensile behavior of the samples was compared to continuously printed samples with no stitch. A custom micromechanical setup was used to complete tensile testing of the structures, and no pronounced tensile strength difference between stitched and continuously printed samples was found, although stitched samples regularly failed along the stitch interface. ❧ Combined, these results have addressed three critical aspects of the emerging field of hierarchical materials, including: 1) developing novel methods to incorporate nanoscale features in hierarchical materials

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Asset Metadata

Creator Juarez, Theresa (author)

Core Title Development and characterization of hierarchical cellular structures

School Viterbi School of Engineering

Degree Doctor of Philosophy

Degree Program Mechanical Engineering

Publication Date 09/29/2017

Defense Date 08/14/2017

Publisher University of Southern California (original), University of Southern California. Libraries (digital)

Tag dealloying,hierarchical materials,microtrusses,nanoporous metals,OAI-PMH Harvest,sputtering

Language English

Contributor Electronically uploaded by the author (provenance)

Advisor Hodge, Andrea (committee chair), Bermejo-Moreno, Ivan (committee member), Gupta, Satyandra (committee member), Kassner, Michael (committee member)

Creator Email theresa.juarez10@gmail.com

Permanent Link (DOI) https://doi.org/10.25549/usctheses-c40-440736

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Legacy Identifier etd-JuarezTher-5798.pdf

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Document Type Dissertation

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Type texts

Source University of Southern California (contributing entity), University of Southern California Dissertations and Theses (collection)

Access Conditions The author retains rights to his/her dissertation, thesis or other graduate work according to U.S. copyright law. Electronic access is being provided by the USC Libraries in agreement with the a...

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