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Understanding the formation and evolution of boundaries and interfaces in nanostructured metallic alloys

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Understanding the formation and evolution of boundaries and interfaces in nanostructured metallic alloys

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Content Copyright 2020 Joel Antunez Bahena

Understanding the formation and evolution of boundaries and interfaces
in nanostructured metallic alloys

by

Joel Antunez Bahena

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

August 2020

ii
Acknowledgements
First, I would like to thank my advisor, Professor Andrea Hodge, for her unwavering guidance and
mentorship throughout this journey. I am grateful for her ability to bring the best out of me as an
individual, a Hispanic in STEM, and as a scientist.
I am also grateful to the staff and faculty that have helped me navigate my time at USC. Kim Klotz,
thank you for our semi-weekly talks and for always being willing to share a sugary treat with me.
Thank you Professor Lessa Grunenfelder for showing me what amazing teaching looks like and
allowing me to flourish as a TA. I would also like to acknowledge the staff at the Core Center of
Excellence in Nano Imaging (CNI) for their assistance and for taking the time to answer my
questions. I would also like to thank Professor Paulo Branicio and Professor Mitul Luhar for
serving on my dissertation committee as well as their guidance throughout the years.
The studies presented in this thesis would not have been completed without the efforts from my
distinguished collaborators. From the Karlsruhe Institute of Technology, I would like to thank Dr.
Torben Boll and Dr. Reda Chellali for their tremendous Atom Probe Tomography (APT) work and
for making me feel welcome in Germany. I am also grateful for my time spent at Sandia National
Labs working with Dr. Brad Boyce, Dr. Khalid Hatter, and Dr. Christopher Barr. As a result of
their mentorship, I am now versed in several advanced characterization techniques and more
importantly, have grown as a scientist. Lastly, a special thanks to my previous lab mates and
mentors, Dr. Theresa Juarez, Dr. Leonardo Velasco, Dr. Nathan Heckman, and Dr. Sebastián
Riaño for their scientific contributions and providing me with the tools to navigate my PhD.
I want to thank the Frank H. & Eva Buck Foundation for always encouraging me to continue my
education and gifting me the financial flexibility to pursue my goals. I also want to acknowledge
the funding from the Office of Naval Research (N00014-18-1-2617) and the National Science
Foundation (DMR-1709771 and OISE-1460006) for making this research possible.
My time at USC was enhanced by the friends I have developed throughout the years. Thank you
to the Hodge Group members for all the memories and scientific discoveries we have made along
the way. Working with all of you has been a pleasure and I wish you the best in your future
endeavors. To USC SHPE, thank you for providing me with a community and the continued
support throughout my PhD. I especially want to thank Dr. Chelsea Appleget for going through
this journey with me starting day one. Your friendship is invaluable and have made these past five
years truly memorable.
I want to thank my family for all their love and support. Without all of you, I would not have had
the courage to purse or continue my PhD. You are all my inspiration and have shown me the
importance of hard work and dedication. Thank you for always encouraging me to pursue my
dreams. To the best mom in the world, Juana Antunez, you have always believed in me and I
would not be the person I am today without your support and sacrifices.

iii
Table of Contents

Acknowledgements ____________________________________________________________ ii
List of Figures ________________________________________________________________ v
Abstract ____________________________________________________________________ xii
1 Introduction ______________________________________________________________ 1
2 Background ______________________________________________________________ 2
2.1 Nanostructured materials _______________________________________________ 2
2.2 Interfaces in nanostructured materials _____________________________________ 6
2.2.1 Grain boundaries __________________________________________________ 7
2.2.2 Interphase boundaries _____________________________________________ 10
2.3 Tailoring interfaces ___________________________________________________ 13
2.3.1 Grain boundary engineering (GBE) __________________________________ 13
2.3.2 Nanotwinned metals ______________________________________________ 16
2.4 Grain boundary segregation engineering (GBSE) ___________________________ 21
2.4.1 Thermodynamic approach _________________________________________ 24
2.4.2 Kinetic approach _________________________________________________ 26
2.4.3 Nanometallic multilayers (NMMs) ___________________________________ 29
2.5 Thermal Transformations in Nanostructured materials _______________________ 32
2.5.1 Recovery _______________________________________________________ 33
2.5.2 Recrystallization _________________________________________________ 34
2.5.3 Grain Growth ___________________________________________________ 37
2.6 Heterogenous nanostructured materials ___________________________________ 39
3 Experimental Techniques __________________________________________________ 44
3.1 Magnetron sputtering of nanostructured metallic alloys ______________________ 44
3.2 Characterization methods ______________________________________________ 50
3.2.1 Differential scanning calorimetry (DSC) ______________________________ 50
3.2.2 Scanning electron microscopy (SEM) and electron back scattering diffraction
(EBSD) 51
3.2.3 Focused ion beam (FIB) ___________________________________________ 56

iv
3.2.4 Transmission electron microscopy (TEM) _____________________________ 58
3.2.5 Atom probe tomography (APT) _____________________________________ 63
4 GBE – Effect of initial twinned microstructure on the thermal evolution of highly NT alloys
67
4.1 Synthesis of NT Cu Alloys _____________________________________________ 68
4.2 Microstructural evolution of tailored microstructures ________________________ 71
4.3 Contributions from stored energy ________________________________________ 73
4.4 Mechanisms controlling grain growth behavior _____________________________ 76
4.5 Summary ___________________________________________________________ 77
5 GBSE – An investigation of the of the Mo-Au system ___________________________ 79
5.1 As-sputtered microstructure ____________________________________________ 81
5.2 Multilayer degradation region, T= 350˚C __________________________________ 86
5.3 Recrystallization region, T=550˚C _______________________________________ 88
5.4 Grain growth region, T= 800˚C _________________________________________ 92
5.5 Grain growth region limit, T=1000˚C _____________________________________ 96
5.6 Summary ___________________________________________________________ 99
6 Manipulation of interfaces to produce a heterogenous nanostructured superalloy _____ 101
6.1 As-Sputtered Microstructure ___________________________________________ 105
6.2 Three hours ________________________________________________________ 110
6.3 Five and Eight Hour Heat Treatments: Heterogeneous Nanostructured Material __ 112
6.4 AR Region: Five Hour Treatment _______________________________________ 115
6.5 AR Region: Eight Hour Treatment ______________________________________ 121
6.6 Summary __________________________________________________________ 124
7 General conclusions and future research outlook _______________________________ 127
References _________________________________________________________________ 132
Appendix A: Additional APT characterization of annealed NMMs ____________________ 145
Appendix B: Complementary characterization of sputtered Inconel 725 _________________ 151
Appendix C: Summary of sputtered samples ______________________________________ 156

v
List of Figures
Figure 1: Hardness plotted as a function of grain size for Ni–W alloys, where the dash lined
indicates the classical trend for Hall-Petch strengthening. [4] ............................................. 3
Figure 2: TEM micrographs of nanostructured Ni samples synthesized through a) ball milling, b)
electrodeposition, and c) sputtering. [6] ................................................................................ 4
Figure 3: TEM micrographs of the grown observed in nanocrystalline Ni samples annealed at
420˚C at different time intervals. [9] ...................................................................................... 5
Figure 4. Schematic illustration of defects driven strengthening mechanisms. [13] ..................... 6
Figure 5. Schematic of a HAGB and LAGB, showing the structural difference between the two
type of grain boundaries. [11] ................................................................................................ 8
Figure 6: Schematic of grain 1(blue) and grain 2(black) extending over one another, showing
the points of coincidence. ........................................................................................................ 8
Figure 7. Schematic of the different type of interphase boundaries a) incoherent, b) semi-
coherent, and c) coherent. [11] ............................................................................................ 10
Figure 8. TEM micrograph of an Al-Pb nanometallic multilayer displaying the alternating
layered structure and interphase boundaries. [25] .............................................................. 11
Figure 9. TEM micorgraph of an Inconel 718 alloy heat-treated to form the g” precipitate phase.
[28] ....................................................................................................................................... 12
Figure 10: Plot illustrating the effects of misorientation angle on several properties listed on the
right. [22] ............................................................................................................................. 14
Figure 11: A molecular dynamics simulation of grain boundary energy for Ni, Cu, and Al. [36]
............................................................................................................................................... 15
Figure 12: Top surface SEM of 304 stainless (a) as received and (b) thermomechanically
processed. [37] ..................................................................................................................... 16
Figure 13: Cross-sectional TEM of a sputtered Cu alloy sample with a nanotwinned
microstructure. [41] .............................................................................................................. 17
Figure 14: Stress-strain curves of several nanowtinned Cu samples with different twin
thicknesses. [46] ................................................................................................................... 18
Figure 15: Top-surface EBSD scans of UFG Cu (right) and nt Cu (left) at temperatures from
room-temperature to 400˚C. [47] ......................................................................................... 19

vi
Figure 16: Cross sectional grain boundary maps of the NT-Cu annealed at 300˚C for a) 1 Hr, b)
2 hrs, c) 4 hrs, and d) 24 hrs [49] ......................................................................................... 21
Figure 17: TEM and representative APT map of alloy 617 showing B segregation and quantified
in the 1D concentration profile across the B enriched region. [56] .................................... 22
Figure 18: Plot of impact energy versue annealing time at 450°C of an Fe alloy, showig
increased toughness as the annealing time is increased. [34] ............................................. 23
Figure 19: (a) Grain size of ball-milled W and a W-Ti alloy. The TEM micrographs of (b) as-
milled structure and of (c) W and (d) W-Ti after week heat-treatment ay 1100˚C. [58] ...... 25
Figure 20: Thermal stability maps, indicating alloy pairs that have a proclivity for nano grain
stabilty(green) or instability (red). [60] ............................................................................... 26
Figure 21: Illustration of the Zener pinning mechanisms, showing the a) attractive and b)
repulsive force between a particle and a grain boundary. [61] ........................................... 27
Figure 22: TEM micrographs at different magnification showing the microstructure of a ball-
milled Cu-Ta sample heat treated at 900°C and the presence of Ta clusters. [63] ............. 28
Figure 23: Plot of hardness as a function of layer thickness for several NMMs. [73] ................ 30
Figure 24. Top and cross sectional SEM of a Cu-W, demonstrating the breakdown of the layered
structure at elevated temperatures. [81] .............................................................................. 31
Figure 25. Cross sectional TEM of a Hf-Ti NMMs layers being annealed and showing how the
microstructure evolves at different temperatures. [83] ........................................................ 32
Figure 26: Stress-strain curve of a coarse-grained Cu sample and nanocrystalline Cu sample
highlighting the differences in ductility. [85] ....................................................................... 33
Figure 27: Micrograph of an etched Al surface identify the migration of grain A through strain
induced boundary migration (SIBM). [86] ........................................................................... 35
Figure 28: SEM images of a Ni superalloy strained at a) 0%, b) 1.8%, c) 4.7%, and d) 11% after
1 min heat treatments at 1150˚C. [88] .................................................................................. 36
Figure 29: The microstructural evolution of a W-Cr NMM in the after being annealed at a)
550˚C, b) 800˚C and c) 1000˚C. [91] ................................................................................... 37
Figure 30: In-situ TEM of a Ni films exhibiting abnormal grain growth. Arrows point to the
formation of a twin (I) and abnormally large grains (II & III) [97] .................................... 38
Figure 31: a) Schematic and b-h) microstructural features of a gradient nanotwinned structure
[102]. .................................................................................................................................... 40

vii
Figure 32: Schematic illustrating the potential mechanical properties of heterogenous
nanostructures (GNGs) in comparison to coarse-grained and nanograined materials [101].
............................................................................................................................................... 41
Figure 33: STEM of an Ta-Hf sample in the as-sputtered condition and after being annealed at
550˚C and 1000˚C [104]. ..................................................................................................... 42
Figure 34: a) Schematic of indent performed on the top surface of a Ni-superalloy and SEM of
the b) top surface, c) cross-section, and d) corresponding cross-sectional EBSD. [89] ..... 43
Figure 35: Schematic of magnetron sputtering chamber, with a one source arrangement. ........ 45
Figure 36: The influence of deposition rate on the critical radius for a twin nucleus and a perfect
nuclei on a stainless steel sample. [107] .............................................................................. 46
Figure 37: Twin boundaries in high stacking fault energy materials a) Al, b) Al-Mg, and c) Ni
which were synthesized via magnetron sputtering. [43] ...................................................... 47
Figure 38: Schematic of a magnetron sputtering chamber, with a two-source arrangement. A
technique useful for the deposition of NMMs. ...................................................................... 48
Figure 39: Cross-sectional TEM of a Hf-Ti NMM synthesize via magnetron sputtering. [82] ... 49
Figure 40: Residual stress profile of a sputtered Cu sample using a FIB technique. [109] ........ 50
Figure 41: DSC scan of a Hf-Ti showing thermal events from 20-1000˚C. [82] ......................... 51
Figure 42: EBSD scan of the a) cross-section of an electrodeposited Cu alloy and b) top surface
of a coarse-grained sample. The insets shows the associated inverse pole figure color key.
[112] ..................................................................................................................................... 52
Figure 43: Cross sectional t-EBSD and TEM of a nanostructured Cu-alloy. [114] ................... 53
Figure 44: GOS maps of an Fe-1%Si sample in the a) as deformed state and after annealing at
787˚C for b) 15, c) 30, d) 120, and e) 180 min. [115]. ......................................................... 55
Figure 45: EBSD maps of twin related domains observed in a Ni sample, showing the a) IPF
map, b) GOS map, and c) boundary character. [120] ......................................................... 56
Figure 46: SEM images of a sample being prepared by FIB lift-out method (a-f), being welded
onto an APT post (g-j) and milled into a needle shape (k-n). [122] ..................................... 58
Figure 47: Illustration of different TEM modes a) bright field, b) dark field, c) and diffraction
mode, with corresponding micrographs of a Hf-Ti multilayer below each illustration (d-f)
[124] ..................................................................................................................................... 60

viii
Figure 48: STEM EDS maps of a nanostructured Hf-Ti sample depicting a) Hf b) Ti, c) the
representative HAADF image and d) composition line scan across region A-B. [72] ........ 62
Figure 49: ACOM-TEM maps of a Pd sample a) of the cross-section and b) plan view of a
sputtered Pd sample. C) and d) show the respective quality maps of each scan. [128] ....... 63
Figure 50: Illustration of a pulsed laser atom probe, where a femtosecond laser causes the
evaporation of individual ions that get collected by a 2d X-Y detector. ............................... 64
Figure 51: a) APT reconstruction of a ball-milled CuNb system, where Nb rich clusters are
highlighted in the b) proximity histogram c) applied isosurface reconstruction, and d) 2d
density plots. [68] ................................................................................................................. 65
Figure 52: Cross-sectional bright field TEM comparing as-sputtered a) Cu-Al (λ=5), b) Cu-Al
(λ=18), c) Cu-Ni (λ=13), and d) Cu-Ni (λ=31); and E) representative XRD scans of the NT
alloys. The arrow indicates film growth direction. ............................................................... 70
Figure 53: Top surface EBSD IPF maps showing grain evolution of as-sputtered A) Cu2Al-A, B)
Cu2Al-B, C) Cu10Ni-A, and D) Cu10Ni-B after heat treatments performed at 200°C,
400°C, and 600°C. Dashed boxes highlight instances of significant grain growth, note the
changes in scale bar. The IPF triangle is shown to the right of the scans. .......................... 72
Figure 54: Top surface EBSD maps showing the GOS distribution for samples annealed at
200˚C and 600˚C. Σ3 grain boundaries are outlined in red. ................................................ 75
Figure 55: Mo-Au NNMs (a) bright field TEM with inset SAED patterns (top), dark field TEM
(middle), grain size distribution (bottom), and (b) differential scanning calorimetry scan of
the Mo-Au multilayer from 20°C to 1000ºC. From the scan, thermal events were indicated
and enthalpy changes (ΔH) were calculated. ....................................................................... 82
Figure 56: Cross-sectional bright field TEM with insets of respective SAED patterns (top),
corresponding EDS maps (middle), and grain size distribution (bottom) of Mo-Au samples
heat-treated at (a) 350ºC, (b) 550ºC, (c) 800ºC, and (d) 1000ºC. ........................................ 84
Figure 57: Integrated radial Intensity profiles interpolated from SAED patterns obtained from
Mo-Au samples (a) as-sputtered condition and heat treated at (b) 350ºC, (c) 550ºC, (d)
800ºC and (e) 1000ºC. .......................................................................................................... 85
Figure 58: a) APT map from the Mo-Au sample annealed at 350˚C showing the multilayered
structure with a 17 at.% Au isosurface. The 1D concentration profile (b) obtained from the
region indicated by the dashed arrow shows the composition across several layers. ......... 87

ix
Figure 59: APT maps from the Mo-Au sample annealed at 550˚C showing solute behavior at
GBs. The first APT map (a.1) shows a Au enriched GB (segregation zone) with insets
showing individual Au and Mo ion distribution. The 1D concentration profile (b.1) is
obtained through the region indicated by the dashed arrow, where the inset shows
concentration of (C + O). The second APT map (a.2) shows clusters along a GB with insets
showing individual Au and Mo ion distribution. The 1D concentration profile of (b.1) is
obtained through the region indicated by the dashed arrow, where the inset shows the Au
content through “Cluster 2” obtained from the region indicated by the blue cylinder. ...... 90
Figure 60: APT maps from the Mo-Au sample annealed at 800˚C, where (a) shows the ion
distribution of Mo (30 at.% isosurface) and Au (90 at.% isosurface), and (b) shows the ion
distribution of C (15 at.% isosurface) and Si (10 at.% isosurface). Regions are separated
where possible micro-fractures have occurred during APT measurements. These interfaces
are not evaluated. The 1D concentration profile (c) obtained from the region indicated by
the dashed arrow spans across a Au precipitate and a Mo-rich grain. ............................... 94
Figure 61: EDS scans of a Au precipitate and Mo-rich grain and the corresponding cross-
sectional bright field TEM image of the Mo-Au sample heat treated at 1000˚C. The red
arrows indicate the respective regions where the scans were taken from. ........................... 98
Figure 62: (A) BF-STEM image showing the nanoscale columnar structure of the as-sputtered
Inconel 725 sample, with (B) bright field TEM images highlighting density of nanotwins.
(C) EDS maps of corresponding STEM image in (A) of Ni, Cr, Ti, Nb, (D) grain width
distribution of columnar grains, and (E) ACOM map showing a strong (111) texture. ..... 106
Figure 63: (A) Schematic of a representative cross-section of the as-sputtered film highlighting
the region expected to undergo abnormal recrystallization delineated by blue lines. Growth
direction is shown by arrow to the left. This region includes a green rectangle representing
(B) and a red rectangle representing (C-F). Where (B) shows the GOS map and
corresponding distribution plots, which reveals a gradient in GOS. (C-F) highlight the
EBSD scans (top) and BF-STEM images (bottom) representing the microstructural
evolution of the abnormal recrystallization region in (C) the as-sputtered condition and
after aging treatments of (D) 3 hours, (E) 5 hours, and (F) 8 hours. IPF triangle is shown to
the bottom right of the scans. .............................................................................................. 109

x
Figure 64: (A) BF-STEM image of the 3-hour aged sample of the upper portion of the film and
(B) corresponding EDS maps highlighting Cr segregation where white scale bars are 1µm.
(C) ACOM scans of representative region, where green insets highlight select GBs with red
dotted lines indicating HAGBs and white dotted lines indicating LAGBs and CSL GBs. Cr
precipitates are encircled with yellow dotted lines. IPF triangle is shown to the right of the
scan. .................................................................................................................................... 112
Figure 65 (A) EBSD maps of cross-section of heterogenous structured achieved at aging times of
5 and 8-hours (top) with respective grain size distributions of grains in the AR region (ii)
(bottom). Insets highlights characteristic microstructures of the region (i). ...................... 115
Figure 66: (A) BF-STEM images and (B) corresponding ACOM scans of the AR region in the 5-
hour aged sample, highlighting ALGs governed by different thermal processes. Randomly
oriented recrystallized grains are highlighted by dashed black and white lines. White scale
bar is 1 µm and IPF triangle is shown to the right of the scan. ......................................... 117
Figure 67: (A) BF-STEM image of the 5-hour aged sample illustrating the heterogenous
precipitate and microstructural behavior in the AR region with inset SAED pattern
confirming the presence of δ precipitates. (B) EDS maps of Ni, Cr, Ti, and Nb highlighting
location of different precipitates where white scale bars are 1 µm. ................................... 120
Figure 68: (A) BF-STEM Image, corresponding (B) EDS maps of Ni, Cr, Nb, and (C) ACOM
maps from the outlined (orange) region in (A) with inset of IPF triangle of the 8-hour aged
sample. White scale bars are 500 nm. ................................................................................ 123
Figure 69: Schematic overview of the microstructural evolution of the sputtered Inconel 725 film
at different aging times that presents the changes in grain morphology, precipitate
formation, and development of a heterogenous microstructure. The arrow to the left
indicates the growth direction of the film and the red dotted box highlights the AR region.
............................................................................................................................................. 126
Figure 70: APT maps of a Mo-Au sample annealed at 800˚C, highlighting the phase separation
occurring of Mo (red) and Au (yellow). Top left to right: Combined APT maps of Mo and
Au, APT map of Mo, and APT map of Au. Below is a concentration plot extracted from the
ROI indicated by blue line. ................................................................................................. 145

xi
Figure 71: APT map (left) of a Hf-Ti NMM sample annealed at 500˚C, highlights the multilayer
structure and segregation of Hf (black) along interfaces. 1D concentration profile (right)
along the z axis displays the elemental fluctuations. .......................................................... 147
Figure 72: APT map (left) of a Hf-Ti NMM sample annealed at 500˚C shows a top view
perspective of Figure 71, where 1D concentration profiles (rights) were extrapolated from
different grain boundaries, showing an Hf enrichment. ..................................................... 147
Figure 73: APT map (left) of a Ta-Hf NMM sample annealed at 550˚C showing a preserved
multilayer structure of Ta (purple) and Hf (yellow), and is further confirmed by the
concentration profile (right) ............................................................................................... 148
Figure 74: Rotated view of the APT map from Figure 73 of a Ta-Hf NMM sample annealed at
550˚C which highlights the breakdown of the multilayer structure and shows the formations
of HF rich regions. .............................................................................................................. 149
Figure 75: APT map (left) of a Ta-Hf NMM sample annealed at 1000˚C shows the a non-
multilayer structure that has phase-separated as indicated by the sharp transitions in the
concentrations profile (right). ............................................................................................. 150
Figure 76: Cross-section of APT map from Figure 75 depicting the formation of Hf clusters
forming in the Ta rich regions. ........................................................................................... 150
Figure 77: The heat-treatment process for conventional Inconel 725. For sputtered films, solid
solution step was skipped, and samples were annealed at 3, 5 and 8 hours. ..................... 152
Figure 78: XRD data of samples annealed at 730˚C at various aging times, collected on a
Rigaku Ultima IV diffractometer. The boxes highlight structural changes. ....................... 153
Figure 79: TEM snapshots of the beginning and end of an in-situ heating experiment performed
on sputtered Inconel 725 along with ACOM showing some of the grains exhibiting a texture
change. ................................................................................................................................ 154
Figure 80: (Left) SEM image of cross-section of as-sputtered Inconel 725 with the various
indentations caused by nanoindentation. (right) Hardness profile from the top of the film
towards the bottom in microns. ........................................................................................... 155

xii
Abstract
Nanostructured alloys have many attractive mechanical, chemical, and thermal properties that are
largely attributed to a high density of interfaces. However, these same interfaces can drive
microstructural instability at relatively low temperatures, resulting in the annihilation of the
nanostructure and subsequent loss of associated properties. Thus, there is significant interest in
investigating routes to prevent or mitigate thermal processes such as grain growth or
recrystallization. To date, two routes have demonstrated some level of effectiveness: (i)
introducing a high density of low energy boundaries that resist grain boundary mobility and (ii)
utilizing solutes to decorate interface that restrict grain boundary motion through either kinetic or
thermodynamic mechanisms. Despite current progress, the thermal processes and mechanisms are
not well understood, particularly the intermediate microstructural transitions that can provide key
insights. Thus, this dissertation examines the thermal evolution of several tailored nanostructured
alloys in order to examine the effects of interfaces and chemical composition on thermal stability,
and also as a means to develop new nanocrystalline systems. Overall the results from these
investigations culminate in three important findings: 1) An interplay between kinetic and
thermodynamics is shown in a Mo-Au nanometallic multilayer system which dictates
microstructural transformations at different temperatures, 2) twin boundaries at high densities,
were shown to initiate the early onset of recrystallization and abnormal grain growth in
nanotwinned Cu alloys, however, abnormal grain growth can be mitigated by the presence of tilted
twins that can induce a random texture, and 3) stored energy gradients in a nanotwinned Ni
superalloy, in the form of defects and intrinsic stresses, can lead to the heterogenous activation of
thermal processes that results in a unique combination of grain size and morphology. Overall, this
dissertation contributes to the understanding of the thermal processes and mechanisms that

xiii
influence the evolution of interfaces and microstructures and provides insight into developing
stable and unique nanostructured materials.

1
1 Introduction
Nanostructured materials possess microstructural features in the nanometer regime (< 100 nm),
inherently creating a high density of interfaces that contribute to enhanced mechanical, chemical,
and thermal properties. However, a consequence of the increased number of interfaces is the
generation of a thermodynamic driving force that can prompt an unpredictable grain growth
behavior. This microstructural instability has limited the use of nanostructured materials in
commercial applications as most notable properties degrade with increasing feature size.
A potential route to stabilize nanostructured materials is through the design of alloys that contain
a high density of low energy or solute decorated interfaces, which can mitigate grain boundary
motion. Despite promising results, interface engineered materials have shown susceptibility to a
series of thermally activated transformations, leading to unique or abnormal microstructures at
elevated temperatures. The thermal processes and mechanisms facilitating these microstructural
transitions are not well understood and have garnered significant interest as it may elucidate the
role of interfaces on thermal stability. Furthermore, insights on thermal phenomena could be
utilized to tailor and develop advanced nanostructured alloys with optimized properties. Thus, to
expand the working space of nanostructured materials, a detailed and comprehensive
characterization of the thermal response of interface engineered nanostructured is required.
The present study investigates the microstructural evolution of a series of interface engineered
alloys at a wide range of temperatures. Specifically, nanometallic multilayers (NMMs) and
nanotwinned (NT) alloys synthesized by magnetron sputtering are presently examined as the
highly dense and tunable interfaces provide model systems for investigating the contribution of
interfaces to the activation of thermal phenomena. Complementary characterization techniques

2
detailing elemental segregation, crystallographic information, and grain boundary character at
transitional temperatures provide fundamental insights on the underlying processes governing the
thermal transformations of nanostructured alloys. Ultimately, the results from the comprehensive
study performed in this dissertation will enhance the design of nanostructured materials by
elucidating the role of interfaces and thermal processes and mechanisms in relation to
microstructural evolution.
2 Background
2.1 Nanostructured materials
Nanostructured materials have feature sizes that are on average less than 100 nm, which result in
a high density of interfaces that can make up a large volume fraction of the constituent
microstructure and dictate several structure-property relationships [1, 2]. For instance, the yield
strength of nanostructured materials has been shown to be an order of magnitude higher than
similarly composed coarse grained materials, which is primarily attributed to a higher number of
interfaces impeding dislocation motion [3]. The relationship between grain size and yield strength
can be expressed by the Hall-Petch equation, which is as follows:
𝜎
!
=𝜎+
"
√$
(1)
sy is the resulting yield strength, s is the yield strength of the bulk material, k is the strengthening
coefficient, and d is the average grain diameter. Several experimental results have validated the
expected trends of the Hall-Petch relationship. For example, electrodeposited Ni-W alloys
demonstrate this linear trend in strength as the grain size is decreased, which is illustrated in Figure
1 [4]. It should be noted that as the grain size reaches below 10 nm, the trends are not clear as other

3
mechanisms are expected to contribute. Experimental studies have shown both a plateau and a
decrease in strength for a grain size below 10 nm, and thus, is still an area of scientific investigation
[5].

Figure 1: Hardness plotted as a function of grain size for Ni–W alloys, where the dash lined
indicates the classical trend for Hall-Petch strengthening. [4]
As more studies have shown the expansive and attractive properties of nanostructured materials,
the techniques to synthesize nanostructured materials have become extensive. While the features
size for each synthesis technique can be comparable, each technique can generate materials with
deviations in microstructures (dislocation, impurities, vacancies, and etc.) that can directly
influence properties. Such microstructural differences can be observed in the synthesized Ni
samples shown in Figure 2, where three different techniques (ball-milling, electrodeposition, and
sputtering) were utilized [6]. Each synthesis route provides certain benefits and drawbacks, but
ultimately the range of techniques increases the potential for the use of nanostructured materials
in engineering applications. However, the main obstacle for widespread use of nanostructured
materials, regardless of synthesis technique, is an inherent thermal instability.

4

Figure 2: TEM micrographs of nanostructured Ni samples synthesized through a) ball milling, b)
electrodeposition, and c) sputtering. [6]
Despite recent progress, nanostructured materials suffer from the drawback of being
microstructurally unstable at elevated temperatures [7], where some materials have even
demonstrated instability at ambient temperatures [8]. While interfaces are a source for many
desirable properties, a large density of interfaces can also be detrimental since they carry an
energetic penalty that drives grain boundary mobility. This is especially a concern in
nanostructured materials where the volume fraction of interfaces can be as much as 50% for 5 nm
grains, 30% for 10 nm grains, and about 3% for 100 nm grains [1]. The magnitude of the effects
of microstructural and grain size changes are exemplified in the case of nanocrystalline Ni, where
significant growth is observed at 420˚C, even after short annealing times (Figure 3) [9]. It should
be mentioned that the melting temperature (Tm) of Ni is 1455˚C and thermal processes such as
grain growth and recrystallization are not typically observed below 0.5 Tm. While prominent in
many pure systems, nanostructured alloys have also shown unstable performance at elevated

5
temperatures [10]. Thus, thermally driven grain size instability limits the applications of the
majority of nanostructured materials as properties can degrade even at lower temperatures ranges.

Figure 3: TEM micrographs of the grown observed in nanocrystalline Ni samples annealed at
420˚C at different time intervals. [9]
A high density of interfaces is an intrinsic characteristic of nanostructured materials that gives rise
to exceptional properties yet contributes to an overall thermodynamically unstable microstructure.
While it may seem to be an unavoidable issue, the subsequent driving force for grain growth varies
greatly depending on interface type and structures [11, 12]. Furthermore, segregation of secondary
elements to interfaces has been shown to either reduce grain boundary energy or mitigate grain
boundary mobility [7]. There are several routes to potentially stabilizing microstructure through
manipulation of interfaces, but the mechanisms controlling such thermal stability requires further
investigation to make these materials viable in engineering applications.
as-deposited 30 s
1 h 11 h
1 s
120 h

6
2.2 Interfaces in nanostructured materials
Crystalline solid contains several types of intrinsic defects such as point defects (vacancies,
interstitial, etc.), line defects (dislocations), and interfacial defects (grain boundaries, interphase
boundaries, free surfaces, etc). Figure 4 highlights many of these defects and how they can impede
dislocation motion to strengthen a material [13]. Of the listed defects, interfaces are typically
regarded as the primary contribution to mechanical properties, due to their effectiveness as barriers
to dislocation slip [11]. Interfaces are present in both pure elements and alloyed systems, where
interfaces between grains of the same phase are commonly referred to as grain boundaries and
those between two different phases are referred to as interphase boundaries.

Figure 4. Schematic illustration of defects driven strengthening mechanisms. [13]
Aside from mechanical strengthening, interfaces are essential to the diffusion and segregation of
solute elements as they can act as high diffusivity paths, which can control the evolution of

7
microstructure and properties at higher temperatures [14]. Interfaces also play important roles in
thermally activated processes such as recrystallization and grain growth , as well as the nucleation
and formation of secondary phases [15]. Furthermore, the composition and structure of the
adjoining mediums can give rise to microstructures with a wide spectrum of energy states.
Accordingly, interfaces are essential to understanding material properties and behaviors, especially
as the grain size of engineering alloys is trending towards the nanoscale and the volume of
interfaces increases.
2.2.1 Grain boundaries
A grain boundary (GB) is defined as the interface between two grains in polycrystalline materials,
where grains are defined as the regions of atoms ordered in a certain crystallographic structure and
orientation [16]. GBs have five degrees of freedom (DOF); three which specify rotations between
adjacent grains, known as the misorientation angle, and two that characterize the boundary plane.
As a general approach, GBs are typically defined by the misorientation angle and depending on
the angle can be categorized as low angle grain boundaries (LAGB) or high angle grain boundaries
(HAGB) [16]. LAGB are defined as GBs with misorientation angles less than 15˚ and are
composed of arrays of dislocations. HAGB, on the other hand, have misorientation angles greater
than 15˚ and dislocations overlap causing them to no longer be distinguishable [17]. The structural
difference between these two types of GBs is illustrated in Figure 5. As the structure changes, there
are there are also energetic discrepancies that need to be considered [18]. For instance, LAGB have
improved lattice matching between grains leading to a relatively lower grain boundary energy than
HAGB. However, several HAGBs have specific geometric orientations that allow for low energy
configurations [19].

8

Figure 5. Schematic of a HAGB and LAGB, showing the structural difference between the two
type of grain boundaries. [11]
A basic descriptor geometric fit of HAGBs has been generated based on the concept of coincident
site lattice theory (CSL), where it is postulated that GBs with ordered structure and low atomic
misfit can exhibit lower energies. The basic of CSL theory centers around the shared atomic sites
between two adjacent lattices [16]. In Figure 6, two grain shown in black and blue are
superimposed on top of one another, and the shared points highlighted in orange are the
coincidence sites. The outbox, shown in black, represent the unit cell of the CSL.

Figure 6: Schematic of grain 1(blue) and grain 2(black) extending over one another, showing
the points of coincidence.

9
At certain misorientation angles, a higher density of coincidence sites is achieved leading to an
improved lattice matching between neighboring grains. Based on the density of coincident sites,
the parameter S, which classifies different CSL GBs, can be calculated as presented in Eq. 2.
Σ=
%&'( *+,%-. +/ 012
%&'( *+,%-. +/ 34!5(6, ,6(('3.
(2)
The lower the S value, the better the lattice fit between two neighboring grains, which would
ideally lead to a lower GB energy. For example, S1 GBs have near perfect crystal symmetry, and
are typically classified as LAGBs, exhibit the lowest energy since the misfit between grains is
accommodated by dislocations. S3 GBs, which are also known as twin boundaries, typically
exhibit the lowest energies of any HAGB [20]. The range of low-S GBs is typically characterized
by a criterion that can be expressed as:
∆𝜃
-67
=𝜃
°
/Σ
9&
(3)
The most widely used conditions for the expression are based on Brandon’s criterion, where θ◦=
15 and n = ½ [21]. Under these conditions, the conventional range for S GBs that exhibit low
energies and which are considered “special grain boundaries” is S ≤ 25. GBs that fall outside of
this range are termed random grain boundaries. While Brandon’s criterion is commonly used to
categorize special or low energy boundaries, several studies have suggested that not all low S value
boundaries exhibit special properties as CSL theory does not consider plane orientation [22]. A
theory considering all five DOF has yet to be established. However, of these boundaries, S3 twin
boundaries are consistently identified as having special properties [23]. Therefore, not only is the

10
density of interfaces important but also the types of GB present within a microstructure in
governing the overall properties.
2.2.2 Interphase boundaries
While the interfaces between pure element materials are of considerable interest, many engineering
alloys contain multiple phases which creates interfaces between dissimilar phases or composition.
These interphases are particularly prominent in metallic multilayers, where alternate layers of
compositions can play a critical role on the material properties depending on the coherency
between the two phases [24]. These three different categories, which are summarized in Figure 7,
are coherent, semi-coherent, or incoherent interfaces. Depending on the atomic matching, different
boundaries will have drastic ranges in energy. For coherent interfaces, where the lattice mismatch
is minimal, the energies range from 5-200 mJ m
−2
. On the other hand, semi-coherent (200-800 mJ
m
−2
) and incoherent interphases (800-2500 mJ m
−2
) demonstrate higher energies due to increases
atomic misfit [11].

Figure 7. Schematic of the different type of interphase boundaries a) incoherent, b) semi-
coherent, and c) coherent. [11]

11
Metallic multilayers and specifically those with nanoscale layers have generated interest due to the
high density of interphase that have shown improved thermal stability depending on coherency
between the different layers. For example, an Al-Pb multilayer structure synthesized by
accumulative roll bonding formed nanometer layers of Al and Pb that formed epitaxial semi-
coherent interphases, as seen in Figure 8. The semi-coherent interphase boundaries provided
stability and led to an elevated melting temperature for the Pb nanolayers [25]. The integrity of
these layers and the thermal stability of several other multilayer systems have been explored,
showing similar improvements when the alternating layers are either coherent or semi-coherent
[26, 27].

Figure 8. TEM micrograph of an Al-Pb nanometallic multilayer displaying the alternating
layered structure and interphase boundaries. [25]
The formation of secondary phases in many alloy and superalloy systems is another example of
the prominence of interphase and coherency. In nickel-based superalloys, such as Inconel 718, the
formation of Ni-based precipitates is a common processing mechanism that leads to significant

12
strengthening. The major strengthening phase in this specific alloy is the g” phase, which has a
BCT structure and is coherent with the g matrix. These intermetallic which are often observed in
the nanoscale (~10 to 20 nm) as illustrated in Figure 9, produces a significant source of
strengthening through coherency strain. Additionally, this strain can stabilize precipitates and
corresponding matrix that increases the thermal stability of the alloy [28].

Figure 9. TEM micorgraph of an Inconel 718 alloy heat-treated to form the g” precipitate phase.
[28]
Thus, in nanostructured alloys, the interface interactions become much more complex as both grain
boundaries and interphases are present. Both types of interfaces contribute to the mechanical and
thermal properties and act as fast diffusion paths for segregation. While there has been significant
work on understanding GBs of interphase behavior, tailoring of these interfaces though synthesis
techniques have not been as widely explored.

13
2.3 Tailoring interfaces
Interfaces have a wide spectrum of energies and properties, making the distribution and structure
of interfaces as important as the quantity [16, 29]. Thus, incorporating specific types of interfaces
into a microstructure has been of particular interest in recent years. Advancements in the synthesis
and processing of nanostructured materials has allowed for more control of microstructural
features such as grain size, impurities, and defects [30]. More importantly, current synthesis
techniques have shown promise in tailoring the structure and distribution of interfaces at the
nanoscale, allowing for the development of materials with combinations of properties not
previously achievable [31].
Two general approaches to engineering alloys with tailored interfaces have been explored. The
first approach involves producing microstructures with a high density of low-energy boundaries
that exhibit special properties and contribute to improving bulk properties [32, 33]. The second
approach utilizes segregation as a means to decorate interfaces with solutes, which can influence
the energy, mobility, and structure of the interfaces [34, 35]. Both methods provide nanoscale level
control of the interfaces and have demonstrated potential in generating more thermally stable
nanostructures along with other interesting properties, that will be discussed in more detail in the
following sections.
2.3.1 Grain boundary engineering (GBE)
The concept of grain boundary engineering (GBE) was first proposed in 1980 by Wattanabe with
the premise of producing high performance materials by tuning interfaces and more specifically,
grain boundaries [22]. Specifically, the intent of GBE is to fabricate materials with a high fraction

14
of low energy or “special” GBs. As previously mentioned, boundaries with sigma values S< 27
are considered to have special characteristics and therefore, increasing the density of these
boundaries should result in improved properties. A summary of these improved properties is listed
in Figure 10.

Figure 10: Plot illustrating the effects of misorientation angle on several properties listed on the
right. [22]
The reason for these unique properties directly results from the lowering of the overall GB energy
by introducing more CSL boundaries and reducing the number of randomly oriented grain
boundaries. The plot in Figure 11 shows the simulated grain boundary energy for FCC metals of
Al, Ni, and Cu at different misorientation angles, where when the CSL criteria is fulfilled, drops
in energy are observed [36].

15

Figure 11: A molecular dynamics simulation of grain boundary energy for Ni, Cu, and Al. [36]
A notable property in GBE microstructures, is the improved intergranular corrosion behavior.
Several alloys, while stable at room temperature, solute atoms are known to segregate to the GBs
at elevated temperatures. As a result, the change in chemical composition along the interfaces can
lead to preferential pathways for corrosion under certain environments. This is demonstrated for
304 stainless steel, where sensitization led to carbide precipitation that induced susceptibility to
intergranular corrosion (Figure 12) [37]. However, since GBE materials have an increased fraction
of special GBs that exhibit lower energies, the tendency for solute atoms to preferentially diffuse
to these GBs is reduced. Thus, when thermomechanically processed to increase the fraction of
special GBs, a more corrosion resistant was observed as seen in Figure 12. While the fraction of
special GBs plays a significant role in the corrosion behavior, the network of GBs is also crucial.
A continuous network of special GBs is more likely to arrest the progression of intergranular
corrosion [37, 38].

16

Figure 12: Top surface SEM of 304 stainless (a) as received and (b) thermomechanically
processed. [37]
Even between the different CSL boundaries, differences in energy exist. Thus, from a GBE
standpoint, certain boundaries would be preferred over others. The introduction of only the lowest
energy GBs would potentially lead to the most improved properties [12, 39]. Further discussion
can be found in refs [40] and [22].
2.3.2 Nanotwinned metals
As previously mentioned S3 twin boundaries are considered to have the most unique properties of
all grain boundary types, and recently there has been a surge in designing materials with a high
density of twin boundaries with nanoscale spacing [23]. Figure 13 depicts the typical
microstructure of a nanotwinned (NT) material, where the TEM micrograph of a Cu-4wt.% Al
alloy shows several twin boundaries (TBs), highlighted by yellow lines, that have a spacing of 9.7

17
nm. A TB defines the interface between the matrix lattice and mirror lattice where the stacking
sequence reverses. To further confirm the presence of twin boundary, the inset SAED pattern
shows the hexagonal pattern indicative of a twinned structure [41].

Figure 13: Cross-sectional TEM of a sputtered Cu alloy sample with a nanotwinned
microstructure. [41]
The formation of twinned or nanotwinned structures occurs is primarily observed in FCC metals
with a low stacking-fault energy (SFE), as a stacking fault is more likely to occur which is
conducive for the formation of a twin boundary. Thus, metals such as Cu, Ag, and stainless steel
(SFE 20-40 mJ/m
2
) [42] have a higher probability of forming twins than Ni or Al (SFE 125-166
mJ/m
2
), which have a relatively high SFE [43]. However, there are means to alter the SFE of
materials or to induce twins in moderate to high SFE materials. It has been shown that alloying Cu
with Al or Zn can drastically reduce the SFE [44], it also should be noted that whether an alloying
element will increase or decrease the SFE is still not thoroughly understood. Additionally, many
non-equilibrium synthesis processes have also been shown to promote twins in high SFE materials

18
[43, 45]. Thus, the working space for NT metals is currently being expanded, leading to an
increased knowledge of low-energy interface engineered nanostructures materials. Still, many of
the studies have focused on pure FCC metals and the benefits of these GBE microstructures are
discussed below.
NT materials have demonstrated improved mechanical properties primarily that have been
attributed to strain hardening generated by dislocation-dislocation and dislocation-TB interactions
[13]. A study on fully NT pure Cu samples highlighted the impact of a NT structure and the twin
thickness on the mechanical properties. These differences are illustrated in the stress-strain curves
in Figure 14, where a reduction in twin thickness up to 15 nm reveals an increase in yield strength,
and a Hall-Petch relationship with twin thickness is observed. Beyond 15 nm, strength decreases
but improvements in ductility are seen [46].

Figure 14: Stress-strain curves of several nanowtinned Cu samples with different twin
thicknesses. [46]

19
Beyond mechanical properties, the low energy state of twin boundaries results in metals that are
more thermally stable than a nanograined counterpart. A clear example is shown in in a comparison
study between a magnetron sputtered NT and UFG Cu sample. The top surface of each
microstructure characterized by EBSD from the as-sputtered condition to 400˚C. The EBSD scans
can been see in in Figure 15, where the initial grain size is 328 nm and 697 nm for the UFG and
NT sample, respectively. Upon annealing at 200˚C, 300˚C, and 400˚C, the UFG sample exhibits
grain size increases of 963 nm, 1.8 μm, and 4.3 μm. On the other hand, the NT samples showed
grain size stability at temperatures of 300˚C, with a slight increase to 750 nm. At 400˚C, an average
grain sizes of 1.2 μm was observed with a wide distribution [47].

Figure 15: Top-surface EBSD scans of UFG Cu (right) and nt Cu (left) at temperatures from
room-temperature to 400˚C. [47]

20
The thermal stability of the NT structure over its UFG counterpart was attributed to the high
density of low energy TBs which are less mobile than randomly oriented GBs. Furthermore, the
strong (111) texture and a higher density of triple junctions of the NT Cu sample was also expected
to contribute [47]. As both samples were produced through magnetron sputtering with high purity
targets, the role of impurities is expected to be minimal. It is interesting to note that while the UFG
Cu sample exhibited coarsening at a lower temperature (200˚C), the final grain size distribution
was more uniform. As can be seen in Fig. 15, several larger (100) grains, over 10 um in size,
surround the smaller (111) oriented grains in the NT sample annealed at 400˚C, indicating
abnormal grain growth behavior. Thus, there are still many questions surrounding the
predictability of the grain growth behavior of NT metals in addition to the role of nanotwins. In a
previous study by Misra and Zhang, a sputtered NT Cu sample demonstrated thermal stability up
to 800˚C [48], which is in stark contrast to the observations from the aforementioned study by
Zhao et al. In addition, other studies have shown much lower thermal stability in similarly
produced NT Cu samples as seen in Figure 16, suggesting that a high density of nanotwins is
insufficient to stabilize a nanostructured material [49, 50]. Nanotwin, in high densities, have also
been discussed as a source of excess energy that can promote abnormal grain growth and
recrystallization [51]. As previously discussed, different synthesis mechanisms can generate a
broad gamut of microstructures with different characteristics that can influence thermal stability
along with other properties. Thus, there a more comprehensive examination of the effect of the
initial twinned microstructure of NT metals is still necessary to understand the thermal processes
governing the grain growth behavior of these materials.

21

Figure 16: Cross sectional grain boundary maps of the NT-Cu annealed at 300˚C for a) 1 Hr, b)
2 hrs, c) 4 hrs, and d) 24 hrs [49]
2.4 Grain boundary segregation engineering (GBSE)
Interface segregation occurs when the necessary temperature and time for the diffusion of solute
elements is reached. This process, in many instances, can negatively impact the material properties
through mechanisms such as GB embrittlement, intergranular corrosion, and phase separation.
Preventing the diffusion of solute and impurities to interfaces is an active topic of research and is
a main proponent to the field of GBE. However, the process segregation to interfaces can be
tailored to manipulate interfaces and their properties, such as fracture resistance, grain boundary
energy and mobility [52-55]. For instance, GB segregation is a common technique to stabilize the
grain size of nanostructured alloyed systems, where solute atoms will segregate and enrich GBs or
form precipitates that provide thermodynamic or kinetic stabilization, respectively [7]. However,
the selection of alloying elements and the interface characteristics are crucial to obtaining these
improved properties. Several other studies have utilized similar segregation behavior as a tool to
design alloys with tuned and enhanced properties, where this method of interface manipulation has
recently been termed grain boundary segregation engineering (GBSE) [34]. Through

22
understanding the interface characteristics, GBSE can be used as a microstructural design method
to design materials with improved mechanical, thermal, and chemical properties.
The concept of GBSE has been applied to several common engineering materials which has
resulted in several improvements. For example, Alloy 617, a Ni-based superalloy, has shown a
25% increase in creep resistance at service temperatures of 700˚C through the addition of B, which
would segregate to GBs and interphases to enrich fine carbide precipitates. As seen in Figure 17,
the enrichment along GBs reached upwards of 20 at.%. The addition of B was shown to stabilize
fine coherent carbide precipitates and the overall grain sizes which promoted an improved creep
rupture strength [56].

Figure 17: TEM and representative APT map of alloy 617 showing B segregation and quantified
in the 1D concentration profile across the B enriched region. [56]

23
Another example demonstrating the benefits of GBSE is presented for an Fe-8.8wt.% Mn
quenched alloy. Annealing the alloy for over 100 hours produced an alloy that has improved
toughness (Figure 18), which allowed for sufficient time for the segregation of Mn that ultimately
reduced the GB energy. As a result, the grain size was stabilized that led to the improved toughness
and ductility [34].

Figure 18: Plot of impact energy versue annealing time at 450°C of an Fe alloy, showig
increased toughness as the annealing time is increased. [34]
Recent advancements in the field of GB segregation offers a unique opportunity to exploit GBs
and interfaces to develop alloys with unique and combined properties. Several more examples will
be discussed in the following section, however, emphasis will be placed on mechanisms to stabilize
the microstructure and grain size. As previously mentioned, many properties of nanostructured
materials stem from the nanograin size, so understanding the mechanisms that contribute to
thermal stability are critical. In the next sections, GBSE as a means to achieve thermal stability by
kinetic and thermodynamic mechanisms will be discussed.

24
2.4.1 Thermodynamic approach
The thermodynamic approach involves reducing the excess Gibbs free energy through solute GB
segregation until an equilibrium boundary concentration is reached. The reduction of GB energy
minimizes the driving force for grain growth and leads to the stabilization of grain size [57]. The
alloy selection for solute segregation is dependent on the enthalpy parameters, mainly the enthalpy
of segregation (Hseg) and enthalpy of mixing (Hmix). The most classical example of a
thermodynamic approach through solute segregation stabilizing is the comparison of a pure W and
W-20 at.% Ti sample (Figure 19) [58, 59]. Both samples were annealed under the same conditions
at 1100˚C for 1 week. The pure W sample saw an increase in grain size from 20 nm to 604 nm,
while the alloyed sample saw negligible grain coarsening with an increase from 22 nm to 24 nm
[59]. The main contribution to the observed thermal stability was Ti segregating to the GBs. EDS
maps in Fig. 10 highlight this segregation, where the concentration of solute Ti reached
concentrations of 50 at.%.

25

Figure 19: (a) Grain size of ball-milled W and a W-Ti alloy. The TEM micrographs of (b) as-
milled structure and of (c) W and (d) W-Ti after week heat-treatment ay 1100˚C. [58]
Certain alloys are expected to exhibit nanograin stability, where Murdoch and Schuh [60] have
constructed an expression to predict alloy combinations as follows:
∆;
!"#
(∆;
$%&
)
(
>𝑐 (4)
The constants a and c are calculated based on the critical temperatures for the selected alloy, and
∆Hseg and ∆Hmix represent enthalpy of segregation and enthalpy of mixing, respectively. From Eq.
4, thermal stability maps were produced that outlined alloy combinations as stable, metastable,
and unstable in terms of maintaining a nanostructure at elevated temperatures. In Figure 20, the
solvent elements are presented on the right through various symbols and the solute elements with

26
matching symbols are presented in the maps. Based on where the pairs fall within the maps, a
predication of the stability can be made.

Figure 20: Thermal stability maps, indicating alloy pairs that have a proclivity for nano grain
stabilty(green) or instability (red). [60]
These maps, along with other calculations and modeling works can be utilized to construct binary
alloys that should demonstrate thermal stability and a retention of the nanostructure properties.
2.4.2 Kinetic approach
In the kinetic approach, stabilization of the grain size arises from a reduction in GB mobility
through certain impeding mechanisms. The most typical kinetic mechanisms is Zener pinning,
where second-phase particles exert a pinning force on GBs [61], that is expressed in equation 5.
𝑃
>
=
?/@
A4
(5)

27
𝑃
>
in the equation is the Zener pinning pressure, 𝑓 is the volume fraction of precipitates, r is the
size of the particle, and g is the GB energy. Therefore, the smaller the particles and the greater the
volume fraction the more pressure that is applied to the GBs. The Zener pinning mechanisms
behaves in a two-step method, where initially, an attractive force between the particle and GB is
induced. Secondly, as the GB attempts to move past the particle, the particle generates a pinning
force that immobilized the GB (Figure 21) [62].

Figure 21: Illustration of the Zener pinning mechanisms, showing the a) attractive and b)
repulsive force between a particle and a grain boundary. [61]
Kinetic mechanisms have typically been associated with low-temperature stability of
microstructures, however recent studies have demonstrated that nanoscale precipitates can pin GBs
and retard growth even at higher temperatures. In a Cu-Ta ball-milled nanocrystalline sample
(Figure 22), the precipitation of nanoscale Ta clusters along the GBs generated a stabilized
microstructure that maintained a grain size of 167 nm even after annealing to 97% of the melting
temperature [63].

28

Figure 22: TEM micrographs at different magnification showing the microstructure of a ball-
milled Cu-Ta sample heat treated at 900°C and the presence of Ta clusters. [63]
The selection of alloys with tendencies to phase separate and form cluster or precipitates along
interfaces give way to advancements in GBSE. For example, a study of a Cu-Co immiscible sample
was forced into a solid-solution through high pressure torsion. After annealing, the phase
decomposition led to the formation of pure Co and Cu particles along and near GBs that are
expected to produce pinning pressure on the GBs and influence the thermal stability of the system.
Therefore, the segregation and formation of second phase precipitates in alloys that are expected
to phase separate can be utilized to tailor thermal or mechanical properties [64].
Asides from Zener pinning, other kinetic mechanisms can hinder grain mobility, such as vacancy
drag, chemical ordering, and solute drag. In alloys with small solute additions or impurities, solute

29
drag can have a significant impact on GB mobility, where as a GB migrates, it interacts and entraps
solute atoms that reside in the matrix. The solute atoms provide a drag force that reduce GB
velocity and generates sluggish growth kinetics. In a study performed on Al samples with varying
impurities, it was shown that the GB velocity of decreased from 7.4×10
−8
m
2
s
−1
to 1× 10
−10
m
2
s
−1
as the impurity concentration increase from 0.4 at. ppm to 7.7 at. ppm [65]. Along with several
other experiments [66, 67], this study shows the grain growth retardation and GB velocity decrease
as a function of solute atom concentration. As the solute atoms get entrapped, eventual equilibrium
boundary concentration will be reached and reduce the overall grain boundary energy. Therefore,
solute drag is typically seen as a precursor or a joint mechanism to GB energy reduction, which is
thermodynamically driven.
To date, most studies have focused in isolating either kinetic or thermodynamic mechanisms as
the primary driving force for thermal stability. However, it has also been shown that these two
mechanisms can occur simultaneously or consecutively [68, 69] and thus the intermediate steps
contributing to the overall stability are still unclear. Thermal processes such as recovery,
recrystallization and grain growth should also be examined in the context of thermal evolution in
nanostructured systems and could clarify and define their effects within the aforementioned
mechanisms [70, 71]. Therefore, a model system with fine control over the microstructure is
needed to explore the thermal and microstructural evolution of nanostructured materials.
2.4.3 Nanometallic multilayers (NMMs)
Nanometallic multilayers (NMMs) are ideal candidates for the investigation of GBSE
nanostructured materials due to the tunability of composition and structure of the microstructure
[72]. NMMs consist of alternating metallic multilayers on the order of several nanometers, where

30
the thickness of the layers directly influence several prominent physical properties. These materials
have garnered interest because of exceptional and controlled mechanical properties that have been
shown to be well beyond their bulk counterparts. The hardness of several NMMs samples with
varying compositions and layer thicknesses are presented in. Figure 23 that showcases a general
trend of increasing hardness with decreasing layer thickness [73]. The strengthening behavior seen
in NMMs resembles the Hall-Petch relationships, where layer thickness and grain size play similar
roles, respectively.

Figure 23: Plot of hardness as a function of layer thickness for several NMMs. [73]
In addition to the mechanical behavior of these materials, the nanolaminate geometry of NMMs
are promising materials to study thermally induced microstructural evolution. Their synthesis
allows for the control of the grain size, global composition, elemental distribution, and energetic
state by tuning of the deposition parameters and layer thicknesses [72]. Additionally, these
materials can by synthesized by techniques that minimize the high density of defects, dislocations,
and contaminants [68, 74-77]. Current thermal studies performed on NMMs have primarily

31
focused on the breakdown of the metallic layers by examining mechanisms such as thermal
grooving and layer pinch-off [78-80]. More recent work has investigated the microstructural
evolution of NMMs by annealing at elevated temperatures where distinct morphologies begin to
form [81, 82]. A study on the thermal evolution of Cu-W annealed at temperatures ranging from
400-800˚C revealed several microstructural transitions that led to the eventual breakdown of the
layers and the formation of a microstructure consisting of W particles embedded in a Cu matrix as
seen in Figure 24 [81].

Figure 24. Top and cross sectional SEM of a Cu-W, demonstrating the breakdown of the layered
structure at elevated temperatures. [81]
In another study, Hf-Ti NMMs were subjected to temperatures up to 1000˚C, where key structural
transformations were identified, and the system demonstrated a unique microstructural change
from a NMM to an equiaxed nanocrystalline structure [72, 82], which can be seen in Figure 25.
While the NMM loses the structural integrity of the nanolaminate structure, the microstructural
features, namely the grain size, remain on the nanometer scale at temperatures of 800˚C. The

32
stability that is achieved is attributed to the segregation behavior of Ti solute that enriches the GBs
surrounding the Hf-rich grains.

Figure 25. Cross sectional TEM of a Hf-Ti NMMs layers being annealed and showing how the
microstructure evolves at different temperatures. [83]
The thermal mechanisms and processes that induce these microstructural transformations in
NMMs are still not understood, thereby emphasizing the need to further explore the thermal and
microstructural evolution of these nanostructured materials.
2.5 Thermal Transformations in Nanostructured materials
While it has been mentioned that nanostructured materials are susceptible to grain growth in a
wide range of temperatures, it can occur before or after thermal processes such as recovery and
recrystallization. The driving forces for these processes is the reduction of excess energy, which
makes them prominent in nanostructured materials due to the high density of interfaces. Through
understanding these processes and potentially manipulating them by means of GBE and GBSE, it
is possible to tailor unique microstructures. While studies pertaining these processes in
nanostructured materials is sparse, a general summary including all length scales will be provided
in the following subsections.

33
2.5.1 Recovery
During recovery, thermal vibrations activate the motion of defects like vacancies and dislocations
that leads to rearrangement and partial annihilation of dislocations without significant grain
boundary mobility [84]. The imitation of recovery is highly dependent on the amount of strain
induced by defects, temperature, and material properties. Since nanostructured materials are often
produced via non-equilibrium synthesis techniques, these materials contain a high density of
defects compared to traditional coarse-grained materials, which has been shown to trigger recovery
at room temperatures. In combination with dislocation storage limitation due to grain size effect,
low temperature recovery can restrict the amount of strain-hardening exhibited by nanostructured
materials [85]. The difference in ductility is exemplified in the the stress-strain plot in Figure 26.

Figure 26: Stress-strain curve of a coarse-grained Cu sample and nanocrystalline Cu sample
highlighting the differences in ductility. [85]

34
The thermal processes in nanostructured materials can be convoluted. Grain coarsening has been
observed during the recovery stage that is facilitated by sub grain boundary migration and sub
grain rotation and coalescence. Additionally, reducing the defects during can provide sites for
solute atoms to segregate, especially in supersaturated solid solution alloys [71].
2.5.2 Recrystallization
The topic of recrystallization with regard to nanocrystalline materials is difficult to assess as there
is still much speculation to the mechanism involved. Typically, during recrystallization, strain-free
grains consume the recovered microstructure. However, this can occur through different processes.
In primary recrystallization, new grains appear as a function of nucleation and growth, where the
newly formed grains grow at the expense of the strained microstructure [86]. The second process
is referred to as strain induced boundary migration (SIBM) and is similar to recrystallization but
does not involve the formation of a new nuclei. Instead, grains with a stored energy advantage
consume neighboring grains with higher energy, leading to the formation of strain free regions
with similar textures [84]. An example of SIBM was first seen in a high purity aluminum sample
in 1950 by Beck and Sperry. Figure 27 shows the migration of grain A as it consumes grain B [86].
It should be noted that the driving force for recrystallization and SIBM is similar, but the
mechanism and resultant microstructure are quite different [87].

35

Figure 27: Micrograph of an etched Al surface identify the migration of grain A through strain
induced boundary migration (SIBM). [86]
When a microstructure is uniformly strained, either through processing or synthesis methods,
complete recrystallization is expected. However, there are instances where heterogenous
distribution of stored energy and residual stresses can lead to abnormal recrystallization and
unpredictable microstructures. Such behavior is typically observed under low strains or low stored
energy conditions as seen in Figure 28, where the recrystallization and formation of abnormally
large grains in an annealed Ni superalloy is more prominent when strained to 1.8% (Figure 28b)
and 4.7% (Figure 28c) [88]. Under these types of conditions, only a limited number of nuclei are
thermodynamically favorable to grow, which can then consume neighboring high stored energy
grains. The microstructure then stabilizes once abnormally large grains (low energy) impinge on
one another. Furthermore, there are situation where the distribution of stored energy is non-
uniform across a microstructure which can initiate distinct recrystallization processes in different

36
regions [88, 89]. In the situation where the stored energy conditions are higher, an increased
density of nuclei can form causing more grains to develop impinge on one another that can lead to
a relatively smaller grain size [90].

Figure 28: SEM images of a Ni superalloy strained at a) 0%, b) 1.8%, c) 4.7%, and d) 11% after
1 min heat treatments at 1150˚C. [88]
In addition, recrystallization has been shown to initiate several thermal mechanisms. For example,
during recrystallization of a W-Cr NMM, diffusivities along grain boundaries are expected to be
high enough to allow for the formation of diffusion zones where new phases can nucleate as seen
in Figure 29 at 1000˚C [91]. Similar behavior has been observed in other NMMs, and in most cases
segregation and the formation of secondary phases leads to nano-grain stability [82, 92].

37

Figure 29: The microstructural evolution of a W-Cr NMM in the after being annealed at a)
550˚C, b) 800˚C and c) 1000˚C. [91]
In summary, the mechanisms involved with recrystallization are complex, and could activate
stabilization mechanisms or promote coarsening. Thus, understanding the process is important for
designing stable nanostructures, especially in far from equilibrium microstructures as this area is
still relatively unexplored. Overall, controlling recrystallization is equally as important as
mitigating grain growth for tailoring thermally stable nanostructured materials.
2.5.3 Grain Growth
While grain growth typically follows recrystallization, in many nanostructures materials it has
often been observed to without recrystallization. Since the driving force for grain growth is two

38
orders of magnitude lower than the that of recrystallization, the excess energy generated by a high
density of grain boundaries can initiate grain growth before recrystallization [84]. The most
prominent form of grain growth seen in nanostructured materials is abnormal grain growth, where
select grains grow at the expense of surrounding smaller grains [93-95]. This grain growth
behavior leads to a bimodal grain size distribution (Figure 30) and highlights the unpredictable
nature of nanostructured materials at elevated temperatures [96, 97]. While the density of
interfaces contributes significantly to abnormal grain growth, it is speculated that the heterogenous
characteristics of nanostructured materials in the form of interface energies, chemistry, defects,
residual stresses, and texture also have significant impact [98, 99]. There has yet to be a consensus
identifying the key microstructural features controlling abnormal grain growth and remains an
open area of interest.

Figure 30: In-situ TEM of a Ni films exhibiting abnormal grain growth. Arrows point to the
formation of a twin (I) and abnormally large grains (II & III) [97]

39
It should be noted that the reports of abnormal grain growth may actually be a form of abnormal
recrystallization, and in recent years advanced characterization techniques have been able to detect
the difference between the two processes [88]. Still, both mechanisms are detrimental to the
material properties of nanostructured materials under thermal loads, and there is a need to further
understand the mechanisms controlling these processes.
2.6 Heterogenous nanostructured materials
It has often been observed that nanostructured materials will undergo some sort of abnormal
thermal transformation that lends to the detriment of thermal stability. Recently, tailoring these
drastic changes in grain size to generate materials with varying domains has been of significant
interest and has initiated a new class of materials known as heterogeneous nanostructured materials
(HNMs). More recent studies of HNMs have shown tremendous promise in improving strength-
ductility synergy, amongst other properties, which is lacking in most conventional nanostructured
metals. The lack of ductility in nanostructured materials stems from low strain hardening caused
by dislocations traversing the relatively small grains and annihilate into the surrounding interfaces
[100]. However, HNMs exhibit both hard (smaller grain size) and soft domains (larger grain size),
where the interfaces between these domains are able to accommodate geometrically necessary
dislocations (GNDs) generated during plastic deformation. An example of a HNM is seen in Figure
31, which displays an electrodeposited nanotwinned copper sample with gradients in both grain
size and twin thickness. The mechanical properties of certain HNMs have been assessed and a
general comparison of the strength-ductility synergy can be seen in Figure 32 .HNMs (labeled
GNGs) fall in a white space that had not been occupied by convention nanostructured or coarse-
grained materials [101].

40

Figure 31: a) Schematic and b-h) microstructural features of a gradient nanotwinned structure
[102].
Thus, the promise of HNMs has brought attention to the development of different processing and
synthesis methods to produce materials with controlled heterogeneity. However, many of the
current methods have several drawbacks which include limited grain morphology, topology, and
alloy selection [103].

41

Figure 32: Schematic illustrating the potential mechanical properties of heterogenous
nanostructures (GNGs) in comparison to coarse-grained and nanograined materials [101].
Specifically, insights from previous investigations of GBE and GBSE have become useful design
tools for producing heterogeneous nanostructured. For example, by annealing a Hf-Ti NMM to
1000˚C, a microstructure with Hf-rich grains surrounded by Ti grains arranged in two layers
equidistant from the center of the sample was produced. While the ability to localize precipitation
behavior in this instance is still unclear, it could be useful for designing specific types of HNMs
[82]. Furthermore, a Ta-Hf NMM film showed a bimodal distribution of HAGBs (top portion) and
LAGBs (bottom portion), which when annealed led to the activation of different thermal processes
and stabilization mechanisms resulting in a gradient nanostructure as seen in Figure 33 [104].
Despite these interesting findings, manipulating annealed nanostructured for the design of HNMs
is still relatively unexplored but developments in understanding interfaces and thermal process can
potentially unlock a new design space.

42

Figure 33: STEM of an Ta-Hf sample in the as-sputtered condition and after being annealed at
550˚C and 1000˚C [104].
A noteworthy approach would be to examine heterogeneous behavior observed in coarse grained
materials and translate it to the nanoscale. For instance, recent work of coarse-grained Ni
superalloys has shown an inclination to exhibit abnormal recrystallization in select regions
depending on stored energy conditions [89]. In Figure 34, a Ni-superalloy sample was indented
leading to a gradient in stored energy, which activated different thermal processes along the profile
of the material when heat-treated. Towards the center of the material abnormal recrystallization
was initiated, forming several abnormally large grains. Similar methods could be used in
nanostructured materials to achieve congruent grain size distributions. This is just one example of
a possible design route, but the potential in this field is expanse. In order to purposely engineering
inhomogeneities, a detailed understanding of these thermal processes and microstructural
transformations are necessary to develop effective means to promote the strength–ductility
synergy.

43

Figure 34: a) Schematic of indent performed on the top surface of a Ni-superalloy and SEM of
the b) top surface, c) cross-section, and d) corresponding cross-sectional EBSD. [89]

44
3 Experimental Techniques
3.1 Magnetron sputtering of nanostructured metallic alloys
Physical vapor deposition (PVD) is a common technique used to synthesize films ranging from a
few nanometers to several micrometers. The specific type of PVD utilized in this work is DC
magnetron sputtering, which provides the benefits of high sputtering rates (> 1 nm/s), an ability to
sputter an wide variety of metals or alloys, and the production of high purity films relative to other
synthesis techniques [105]. A schematic highlighting the magnetron sputtering chamber and its jet
components used in these studies is illustrated in Figure 35. Initially, a target material is placed on
the cathode and the chamber is then set to vacuum (< 10
-6
Torr) and filled with argon gas that
ranges from 1-5mTorr. The cathode is negatively biased with sufficient potential to generate an
Ar plasma composed of Ar
+
ions and e
-
. The positive Ar ions are then attracted to the negative bias
of the cathode leading to a bombardment of the target surface, which causes target atoms to be
ejected in the form of a metal vapor. In addition, the presence of a magnetron generates a magnetic
field that increases the density of e
-
, which enhances the bombardment process to produce high
sputtering rates. As a consequence of the PVD nature of this process, target atoms deposit onto the
substrate and the entire chamber layer by layer. This results in the formation of a film adhered to
a substrate, with physical properties including thickness, microstructure and morphology that can
be varied depending on sputtering time and rates [106, 107].

45

Figure 35: Schematic of magnetron sputtering chamber, with a one source arrangement.
A key benefit to magnetron sputtering is the capability to alter several processing parameters which
can directly influence microstructural characteristics. Mainly, this discussion will focus on the
formation of Σ3 twin boundaries and synthesis of materials with a high density of interphases, but
the capabilities of magnetron sputtering are much more expanse [105].
In regard to the formation of Σ3 twin boundaries, the non-equilibrium of process of magnetron
sputtering can increase the probability of developing twins, primarily as function of the deposition
rate, which has been seen for materials with a wide range of SFE’s [42, 43]. During the film growth
process, small islands of atoms nucleate, grow, and coalesce to comprise the resulting
microstructure, which have a probability of developing nuclei with SF defects or TBs. However,
depending on the SFE, the critical radius for a twin nucleus to form and grow is much larger than
that of a perfect nucleus. While unfavorable in equilibrium growth conditions, utilization of high
deposition rates has been shown to lower the critical radius of a twin nucleus, increasing the

46
probability of twin nucleation. Radii stability for twin and prefect grains is shown in Figure 36,
where it is seen that the difference in stable radii decreases with increasing deposition rate.

Figure 36: The influence of deposition rate on the critical radius for a twin nucleus and a perfect
nuclei on a stainless steel sample. [107]
Thus, synthesis by high deposition rates can achieve microstructural designs not previously
observed, such as NT microstructures in high SFE materials as seen for Al (~166 mJ/m
2
) and Ni
(~125 mJ/m
2
) in Figure 37 [43]. Furthermore, by controlling the deposition rate, the density of
twin boundaries or twin thickness can be tailored. Overall, changing sputtering parameters can
have a significant influence on other microstructural features in NT metals, such as grain width,
texture, and intrinsic stresses [42, 108].

47

Figure 37: Twin boundaries in high stacking fault energy materials a) Al, b) Al-Mg, and c) Ni
which were synthesized via magnetron sputtering. [43]
Another type of tailored nanostructure enabled by magnetron sputtering is the synthesis of NMM
systems, which provide tunability of composition and layer thickness [72]. The typical set-up for
producing NMM through sputtering is shown in Figure 38. The main difference from the
conventional set-up shown in Figure 35 is the presence of two cathodes, each with targets of
different composition that allows for deposition of alternating layers. By altering the on-time and
applied voltages of the of the individual power sources the thickness and microstructures of the
different layers can be manipulated. Moreover, power sources can be used in unison to co-deposit
targets to form layers with mixed compositions.

48

Figure 38: Schematic of a magnetron sputtering chamber, with a two-source arrangement. A
technique useful for the deposition of NMMs.
An example of a sputtered NMM is shown in Figure 39, where TEM micrographs of a Hf-Ti
NMMs displays the alternating layers of co-deposited Hf-rich layers and pure Ti layers. The
individual layers were measured to be 48 nm Hf-Ti and 5 nm Ti, where the shifts in contrast
indicate the changes in composition. This specific NMM exhibits a columnar structure that extend
throughout the individual layers, which is common for sputtered materials. Therefore, within each
columnar grain there are several smaller grains formed by intersection layers, which exhibit a
preferred texture.

49

Figure 39: Cross-sectional TEM of a Hf-Ti NMM synthesize via magnetron sputtering. [82]
An important aspect of sputtered microstructures is that films can exhibit gradients in intrinsic
residual stresses throughout the film thickness. During film growth, several microstructural
features such as grain size, texture, defects, and impurities can be altered and are highly dependent
on the sputtering conditions. Recently, high resolutions characterization techniques have been able
to detect the instantaneous stresses along the thickness of a film, which is shown for a Cu sputtering
film in Figure 40 [109]. These gradients can lead to heterogeneous microstructural evolutions
when annealed as has been seen in several sputtered systems [92, 104].

50

Figure 40: Residual stress profile of a sputtered Cu sample using a FIB technique. [109]
3.2 Characterization methods
3.2.1 Differential scanning calorimetry (DSC)
Differential scanning calorimetry (DSC) is a useful technique to determine key thermal transitions,
or thermally activated microstructural transformations in a material. A sample is heated with a
constant increase in temperature, and the changes in heat capacity is measured by tracking the heat
flow. Depending on the thermal transition, the heat flow will fluctuate, generating either
exothermic or endothermic peaks. Primarily, DSC can provide snapshots of thermal events and
determine critical temperatures of microstructural transformations [110]. For example, the DSC
scan in Figure 41 provides a roadmap of several thermal transitions such as recrystallization,
precipitation, and grain growth of a Hf-Ti nanostructure [82].

51

Figure 41: DSC scan of a Hf-Ti showing thermal events from 20-1000˚C. [82]
3.2.2 Scanning electron microscopy (SEM) and electron back scattering
diffraction (EBSD)
In scanning electron microscopy, a filament or field emission gun are utilized to generate an
electron beam that is condensed using multiple electromagnetic lenses and apertures into a beam
with a diameter ranging from 1-10 nm. As the beam rasters across the surface of a specimen, elastic
and inelastic interactions occur between the electrons and the sample. During an elastic interaction,
the energy of the electron is conserved, and the colliding electrons are reflected at high angles,
which are known as back-scattered electrons. When the interaction is inelastic, the incoming
electrons interact with the surface of the sample generating x-rays and secondary electrons. Based
on the type of interactions and emitted signal, different information regarding the sample can be
collected. Images from SEM are generated by the collection of secondary electrons and are
produced through the changes in contrast caused by interactions of electrons along different

52
depths. Elevated areas of the sample will appear brighter, as it can reflect a higher number of
electrons [111].
If combined with an electron backscatter diffraction (EBSD) detector, the GBs, grain orientations,
and grain sizes of a material can be determined. To perform EBSD, there is a traditional geometry
where the sample is tilted 70° to the incident electron beam. The interaction of the beam with
surface atoms causes the interacting electrons to scatter. Certain electrons will satisfy Bragg’s law
of a specific crystallographic plane, which will generate a Kikuchi diffraction pattern on the
phosphorous screen. From these patterns, the software can then determine the orientation and size
of specific grains. Different colors are associated with the different grain orientations as seen in
Figure 42, based on the orientation map in the inset of Figure 42b [112]. Specifically, Figure 42
shows the EBSD scan from the cross section of an electrodeposited Cu sample (Figure 42a) and
the top surface of a coarse-grained Cu sample (Figure 42b), highlighting the different grain
morphologies that can be achieved through different synthesis methods.

Figure 42: EBSD scan of the a) cross-section of an electrodeposited Cu alloy and b) top surface
of a coarse-grained sample. The insets shows the associated inverse pole figure color key. [112]

53
EBSD also allows for the characterization of GBs, which is specifically crucial for the
investigation of GBE materials. The GB character can be determined from the GB misorientation
angle, which only accounts for 3 of the 5 GB parameters. However, it is also possible to determine
the GB trace on the surface which can assist in differentiating coherent and incoherent S3 GBs.
The maximum achievable resolution for traditional EBSD is 20 nm, but this is only attainable
under ideal conditions [113]. To overcome these limitations, a technique known as transmission
electron back scattering diffraction (t-EBSD) was designed which can resolve nanoscale features
and grains. The geometric configuration and sample preparation is different than conventional
EBSD, as the sample must be electron transparent (< 100 nm) and placed in a different orientation
with respect to the electron beam [113]. The thinner specimen size reduces the interaction volume
and increases resolution. An example of a t-EBSD scan is displayed in Figure 43 of a NT Cu
sample with corresponding TEM image.

Figure 43: Cross sectional t-EBSD and TEM of a nanostructured Cu-alloy. [114]

54
In addition to orientation and grain size, long and short range misorientations along grains can be
determined from data collected from EBSD. These measurements, known as intragranular
orientation spreads have been particularly useful for determining variations in energy stored in
different grains which can govern many thermal processes. The average misorientation profile
across a grain can be reduced to a single average value such as grain orientation spread (GOS),
which averages the misorientation angle between each pixel to the average grain orientation [115].
It is estimated that stored energy in the form of dislocations or defects must be present to
accommodate the orientation spread. It should be noted that GOS is only a semi-qualitative
representation of stored energy, but studies have shown correlations between orientation spread
and stored energy [116]. Below, in Figure 44, a GOS map of a Fe-1%Si sample was annealed at
787˚C for times from 15 min to 180 min, highlighting the evolution of GOS. The orientation spread
ranges from 0.06˚ to 7.66˚, where grains closer to a red hue would indicate a larger orientation
spread and presumed stored energy. Those with a blue hue would have an orientation spread closer
to 0˚.

55

Figure 44: GOS maps of an Fe-1%Si sample in the a) as deformed state and after annealing at
787˚C for b) 15, c) 30, d) 120, and e) 180 min. [115].
Specifically, GOS maps have been central to many studies regarding abnormal grain growth or
abnormal recrystallization as grains with lower stored energies would have a competitive growth
advantage [89, 117, 118]. Furthermore, it has been a useful technique in examining partially
recrystallized microstructures, where recrystallized grains will typically exhibit GOS values below
1˚ [119]. In Figure 45, the IPF, GOS, and boundary character of a thermomechanically treated and

56
annealed Ni sample are highlighted. Several larger grains with a high density of twin boundaries
were observed and identified as twin related domains. The low GOS value and the networks of
twin boundaries point to certain portion of the microstructure undergoing recrystallization. Such
studies have provided valuable insight into the processing mechanisms controlling grain boundary
engineering.

Figure 45: EBSD maps of twin related domains observed in a Ni sample, showing the a) IPF
map, b) GOS map, and c) boundary character. [120]
3.2.3 Focused ion beam (FIB)
In order to image cross-sections of samples or to prepare specimens for other characterization
techniques, focused ion beam (FIB) microscopy is a valuable tool. The FIB works by generating
Ga ions through field emission, which interact and erode the specimen surface. This interaction
generates the emission of back-scattered ions, radiation, and ion induced secondary electrons.
Similar, to SEM, FIB utilizes secondary electrons to generate images that exhibit improved

57
contrast as compared to SEM. The enhanced contrast can highlight grains with different
crystallographic orientations. Moreover, the interaction of Ga with the surface generates sputtered
ions that can be utilized to mill features into a specimen [111].
When combined with a micromanipulator and gas injection system, the FIB can produce samples
with dimensions suitable for transmission electron microscopy (TEM) and atom probe tomography
(APT). For TEM, sample lamellae need to be produced with thicknesses less than 100 nm so
electrons can transmit through the specimen [121, 122]. To accomplish this, the sample is milled
from a region of interest and placed on TEM lift-out grids, where milling continues to further
reduce the thickness of the specimen. A similar technique can be used to prepare APT specimens,
where the initial steps require the lift-out samples to be placed onto specific APT post. Samples
are then thinned by a dual beam FIB in a confocal geometry to obtain sub 100 nm radial tips. The
steps for this type of procedure can be seen in Figure 46. In this study, samples for TEM and APT
analysis were prepared using a JIB-4500 multi-beam FIB.

58

Figure 46: SEM images of a sample being prepared by FIB lift-out method (a-f), being welded
onto an APT post (g-j) and milled into a needle shape (k-n). [122]
3.2.4 Transmission electron microscopy (TEM)
In transmission electron microscopy, an accelerated parallel beam of electrons transmitted through
thin (< 100 nm) samples allows for the observation of microstructures with nanometer size
features. Electrons are generated by a field emission gun, which are then condensed through a

59
series of lenses and transmitted through the specimen. The transmitted electrons are refocused and
then magnified by a series of electromagnetic lenses after passing through the specimen. Following
this, the electron are then projected onto a phosphor screen or a charged couple device (CCD)
camera to obtain the resulting image [123].
Depending on the energy and type of collision of the transmitted electrons, different
microstructural information can be gathered. The resulting energy of the electron is dependent on
whether the interaction between the electrons and specimen is elastic or inelastic. In an elastic
interaction, the energy of the transmitted electrons is conserved, while an inelastic interaction
involves the loss of energy due to scattering processes. It is possible to collect either elastic of
inelastic electrons by altering the position of the objective aperture to either above or below the
sample [123].Thus, different TEM modes can be achieved depending on position of the objective
aperture and the electron signal collected as illustrated in Figure 47. The three modes presented
are bright field (Figure 47a), dark-field (Figure 47b), and selected area electron diffraction (Figure
47c), and the resulting images of each mode for a Hf-Ti multilayer are seen in Figure 47d-f.

60

Figure 47: Illustration of different TEM modes a) bright field, b) dark field, c) and diffraction
mode, with corresponding micrographs of a Hf-Ti multilayer below each illustration (d-f) [124]
The bright-field mode, where the objective aperture is centered in the back focal plane of the lens,
is the conventional TEM mode. In this mode, the objective aperture blocks all diffracted electrons
allowing only for elastic electrons to contribute to the final image. Contrast in bright-field mode
results from local variations in the microstructure. In dark-field mode, the objective aperture is
positioned to block all but a selected diffracted beam [123]. As a result, only grains that diffract
strongly in the selected direction will be visible in the image, making this technique suitable for
studying grain size. Furthermore, diffracted electrons strongly interact with the microstructure
allowing for the study of planar defects. Another useful mode in TEM is selected area electron

61
diffraction (SAED), which generates diffraction patterns that provide crystallographic
information. SAED patterns can be used to determine lattice parameters, identify crystal structures,
texture, and phase fractions [125].
Scanning transmission electron microscopy (STEM)
In scanning transmission electron microscopy (STEM), a focused electron beam is rastered across
a thin specimen in a rectangular pattern. The electrons transmitted through the sample can be
gathered by two different detectors. The bright field detector collects the electrons transmitted
parallel to the beam, generating bright field STEM images. A high angular dark field detector can
collect high angle incoherent scattered electrons, which are used to from high angle annular dark
field (HAADF) STEM micrographs [126]. Additionally, interactions between the sample and
beam generate x-rays that contain compositional information regarding the specimen. Thus, energy
dispersive x-ray spectroscopy (EDS) can be utilized to obtain compositional maps with resolution
similar to the spot size. Figure 48 shows the compositional maps (Figure 48a-b), compositional
scan (Figure 48d) and the HAADF micrograph (Figure 48c) of a Hf-Ti nanostructure [72].

62

Figure 48: STEM EDS maps of a nanostructured Hf-Ti sample depicting a) Hf b) Ti, c) the
representative HAADF image and d) composition line scan across region A-B. [72]
Automated Crystallographic Orientation Mapping (ACOM)
Similar to t-EBSD, automated crystallographic orientation mapping (ACOM) can be utilized to
obtain phase and crystallographic information at the nanoscale. Specifically, this technique
involved the scanning and precessing of the electron beam along different points of the sample to
generate electron diffraction spot patterns that are collected by a charged coupled device (CCD).
The capabilities of the CCD enables the collection of up to 180 diffraction patterns per second
making it possible to scan relatively large regions [127]. The value of such a technique is
highlighted in Figure 49, where ACOM-TEM allows for the characterization of a Pd sputtered film
with an average grain size of 35 nm [128]. Thus, ACOM-TEM enables orientation mapping with

63
the high resolution of TEM. As such, this technique is crucial for the characterization of grain
boundaries and orientation of nanostructured materials.

Figure 49: ACOM-TEM maps of a Pd sample a) of the cross-section and b) plan view of a
sputtered Pd sample. C) and d) show the respective quality maps of each scan. [128]
3.2.5 Atom probe tomography (APT)
Atom probe tomography (APT) is the only characterization technique able to detect individual
atoms of elements in a three-dimensional space [129, 130]. Thus, making it a powerful tool to
examine the segregation behavior of nanostructured materials. To perform an APT measurement,
extensive sample preparation is needed, where the specimen must be formed into a needle shape
with a tip radius less than 100 nm. This can be done in a variety of ways, but is typically performed

64
through FIB milling or electropolishing [121]. Once inside the APT chamber, a DC voltage is
applied to the specimen below the evaporation condition for the material. A femtosecond laser is
then used to evaporate one or more ions from the surface of the specimen tip that reach a 2D
position sensitive detector. The detector can determine the X and Y coordinates of the ions and
also measure the time of flight from the onset of the laser pulse to the point when the ions hit the
detector. With this information, mass over charge ratios can be determined and three-dimensional
reconstructions of the specimen can be generated with atomic chemical sensitivity and high spatial
resolution. Figure 50 specifically displays the principles of laser-assisted APT.

Figure 50: Illustration of a pulsed laser atom probe, where a femtosecond laser causes the
evaporation of individual ions that get collected by a 2d X-Y detector.
From an APT measurement, a 3D reconstruction of an atomic maps consisting of all the detected
ions is generated. An atomic map of a nanocrystalline Cu-Nb sample upon annealing to 400˚C is
shown in Figure 51a. The individual ions are color coordinated, where in this specific example
only Cu and NbOx, ions are highlighted in red and purple respectively. From that given data,

65
proxigrams or 1D concentration profiles, shown in Figure 51b, can be produced to measure the
elemental composition across a region of interest. Isoconcentrations (isosurfaces) and 2D density
plots can be created to emphasize and highlight elemental variants (Figure 51c).

Figure 51: a) APT reconstruction of a ball-milled CuNb system, where Nb rich clusters are
highlighted in the b) proximity histogram c) applied isosurface reconstruction, and d) 2d density
plots. [68]
APT has been fundamental in understanding the elemental evolution of nanostructured alloys and
detect trace elements that can influence material properties. [68]. For this reason, APT has become
an indispensable analytical tool in studying the thermal stability, and more importantly the role of

66
thermal stability mechanisms of nanostructured materials. Another benefit of APT, is the ability
to detect nanocluster consisting of a few atoms that indicate early stage of phase separation and
precipitate formation. In the case of ball milled NC Cu-Nb sample upon annealing to 400˚C,
several Nb rich clusters were identified using APT, which while hypothesized in other studies had
not been previously observed [68]. The atomic map is shown in Figure 51a, where the detected
ions are coordinated to a specific color, to the right of the atomic maps is a proximity histogram,
which gives the 3D concentration from the start of a selected interface, which in this case is the
Nb rich clusters. These Nb rich clusters, along with clusters detected in other alloys, are expected
to provide a source of thermal stability [110]. Furthermore, these clusters are highlighted in Figure
51c, through 2D density plots or isoconcentration maps.

Copyright 2020 Joel Antunez Bahena
4 GBE – Effect of initial twinned microstructure on the thermal
evolution of highly NT alloys
The following section is adapted from the work titled Grain Boundary Evolution of Highly
Nanotwinned Alloys: Effect of Initial Twinned Microstructure that has been submitted for peer
review to the journal Scripta Materialia.
In this study, a comprehensive examination of the initial NT microstructure is conducted, in terms
of grain growth and thermal stability. Previous studies have shown that NT metals are more
thermally stable than their nanocrystalline (NC) or ultra-fined grained (UFG) counterparts. This is
a result of the as a function of twin boundaries pinning columnar grain boundaries and impeding
grain boundary migration [47, 48]. However, other studies have suggested that a high density of
twins is insufficient to stabilize the microstructure at elevated temperatures, requiring an
investigation of the initial twinned microstructure [49, 50]. To date, a comparison across similar
NT microstructures has not been conducted especially since different synthesis methods can
dictate characteristics (grain size, texture, twin thickness, residual stresses, etc) that impact thermal
evolution. Although the role of these microstructural features is not well understood, it has been
proposed that decreasing the twin thickness could improve thermal stability by increasing the
number of triple junctions and constraining grain boundary motion [131]. Nevertheless, despite
twin boundaries typically exhibiting an order of magnitude less energy than high angle grain
boundaries, a large twin boundary density could increase the overall grain boundary energy and
initiate the onset of grain growth at lower temperatures [47]. Reducing the grain size and/or grain
width could have a similar affect and encourage a tendency towards grain coarsening, especially
at the nanoscale [7]. Furthermore, experimental results point towards certain crystallographic

68
textures and intrinsic residual stresses promoting grain coarsening and abnormal grain growth
behavior in nanostructured materials [93, 132]. Thus, there is potential that particular grain
orientations and stored energy distributions could be more favorable in resisting grain boundary
mobility.
In this investigation, the influence of twin thickness, grain width, texture, and intragranular
orientation spread are evaluated to investigate the factors affecting grain growth in sputtered NT
Cu alloys. Specifically, transmission electron microscopy (TEM) was used to thoroughly
characterize as-deposited films and the thermal microstructural evolution was analyzed through
electron back scattering diffraction (EBSD). Overall, this study examines how the characteristics
of an initial twinned microstructure can lead to the activation of distinct thermal processes and
contribute to mitigating grain growth in NT metals.
4.1 Synthesis of NT Cu Alloys
NT Cu alloys were synthesized by magnetron sputtering Cu-2wt.%Al and Cu-10wt.%Ni targets
onto Si (100) substrates, where the different samples of each composition will be referred to as
Cu2Al and Cu10Ni. Two samples of each composition were produced with tailored
microstructures and twin thicknesses (λ) by altering the sputtering parameters similar to reports by
Velasco et. al [42]. An XP-2 stylus profilometer (AMBiOS) was used to measure the thicknesses
of each film which ranged from 14-19 μm. Electron transparent lamellae were prepared with a
Fischione 1050 ion mill and were characterized via TEM using a JEM-2100F (JEOL) microscope.
In addition, X-ray diffractometry measurements were performed using a Rigaku Ultima IV
diffractometer to obtain crystallographic information. In order to investigate the microstructural
evolution of the NT alloys at elevated temperatures, free-standing as-sputtered films were

69
successively heat treated at temperatures of 200˚C, 400˚C, and 600˚C for two hours in a vacuum
(< 10
-5
torr) tube furnace at a heating rate of 10˚C/min. Top-surface EBSD scans of the as-sputtered
and heat treated films were performed using a JEOL-7001 SEM with an EDAX Hitari Detector
and a FEI Helios G4 with an Oxford Symmetry Detector. EBSD datasets were analyzed with the
MTEX software [133]. To complement the study, micro-hardness testing was performed using a
Vickers tip with 10 g indents, totaling five indents for each sample.
Representative cross-sectional TEM images of the as-sputtered Cu alloys, grain width and twin
thickness distributions, and XRD scans are displayed in Figure 52. The TEM images, shown in the
top portion of Figure 52A-D, reveal microstructures of mostly vertical columnar grains containing
a high-density of twin boundaries perpendicular to the growth direction. Different magnifications
were utilized to capture key microstructural features. Below, the respective twin boundary and
grain width (inset) distributions are displayed for the four types of films. To obtain these plots, a
minimum of 200 grains and 500 twins were measured for each sample. Cu2Al-A, Cu2Al-B,
Cu10Ni-A, and Cu10Ni-B have an average grain width of 140 ± 40 nm, 130 ± 70 nm, 50 ± 20, and
260 ± 130 and an average twin thickness of 5 nm, 18 nm, 13 nm, and 31 nm, respectively. The
XRD patterns in Figure 52E, reveal a strong (111) texture for Cu2Al-A, Cu2Al-B, and Cu10Ni-B,
which is typical for sputtered FCC metals, while Cu10Ni-A exhibits a random texture and inclined
twins.

70

Figure 52: Cross-sectional bright field TEM comparing as-sputtered a) Cu-Al (λ=5), b) Cu-Al
(λ=18), c) Cu-Ni (λ=13), and d) Cu-Ni (λ=31); and E) representative XRD scans of the NT
alloys. The arrow indicates film growth direction.

71
4.2 Microstructural evolution of tailored microstructures
The grain size evolution of each sample was evaluated by EBSD in the as-sputtered condition and
after heat-treatments at 200˚C, 400˚C, and 600˚C, which is presented in Table 1. Additionally,
Vickers micro-hardness was used as a complementary technique to analyze microstructural
stability and provide insight to microstructural changes not detectable through top-surface
characterization [50]. It should be noted that the EBSD measurements of the planar grain sizes of
the as-sputtered films are consistent with the grain widths determined by cross-sectional TEM.
The grain size data from Table 1 is complementary to Figure 53, which displays the inverse pole
figure (IPF) maps of the top-surface in the as-deposited and annealed states with scans performed
near or at the same region after successive heat treatments. In addition, instances of significant
grain growth during the thermal microstructural evolution of the four sputtered films are indicated
by dashed boxes outlining scans in Figure 53. Note that the ~50 nm grain size of Cu10Ni-A in the
as-deposited film was not detectable by EBSD, therefore a planar-view TEM micrograph is
substituted to illustrate the microstructure. Overall at 200˚C few changes are observed for each
microstructure except for Cu10Ni-A, where the grain size increased from ~50 nm to ~120 nm
while maintaining a random texture. At 400˚C, Cu2Al-A shows sudden grain growth with larger
(100) grains consuming the majority of the microstructure, producing an average grain size of
~1550 nm and having large clusters comprised of randomly oriented grains dispersed throughout
film. The remaining three samples do not display significant grain coarsening until 600˚C where
Cu2Al-B and Cu10Ni-B undergo similar abnormal grain growth as seen in Cu2Al-A at 400˚C.
with an average grain size for both films increasing to over 5000 nm. In contrast, Cu10Ni-A
undergoes a second instance of coarsening at 600˚C but retains a relatively fine grain size of ~ 350
nm and a random texture.

72

Figure 53: Top surface EBSD IPF maps showing grain evolution of as-sputtered A) Cu2Al-A, B)
Cu2Al-B, C) Cu10Ni-A, and D) Cu10Ni-B after heat treatments performed at 200°C, 400°C, and
600°C. Dashed boxes highlight instances of significant grain growth, note the changes in scale
bar. The IPF triangle is shown to the right of the scans.
The evolution of hardness with increasing temperature closely correlates with the observed grain
size changes. In the as-deposited state, the hardness values were all greater than 300 HV, with
Cu10Ni-A exhibiting the highest hardness at 400 HV. The recorded values are comparable to those
of previous NT Cu studies with similar grain widths [48]. At 200˚C, Cu10Ni-A shows a minimal

73
decrease in hardness, which would suggest that the nanotwinned structure is retained despite this
initial increase in grain size. Furthermore, at 400˚C the hardness for Cu2Al-A drops below 100
HV corresponding with the observed abnormal grain growth. In addition, moderate drops in
hardness for both Cu2Al-B and Cu10Ni-B are recorded despite no significant change in grain size
and could indicate the onset of detwinning and relaxation of residual stresses [49, 134, 135]. A
more significant decrease is observed at 600˚C at which abnormal grain growth occurs for both
films. The hardness for Cu10Ni-A, even at 600˚C, remains relatively high at 170 HV, reflecting a
more thermally stable sample.
Table 1: Summary of hardness (HV) and grain size from EBSD in sputtered Cu Alloys with
successive heat treatments at 200°C, 400°C, and 600°C.

Cu2Al-A
(λ =5 nm)
Cu2Al-B
(λ=18 nm)
Cu10Ni-A
(λ=13 nm)
Cu10Ni-B
(λ=31 nm)
Hardness
(HV)
Grain
Size
(nm)
Hardness
(HV)
Grain
Size
(nm)
Hardness
(HV)
Grain
Size
(nm)
Hardness
(HV)
Grain
Size
(nm)
AS 340 ± 2 130 310 ± 6 140 400 ± 3 50 300 ± 6 250
200 °C 340 ± 3 140 300 ± 4 150 370 ± 4 120 280 ± 2 270
400 °C 90 ± 8 1550 210 ± 5 165 350 ± 2 150 250 ± 4 290
600 °C 90 ± 3 1580 100 ± 3 5100 170 ± 2 340 80 ± 1 7800
4.3 Contributions from stored energy
In addition to grain size and twin thickness, microstructural heterogeneity has also been shown to
influence grain growth behavior of NT materials [47, 49]. For instance, it has been previously
reported that grain coarsening and recrystallization processes in NT Cu can be selective depending
on grain characteristics [50]. Likewise, a heterogenous distribution of stored energy has been
shown to activate distinct grain growth or recrystallization mechanisms in many coarse-grained
alloys, which could provide more insight into the varied microstructural evolutions observed in

74
this study [115, 136, 137]. In order to obtain a qualitative representation of the evolution of stored
energy and insight into thermal processes occurring in these films, grain orientation spread (GOS)
maps of the samples annealed at 200˚C and 600˚C are presented in Figure 54. The GOS spread is
colored from 0˚ to 3˚, where studies have demonstrated that grains with larger orientation spreads
(red) are expected to have higher stored energy [115]. Considering that not all samples were
resolved by EBSD in the as-sputtered condition, 200˚C was selected as the initial temperature for
comparison purposes. At 200˚C, the GOS maps indicate a relatively wide GOS distribution for
both CuAl samples (Figure 54A and Figure 54B) as illustrated by the heterogenous collection of
grains with low (low stored energy) and large (high stored energy) orientation spreads. Meanwhile,
the GOS map for Cu10Ni-A (Figure 54C) displays a mostly homogeneous distribution with
relatively low spreads, whereas Cu10Ni-B (Figure 54D) exhibits moderate heterogeneity. In terms
of abnormal grain growth, grains with lower stored energy have a competitive advantage over
those with higher stored energy [115], which could contribute to the formation of abnormally
large grains in Cu2Al-A, Cu2Al-B, and Cu10Ni-B. As previously discussed, these three films
exhibit strong (111) textures which points to a correlation between stored energy and the initial
texture, as seen in other studies [138-140]. Upon annealing at 600˚C, the overall GOS values
decreased across all films, but several abnormally large grains retain relatively large orientation
spreads, implying the microstructures have yet to reach a steady state [84]. In the case of Cu2Al-
A, the clusters of randomly oriented grains are characterized (encircled) by large networks of
annealing twins and low orientation spreads, which indicate recrystallization and the presence of
twin related domains (TRDs) [119, 120]. Cu10Ni-A also exhibits low orientation spreads with
large network of annealing twins throughout the entire microstructure, suggesting recrystallization
[120].

75

Figure 54: Top surface EBSD maps showing the GOS distribution for samples annealed at
200˚C and 600˚C. Σ3 grain boundaries are outlined in red.

76
4.4 Mechanisms controlling grain growth behavior
Based on observations from the microstructural evolution and hardness measurements, the initial
NT microstructure appears to have a pronounced effect on the grain growth behavior. Through
comparing the different alloy pairs, the influence of distinct microstructural features can be
identified. Cu2Al-A and Cu2Al-B share similar microstructural characteristics, but the twin
density is ~3 times greater for Cu2Al-A. It was shown that both samples exhibit some level of
abnormal grain growth dominated by the formation of large (100) grains, but the onset occurs at
400˚C for Cu2Al-A and at 600˚C for Cu2Al-B. It is possible that a decrease in twin thickness could
initiate thermal process at lower temperatures by the increase in excess energy associated with
twin boundaries [47, 49, 50]. For example, in a previous study examining NT Ag (λ=10 nm), the
calculated driving force of 1000 kJ/m
3
, was predicted to significantly contribute to the activation
of abnormal grain growth [51]. In the case of Cu2Al-A and Cu2Al-B, similar calculation were
performed assuming a twin boundary energy of 24 mJ/m
2
[141], which yields estimated driving
force values of 4800 kJ/m
3
and 1333 kJ/m
3
, respectively. The increased driving force produced by
the higher density of twin boundaries is further highlighted by the presence of TRDs in select
regions of the Cu2Al-A microstructure, which typically require a higher driving force to initiate
when compared to abnormal grain growth [119]. Although the onset of grain growth in Cu2Al-A
starts at a lower temperature, this does not necessarily result in a larger grain size, since there are
always multiple mechanism present. For example, the formation of TRDs in Cu2Al-A impinge
and restrict the growth of abnormally large grains that would otherwise consume higher stored
energy regions, leading to a smaller grain size by a factor of ~3. However, these observations do
indicate that a critical twin thickness exists that can reduce thermal stability by driving the early
onset of abnormal grain growth and recrystallization as shown for Cu2Al-A at 400˚C.

77
Further insight into the implications of initial grain size and texture is gained by extending
observations to the grain growth behavior of Cu10Ni-A and Cu10Ni-B. In terms of these
microstructural features, Cu10Ni-A has a grain size of ~ 50 nm with a random texture, while
Cu10Ni-B has a grain size of ~250 nm and is strongly oriented in the (111) direction. The effects
of differing initial grain size is evident after the 200˚C heat treatment, where Cu10Ni-B remains
microstructurally similar, while grain growth is observed for Cu10Ni-A as a function of the ~5-
fold increase in the grain boundary density [7]. Despite the early onset of coarsening and the larger
driving force associated with a higher grain boundary energy, at 600˚C, Cu10Ni-A retains a
relatively small grain size. This discrepancy in grain growth behavior appears to be influenced by
the texture evolution of each film, where EBSD revealed that Cu10Ni-A preserves a random
texture and Cu10Ni-B exhibits a (111)-to-(100) texture shift. Although a strong (111) texture can
provide improved thermal stability owing to a high fraction of low angle grain boundaries [47, 48,
142], the presence of a few randomly oriented grains can introduce boundaries with higher
mobilities that facilitate abnormal grain growth [84, 143]. In contrast, Cu10Ni-A exhibits reduced
anisotropy due to inclined twin boundaries promoting changes in orientation of the columnar
grains [43]. As a result, the random texture diminishes the competitive advantage of any specific
grain and prompts a more gradual grain growth behavior. From our observations across all
samples, initial film texture appears to be the most dominant microstructural characteristic
influencing thermal stability followed by twin density
4.5 Summary
The thermal evolution of four sputtered NT Cu alloys have been compared to understand the effect
of initial twinned microstructure on grain growth behavior at elevated temperatures. Texture

78
induced by inclined TBs appears to be the most crucial factor in facilitating grain growth behavior,
where NT films with a strong (111) texture demonstrated the tendency to form abnormally large
(100) grains and the film with a random texture (Cu10Ni-A) demonstrated a stepwise growth
behavior. Additionally, previous studies theorized that a large twin density can stabilize NT metals,
but the present work indicates that there is a threshold to this stabilizing effect, where a smaller
twin thickness can drive the early onset of abnormal grain growth and recrystallization. Overall,
the findings from this study provide insight into the role of the specific microstructural
characteristic that can be leveraged to design more stable NT systems.

79
5 GBSE – An investigation of the of the Mo-Au system
A version of the following work is published as a journal article titled Thermally Activated
Microstructural Evolution of a Sputtered Nanostructure Mo-Au in Materialia, Volume 4, pages
157-165 (DOI: 10.1016/j.mtla.2018.09.019).
In the present study, the thermal processes that govern the microstructural evolution of sputtered
Mo-Au NMMs systems are investigated. Specifically, Mo-Au nanostructures are examined at
annealing temperatures ranging from 350°C to 1000°C, while identifying key thermally activated
mechanisms. Differential scanning calorimetry (DSC) is used as a guide to determine critical
thermal events where crucial microstructural changes could occur. These type of scans have been
used to detect phase transformation [144, 145] and thermal microstructural evolution [82] in
NMMs systems. The selection of the Mo-Au alloy was based on nanocrystalline thermodynamic
stability maps, where Mo-Au is predicted to achieve a stable nanograin configuration through
solute-stabilized grain boundaries [10, 60]. Complimentary characterization techniques of
transmission electron microscopy (TEM) and atom probe tomography (APT) provide a
comprehensive analysis of the microstructural transitions at different length scales over the wide
range of temperatures [129]. In combining these techniques with a predicted thermally stable
system, the underlying processes and mechanisms that contribute to nanograin stability can be
better understood.
The decision to select the Mo-Au system was guided by the thermodynamic stability maps
constructed by Murdoch and Schuh. In their work, binary alloys were classified based on their
proclivity for nanograin stability in stable, metastable, and unstable regions. These Mo-Au alloys
was predicted to show nanograin stability, and therefore was selected to be a model system for the

80
study of the thermal evolution of a nanostructured system. Furthermore, the relatively low melting
temperature of Au (1000˚C) allows for a thorough investigation of the solute behavior at lower
temperatures, which can be achieved under experimental conditions. Upon selection of the binary
selection, NMM samples with alternating Au and Mo-Au layers were synthesized using DC
magnetron sputtering from high purity Au (99.99%) and Mo (99.95%) targets. A dual sputtering
source system allowed for the deposition of Mo and Au onto (100) silicon substrates, with
sputtering powers of 300 W and 7 W for each source, respectively. During deposition, an argon
working pressure of 0.8 Pa was utilized. The aim was to produce NMMs, with compositions of
approximately 20 at.% solute content. As a design choice, the individual layer thickness was
restricted to be less than 50 nm to increase the density of interfaces in the materials. The total
overall film thickness of the samples was 2 μm.
The microstructural evolution of the Mo-Au NMMs were examined over an extended temperature
profile through DSC scans performed with a Labsys thermal analyzer. The enthalpy change (ΔH)
of the endothermic and exothermic peaks was measured over a baseline that was removed from
the scans [146]. The scans were conducted on freestanding films from 20°C to 1000°C at a rate of
10°C/min under a constant argon flow of 40 ml/min. Individual pieces of the free-standing films
were heat treated in a GSL1100X tube furnace (MTI Corporation) at temperatures of 350°C,
550°C, 800°C, and 1000°C for 96 h under a vacuum pressure of 5 x 10
-4
Pa. Upon completion of
the heat treatment, the samples were quenched in a low vapor pressure oil, (Invoil 705, Inland
Vacuum Industries). A vertical tube furnace geometry was utilized that maintained quench oil
temperatures near 25˚C and allowed for quenching under vacuum conditions. As a final step,
samples were cleaned with ethanol after the heat-treatment process.

81
The microstructure and elemental compositions of the heat-treated samples were characterized by
TEM and STEM using a JEOL JEM-2100F microscope combined with an EDS detector. TEM
and APT lamellae of the films were prepared by Focused Ion Beam (FIB) liftouts using a JEOL
FIB-4500 microscope. A 150 nm platinum protective layer was deposited onto the region of
interest to minimize FIB Ga beam damage. Furthermore, APT tip preparation was performed with
a ZEISS Auriga 60 FIB, resulting in needle shaped specimens with end diameters less than 100
nm. APT measurements were carried out on a CAMECA-LEAP 4000X HR instrument utilizing a
laser pulsing mode with a wavelength of 355 nm. The APT parameters for samples annealed at
550˚C and 800˚C were set to a pulse energy of 100 pJ, pulse repetition rate of 100 kHz, and pulse
evaporation rates of 0.3% at a base temperature of 60 K. For the measurements of the 350˚C
annealed samples, the pulse energy and base temperature were decreased to 10 pJ and 25 K,
respectively. Reconstruction and atom probe analysis was evaluated using the CAMECA
Integrated Visualization and Analysis Software (IVAS 3.6.1). A JSM-7001 scanning electron
microscope (SEM) was used to determine global composition by EDS.
5.1 As-sputtered microstructure
The as-sputtered Mo-Au microstructure is shown in the cross-sectional bright field TEM images
in Figure 55a. The bright regions depict the 17 nm thick Mo-Au (11 at.% Au) co-sputtered layers,
while the darker regions represent the 2.5 nm thick Au layers (99.95 at.% Au). The composition
of the overall structure, as measured through SEM/EDS, was approximately 21.3 at.% Au. The
accompanying dark field images in Figure 55b. illustrate the columnar structure of the as-sputtered
NMMs. Over 200 grains were examined and the average grain size was determined to be 9 nm,
where the grain size distribution is presented in Figure 55c. Grain size was calculated using similar

82
procedures detailed in a recent study by Riano et al [82]. The SAED pattern in the inset of Fig. 37
shows a streaked pattern with strong Mo (110) diffraction spots indicating that the sample exhibits
a (110) texture. The DSC scan of the Mo-Au system is shown in Figure 55d from 20-1000˚C. From
these thermal measurement, two evident exothermic peaks were observed between 400˚C and
1000˚C, which highlight microstructural transformations in the Mo-Au sample. A peak indicating
recrystallization appears between temperatures of 400˚C and 500˚C with ∆𝐻
4.34!5(6,,'>6('+&
= -
1.7 kJ/mol. A second exothermic peak is observed starting at 620˚C and continues until the
maximum tested temperature of 1000˚C suggesting grain growth with ∆𝐻
B46'&9B4+C(D
= -21
kJ/mol.

Figure 55: Mo-Au NNMs (a) bright field TEM with inset SAED patterns (top), dark field TEM
(middle), grain size distribution (bottom), and (b) differential scanning calorimetry scan of the

83
Mo-Au multilayer from 20°C to 1000ºC. From the scan, thermal events were indicated and
enthalpy changes (ΔH) were calculated.
In order to determine the thermal processes associated with the identified microstructural
transformations, analysis from the DSC scans was combined with direct observation of the
microstructure at select annealing temperatures. Specifically, samples were heat treated to
temperatures of 350˚C, 550˚C, 800˚C, and 1000˚C to capture the microstructural transformations
that encompass the recrystallization and grain growth processes. Characterization of
microstructural changes such as grain size, morphology, and elemental distribution throughout the
microstructural evolution of the Mo-Au sample were performed through TEM and EDS. Cross-
sectional TEM micrographs, EDS maps, and grain size distribution charts of the samples at
different annealing temperatures are presented in Figure 56a-d. Texture and phase changes of the
microstructure were also examined as a function of the annealing temperature through normalized
integrated radial intensity profiles extrapolated from SAED patterns from the bright field TEM
images. The respective integrated radial intensity plots are highlighted in Figure 57a-e.

84

Figure 56: Cross-sectional bright field TEM with insets of respective SAED patterns (top),
corresponding EDS maps (middle), and grain size distribution (bottom) of Mo-Au samples heat-
treated at (a) 350ºC, (b) 550ºC, (c) 800ºC, and (d) 1000ºC.
Furthermore, to complement TEM characterization, APT was performed on specific heat-treated
samples to further examine the microstructure and to gain chemical information at a near atomic
spatial resolution. In particular, APT enables the effective study of the chemical analysis of
multilayer interfaces, solute segregation behaviors, and provides prevalent insights regarding
microstructural evolution [147, 148]. The atom probe reconstruction and corresponding
concentration profiles of selected regions are presented in Figure 58, Figure 59, and Figure 60.
which display Mo atoms (shown in blue) and Au atoms (shown in yellow) for the 350˚C, 550˚C,

85
and 800˚C annealed samples. A comprehensive overview of the thermal evolution of the Mo-Au
nanostructure, composition, microstructural changes and active thermal mechanisms will be
discussed in the next sections at each selected temperature in terms of the processes

Figure 57: Integrated radial Intensity profiles interpolated from SAED patterns obtained from
Mo-Au samples (a) as-sputtered condition and heat treated at (b) 350ºC, (c) 550ºC, (d) 800ºC
and (e) 1000ºC.

86
5.2 Multilayer degradation region, T= 350˚C
The first selected temperature for this study was 350˚C; TEM images and EDS maps (Figure 57a)
show a multilayer structure with no noticeable changes in the average grain size (9 nm) or grain
size distribution when compared to the as-sputtered sample. At this annealing temperature, the
DSC scan did not indicate any apparent thermal events. Investigation of the multilayer interface
through APT compositional measurements revealed roughening and intermixing of the layers. The
atom map in Figure 58a reveals a multilayer structure where a 17 at.% Au isosurface is applied to
highlight the interfaces of the Mo-rich and Au-rich layers. Figure 58b presents a 1D compositional
measurement across several layers which follows the dashed arrow in Figure 58a. The Mo-rich
layers had an average concentration of 9.1 ± 1.5 at.% Au and 90.5 ± 1.6 at.% Mo which matches
closely to the values obtained through EDS (11 at.% Au). While it is anticipated that Mo and Au
should demonstrate a strong tendency to separate at this composition, the Mo-rich layers remained
in a solid solution at 350˚C. Additionally, the concentration profile reveals a wide inter-diffusion
region from the center of the Au-rich layers, indicating that the as-sputtered 2.5 nm Au layers have
diffused and intermixed with the Mo-rich layers as a result of the increase in temperature. The
center of the Au-rich layers have an average concentration of 19.3 ± 1.9 at.% Au, but the Au
content gradually decreases as the measurements mover further from the center. Intermixing
between layers has been reported in other annealed immiscible NMMs and has been attributed to
strain and interfacial contributions introduced during sputtering [149, 150]. Furthermore,
intermixing at the interface of the layers at increased temperatures is considered a precursor for
layer breakthrough which initiates degradation of the multilayer configuration [151, 152]. These
findings are further supported by the SAED patterns and integrated radial intensity profiles in
Figure 56a and Figure 57b, where a decrease in the intensity of the peaks and blurring of the

87
diffraction rings indicates a reduction in the overall crystallinity of the layered grains [153]. The
streaked diffraction pattern indicates that the (110) texture of the sample was still preserved. Thus,
350˚C marks the onset of the initial stages of multilayer structure degradation in the Mo-Au
system.

Figure 58: a) APT map from the Mo-Au sample annealed at 350˚C showing the multilayered
structure with a 17 at.% Au isosurface. The 1D concentration profile (b) obtained from the
region indicated by the dashed arrow shows the composition across several layers.

88
5.3 Recrystallization region, T=550˚C
As the temperature is further increased to 550˚C, the microstructure of the Mo-Au NMMs
undergoes significant structural and morphological changes. The original layered and columnar
structure has evolved into a nanometallic composite structure consisting of Mo-rich grains
surrounded by Au solute, which is observed in the TEM micrograph and EDS map in Figure 56c.
During this microstructural transition, the average grain size has increased to 38 nm. Similar
microstructural changes have been observed in several other immiscible NMMs systems at
elevated temperatures [72, 78, 79, 81], where breakdown of the layers is generally attributed to
thermal grooving, subsequent breakthrough along the columnar GBs, and eventual pinch-off of
the layers [151]. Those processes are dominated by the mass transport of constituent atoms along
interfaces with no significant changes in texture, suggesting that recrystallization is not associated
with this morphological transformation [81, 154]. However, observations of the SAED pattern
(Figure 56b) and the accompanying radial distribution profiles (Figure 57c) reveal the loss of the
streaked spot patterns and several new Mo reflections. This weakening of texture and reorientation
of the grains seen in the diffraction pattern suggests a recrystallized microstructure [90, 155].
Furthermore, recrystallization at this temperature, is supported by the exothermic peak seen in the
DSC scan (Figure 55d). Overall these findings indicate that the microstructural changes in Mo-Au
at 550˚C are not fully expressed by typical multilayer degradation stages and are significantly
influenced by the recrystallization process. Therefore, it is expected that recrystallization is
initiated by the interfacial energy stored in the NMMs structure due to the high density of phase
boundaries and strain energy generated by differences in radius between Mo and Au [156]. Thus,
the stored energy drives the formations of new crystallites that consume the strained material,

89
leading to the eventual transformation from a NMMs consisting of columnar grains to the observed
nanometallic composite structure [82, 157].
Similar to observations in recent NMMs studies [72, 82], the segregation of the solute element
appears to initiate during recrystallization, where the solute interaction with the GBs tends to
dominate the thermal evolution of nanostructured materials [7, 158]. Segregation of solute atoms
can either enrich or form clusters and precipitates at GBs, with these behaviors contributing to
limiting grain growth. To further understand solute behavior in the Mo-Au system, APT was
utilized to elucidate the formation and evolution of these features along GBs in a three-dimensional
space. Two APT maps were analyzed and presented in Figure 59a.1-a.2. The corresponding 1D
composition profiles of the regions indicated by the dashed arrows are shown in Figure 59b.1-b.2,
where a distinct grain boundary is identified in each scan. Interpretation of GBs was supported by
2D chemical density plots (not shown) and the assumption that segregation species occur along
2D and 3D features [68]. A 5 at.% isosurface was applied to the APT maps to further highlight
and identify segregation zones [58, 68, 159] and clustering of Au solute atoms.
Analysis of the first APT map in Fig. 5a.1, highlights a Au segregation zone that has a peak
concentration of 60.8 ± 1.4 at.% Au and 38.6 ± 1.5 at.% Mo, with < 1 at.% O. The surrounding
Mo grains reach concentrations of 97 ± 1.1 at.% Mo with minimal Au content (< 1 at.%), and 1 ±
0.8 at.% O. The composition within the grains indicates that the Mo-Au solid solution observed in
the multilayer configuration at 350˚C has decomposed after annealing to 550˚C. The detected Au
segregation zone in APT correlates with the observed microstructure in the TEM and EDS maps
(Figure 56b) that show Au solute surrounding Mo-rich grains. From these results, it can be inferred
that as recrystallization occurs, the diffusivity near the grain boundary increases, which drives

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segregation of Au to the GBs resulting in GB solute enrichment and widening [17, 57, 59, 157,
160]. These observations are in agreement with thermal stability maps, where solute segregation
to GBs was predicted to reduce GB energy and inhibit grain growth of the Mo-Au system [10].

Figure 59: APT maps from the Mo-Au sample annealed at 550˚C showing solute behavior at
GBs. The first APT map (a.1) shows a Au enriched GB (segregation zone) with insets showing
individual Au and Mo ion distribution. The 1D concentration profile (b.1) is obtained through
the region indicated by the dashed arrow, where the inset shows concentration of (C + O). The
second APT map (a.2) shows clusters along a GB with insets showing individual Au and Mo ion
distribution. The 1D concentration profile of (b.1) is obtained through the region indicated by

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the dashed arrow, where the inset shows the Au content through “Cluster 2” obtained from the
region indicated by the blue cylinder.
To further analyze the solute behavior in the Mo-Au systems at 550˚C, a second APT map of a tip
extracted from a similar region is presented (Figure 59a.2). In contrast to the GB shown in Figure
59a.1, where solute segregation of Au along GBs was exhibited, the second map (Figure 59a.2)
revealed the presence of Au clusters along a GB. The composition through two of the clusters and
the surrounding Mo-rich grains was analyzed though a 1D composition profile. The concentrations
of the Au-rich clusters were measured to be ~8-9 at.% Au and ~91 at.% Mo with no detectable
presence of contaminants. Additionally, the surrounding Mo-rich grains exhibited concentrations
of 98± 0.5 at.% Mo with traces of C and O. The formation of these solute clusters along the GBs
are indicative of early onset of phase separations, which can facilitate the growth of precipitates
[68, 161, 162]. Presence of solute clusters along GBs and within the matrix have also been shown
to reduce GB mobility through Zener pinning [61, 68, 110]. Therefore, it is possible that the Au-
rich clusters located at GBs have influenced the microstructural evolution of the Mo-Au system
through kinetic thermal stability mechanisms.
These results suggest that the Au segregation behavior along the GBs is heterogeneous and could
be dependent on several factors. Recent studies have demonstrated that immiscible alloys can
exhibit different segregation tendencies depending on GB character [162, 163]. Furthermore,
modeling frameworks of GB solute behavior of immiscible systems have described competing
mechanisms between GB solute segregation and phase separation which can lead to the observed
heterogenous behavior seen in the Mo-Au system [57]. Experimentally, the simultaneous
occurrence of thermodynamic and kinetic stabilization mechanisms have been observed in only a

92
few binary alloy nanostructures. For example, in Cu-Nb, Nb-enriched segregation zones and Nb-
rich clusters were identified along GBs, with both kinetic and thermodynamic stability
mechanisms contributing to thermal stability of the system [68]. However, the reaction of Nb
clusters with oxygen impurities (~15 at.% NbO and NbO2) during synthesis complicated the ability
to accurately assess solute behavior in Cu-Nb. In contrast, for the case of the Mo-Au system in this
study, the APT compositional profile of the Au-enriched clusters at 550˚C show no significant
concentrations of impurities indicating that these do not appear to influence the nucleation and
stabilization of these clusters.
5.4 Grain growth region, T= 800˚C
At temperatures ranging from 620˚C to 800˚C, the DSC scans indicate that grain growth should
be the primary thermal process, which is in agreement with the microstructure seen at 800˚C
(Figure 56c). The TEM micrographs and EDS map show a similar elongated morphology to that
seen at 550 ˚C (Figure 56b), with an increase in the overall average grain size to 67 nm. In addition,
the Au precipitates primarily appeared in the lower and upper portions of the film, which indicates
preferential diffusion of Au solute to these regions. This behavior has also been observed in other
sputtered films, for example, Ti precipitates in a Hf-Ti NMMs formed in two regions equidistance
from the center of the sample when heat-treated at 1000˚C [82].
The resulting microstructure can be viewed as a transitional nanostructure that bridges the
nanocomposite structure (Figure 56b) to the equiaxed grain structure observed at 1000˚C (Figure
56d). Specifically, when examining the microstructural evolution from 550˚C to 800˚C it is
expected that the Au clusters that formed along the GBs coarsen due to increased kinetics and the
high concentration of Au solute atoms present at the GBs. Furthermore, since the Mo-Au system

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is immiscible and therefore thermodynamically unfavorable in a mixed configuration, sufficient
heating encourages further coarsening of Au-rich precipitates until a phase separation between Mo
and Au is eventually reached. The phase separation of the Mo-Au alloys was investigated through
APT, where chemical analysis was performed across a Mo and Au region. The atomic distribution
of the dominant elements is presented in Figure 60a and Figure 60b, while Figure 60c presents the
1D compositional profile through the region marked by the dashed arrow. The atom map reveals
four distinct microstructural regions: two Au-rich precipitates and two Mo-rich grains. The Au and
Mo regions which appear at the top and bottom respectively, are separated from the other two
grains by approximately half-spherical interfaces stemming from reconstruction problems that
may be related to micro-fractures, which means that these interfaces are not considered during
evaluation. The Au precipitates were composed of 98 ± 0.3 at.% Au, with ~ 1 at.% Mo. One of
the Mo-rich grain had a compositions of 75 ± 2.2 at.% Mo and 24 ± 1.9 at.% C, while the other
had a composition of 72 ± 3.6 at.% Mo and 26 ± 3.2 at.% Si. Similarly, APT revealed minimal Au
atoms (< 0.5 at.%) in these Mo-rich grains. While the presence of contaminants, which are
discussed at the end of this section, complicates a thorough analysis, the sharp transitions and
compositions between the Mo-rich grain and Au precipitate seen in the compositional profile
(Figure 60c) suggests phase separation of the alloy has been achieved at 800˚C. Reflections
observed in the radial distribution plots (Figure 57d) also display an increased intensity of Au
peaks relative to the Mo peaks, further supporting that at 800˚C the microstructure is composed of
phase separated bcc Mo grains and fcc Au precipitates.

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Figure 60: APT maps from the Mo-Au sample annealed at 800˚C, where (a) shows the ion
distribution of Mo (30 at.% isosurface) and Au (90 at.% isosurface), and (b) shows the ion
distribution of C (15 at.% isosurface) and Si (10 at.% isosurface). Regions are separated where
possible micro-fractures have occurred during APT measurements. These interfaces are not
evaluated. The 1D concentration profile (c) obtained from the region indicated by the dashed
arrow spans across a Au precipitate and a Mo-rich grain.
In other binary nanostructured alloys, such as Fe-Cu and Cu-Co , the general observation has been
that phase separation and the formation of secondary phase precipitates leads to an accelerated
grain coarsening [64, 160]; however, this study showed that the Mo-Au systems has maintained a
sub 100 nm grain size upon annealing at 800˚C for 96 hrs. The retention of nanograins in the Mo-
Au alloy appears to be attributed to influences from both kinetic and thermodynamic stability
mechanisms. Initially, as observed at 550˚C, thermodynamic and kinetic contributions are present
in the form of Au segregation zones and Au clusters, but as the temperature increased only one

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mechanism becomes prevalent. This switch in mechanisms is presumably facilitated by the
diffusion behavior of Au, where the diffusion along GBs promotes the growth of clusters leading
to the observed Au precipitates in the EDS maps (Figure 56c). As a result of coarsening, the Au
segregation zones are depleted of GB solute and the thermodynamic contribution is ultimately
reduced [69]. From the absence of segregation zones and coarsened Au precipitates, it can be
deduced that the microstructure at 800˚C was primarily stabilized by kinetic mechanisms through
GB pinning.
The extent to which contaminants affect the observed microstructures is always a point of
discussion since both the quantity and time of contamination can greatly alter microstructural
evolution. As stated earlier, no significant contaminants were detected in samples heat-treated
below 550˚C nor in similar as-sputtered samples [72, 82]. However, for samples annealed at
800˚C, APT measurements revealed more than 20 at.% C and Si in the Mo regions; where the most
probable source of contamination results from quenching the samples above 800˚C. This
assumption is based on the composition of the quenching medium, Invoil 705 (C33H34O2Si3), and
the lack of contaminants in samples heat-treated at lower temperatures. To further support this
assumption, basic diffusion calculations for C [164] and Si [165] in Mo at select temperatures were
performed (Table 1). The diffusion distances, L» (Dt)
1/2
, were calculated for a time t = 10s
(approximation of the time required for the sample to cool in the quenching medium). From these
calculations, only surface level diffusion is expected for temperatures below 550˚C; however,
above 800˚C, Si and C are expected to diffuse throughout the 2 um thickness of the sample. These
calculations agree with the APT measurements where C and Si are prominent at 800˚C, and the
measured compositions suggest the formation of MoC and MoSi phases. Therefore, it is assumed
that C and Si contamination occurred during quenching, which altered the final microstructural

96
composition but should have minimal influence on the observed grain sizes and thermal stability
mechanisms.
Table 2: Diffusion distance (L) of C and Si in Mo at different temperatures with time (t) of 10 s

Temperature (°C) D (m
2/s
) L (nm)
C in Mo
550 1.0 x 10
-16
31
800 1.9 x 10
-14
4.3 x 10
2

1000 3.0 x 10
-13
1.7 x 10
3

Si in Mo
550 2.6 x 10
-18
5.1
800 3.2 x 10
-13
1.8 x 10
2

1000 1.3 x 10
-13
1.1 x 10
3

5.5 Grain growth region limit, T=1000˚C
Upon annealing at 1000˚C, the Mo-Au system undergoes a transition from the elongated
morphology seen at 800˚C to a microstructure composed of equiaxed grains, which is observed in
the TEM and EDS maps in Figure 56d. The EDS compositional maps more clearly illustrate the
Au precipitates that appear along Mo GBs, triple junctions, and within grain interiors. At 1000˚C,
it is expected that the overall Au content should decrease as a consequence of the evaporation of
Au at these annealing temperature and vacuum conditions. Thus, several voids are visible in the
microstructure, and the number of Au precipitates has diminished in comparison to the Mo-Au
microstructures observed at lower temperatures. Furthermore, as indicated by the DSC scans, grain
growth is exhibited as the average grain size increased to 190 nm. Specifically, a wide size
distribution for the Au precipitates is seen. Several smaller precipitates are identified within Mo
grains and at triple junctions, but there are also a few Au precipitates that have reached sizes similar
to the Mo grains. Note that given the overall grain size of the sample at 1000˚C, APT was not
performed as TEM and EDS were considered sufficient to characterize the microstructure.

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To further understand the grain growth and precipitation process at 1000˚C, EDS was performed
on a Mo grain and a Au precipitate, where the respective point scans for each region of interest are
shown in Figure 61. From the EDS scan, the presence of strong Au peaks and the absence of any
Mo peaks indicate that the precipitates are mostly pure Au. Likewise, the EDS scan of the selected
Mo rich grain shows strong Mo peaks with no identifiable Au peaks. The detected C and Si peaks
in the scans can be attributed to the quenching process of these samples, as discussed in the
previous section. Therefore, at these elevated temperatures, despite the pronounce changes in grain
size and structure the Mo grains and Au precipitates remain phase separated with no noticeable
changes in composition. These findings suggest that the observed microstructure is formed
primarily through the coalescence and coarsening of the Mo grains and Au precipitates. Based on
growth kinetics of nanophase immiscible systems it is expected that a bimodal growth process
occurs, where the growth of Mo grains is facilitated by GB migration controlled by GB diffusion.
Meanwhile, coarsening of the Au precipitates is promoted by long range volume diffusion which
is similarly controlled by GB diffusion [166].

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Figure 61: EDS scans of a Au precipitate and Mo-rich grain and the corresponding cross-
sectional bright field TEM image of the Mo-Au sample heat treated at 1000˚C. The red arrows
indicate the respective regions where the scans were taken from.
Thus, the sample annealed at 1000˚C exhibited an average grain size ~3 times greater than that
measured at 800˚C. The increase in grain size indicates that the drag effects produced by the Au
precipitates at 800˚C have mostly been overcome at these higher temperatures. Nevertheless,
several nanometer sized Au precipitates are located at GBs, triple junctions, and within Mo grains
and are expected to generate drag forces that counteracted further GB mobility at 1000˚C [64].
Hence, it appears that kinetic stabilization mechanisms are also active at 1000˚C, which limit the
average grain size to 190 nm.

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5.6 Summary
In this study, we have investigated the microstructural evolution of an immiscible binary alloy to
further understand and identify active thermal processes and stability mechanisms as a function of
temperature. More specifically, complementary characterization techniques of APT and TEM
were combined with DSC to explore the contributions of kinetic and thermodynamic mechanisms
during the annealing stages of a Mo-Au NMMs.
Annealing the initial nanostructure to 350˚C results in a multilayer structure with roughening and
intermixing between the Au and Mo-rich layers, indicating the initial stages of degradation of the
nanolaminate structure. Still, no noticeable change in grain size, 9 nm, was observed between these
temperatures. Further annealing to 550˚C led to the recrystallization of the NMMs, producing a
metallic nanocomposite consisting of Mo-rich grains with heterogenous Au solute behavior at the
GBs. The presence of Au clusters and segregation zones along GBs indicated a coupling of kinetic
and thermodynamic stabilization mechanisms which contributed to maintaining an average grain
size of 37 nm. Between temperatures of 550˚C and 800˚C, the prominent thermal process is grain
growth which is facilitated by phase separation of Mo grains and Au precipitates. Furthermore, the
coarsening of Au precipitates through GB diffusion depletes the excess GB solute, promoting a
shift to a predominantly kinetic contribution generated by precipitate pinning. At 800˚C, the Mo-
Au system still retained nanoscale features, with an average grain size of 67 nm. Lastly, at 1000˚C,
the Mo-Au alloy exhibits considerable grain growth leading to the formation of an equiaxed
microstructure with an average grain size of 190 nm. The apparent increase in grain size indicates
that the elevated temperature provides sufficient energy for the grains to overcome the drag forces
generated by the Au precipitates. However, several nanometer sized Au precipitates located at GBs

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and within grain interiors produce an active pinning mechanism that likely mitigated further grain
growth at 1000˚C.
Overall, this study provides a comprehensive guide to the thermal processes and stability
mechanisms responsible for the microstructural changes in Mo-Au NMMs. Annealing of the
NMMs leads to the development of several complex nanostructures that were governed by
recrystallization and grain growth. Furthermore, these thermal processes initiated kinetic and
thermodynamic stability mechanism which appear to be contingent on the segregation and
diffusion behavior of Au atoms. Thus, the utilized methodology in combination with sputtered
NMMs could be useful in identifying the active thermal processes and mechanisms that influence
nanograin stability in other systems.

Copyright 2020 Joel Antunez Bahena
6 Manipulation of interfaces to produce a heterogenous
nanostructured superalloy
A version of the following study is published as an article titled Development of a Heterogeneous
Nanostructure Through Abnormal Recrystallization of a Nanotwinned Ni Superalloy in the
journal Acta Materialia, Volume 195, pages 132-140 (DOI: 10.1016/j.actamat.2020.05.057).
This work explores the development of heterogeneous nanostructured materials (HNMs) through
leveraging abnormal recrystallization, which is a prominent phenomenon in coarse-grained Ni-
based superalloys. The design and development of HNMs have garnered significant interest as
these materials have demonstrated an improved strength-ductility synergy that has otherwise been
lacking in monolithic nanocrystalline metals [100, 103, 167]. HNMs possess nanoscale features
and domains of varying sizes that affect the dominant deformation mechanisms and generate
inhomogeneous plastic deformation that promotes enhanced strain hardening and ductility. To
control heterogeneity from the nanoscale to the macroscale, several processing and fabrication
approaches have been implemented to generate heterogeneous structures. Techniques such as
surface mechanical attrition treatment can produce gradient structures consisting of surface level
nanoscale grains that gradually transition to a coarse-grained core [168]. Additionally, deposition
techniques can offer an effective route for the synthesis of HNMs but with added intricacy [169,
170]. For example, a columnar structured Cu sample was electrodeposited with a consistent
increase in columnar width and twin thickness, producing a hierarchically designed heterogeneous
microstructure by tailoring deposition parameters [102]. However, utilizing electrodeposition
restricts the available systems to single element or binary alloys [171], while using mechanical
treatments localizes the nano-gradient to the surface. Nonetheless, the aforementioned studies do

102
provide the current framework for the design space of HNMs, even though the methodologies
suffer from design limitations.
To date, most HNM studies have focused on generating gradient or bimodal structures, while more
advanced heterogeneous designs, such as those with domains of varying grain morphologies (e.g.
equiaxed, lamellae, columnar, etc.), residual stresses, or chemical composition have mostly been
overlooked [168, 172-174]. In order to synthesize more complex HNMs, there is a need for flexible
techniques that can manipulate several variables and thereby induce both morphological and
microstructural changes. An ideal candidate is magnetron sputtering, as it offers the capability of
depositing nearly any metallic alloy while providing a wide range of deposition parameters that
can be modified to tune the resulting microstructure and morphology. Parameters such as working
distance, working pressure, and target polarization have a direct correlation with microstructural
characteristics, which include feature size, texture, interface type, and residual stresses [175]. In
addition to deposition parameters, material selection is another critical variable as physical
properties can impact the microstructural features and morphology during film growth [106]. For
instance, the stacking fault energy (SFE), which is an intrinsic material property, influences the
propensity to form nanotwins [106]. Depending on the material or alloy selected, the
aforementioned variables can present routes for controlling the microstructural evolution and
potentially introduce heterogeneity in sputtered films [48, 176-178]. An example of this is seen in
age-hardenable Ni-based superalloys which are known to exhibit a bimodal grain size at critical
temperatures and stored energies due to abnormal recrystallization [179-181]. In a study by Wang
et al, it was observed that the annealing of a gradient-strained Ni-based superalloy with an initial
grain size of 2.39 µm resulted in the emergence of several abnormally large grains (ALGs) that
reached sizes up to 208 µm at a critical distance [89]. The ability to tailor abnormal

103
recrystallization in these materials to specific regions can potentially be leveraged with the
microstructural control over nanoscale features and grain morphologies in magnetron sputtering.
The result is an unconventional approach for the manipulation of a sputtered microstructure and
design of a heterogeneous nanostructured superalloy.
In this work, age hardenable Inconel 725 films were synthesized via magnetron sputtering and
subjected to an aging treatment that ultimately induced a transformation from a nanotwinned (NT)
structure to an HNM with a unique and complex gradient grain topology. The aged microstructure
contains domains consisting of columnar NTs, nano-grains, and ALGs with diameters greater than
1 µm, which presents a heterogeneous material with a wide range of grain sizes and morphologies.
Microstructural transitions were highlighted by selecting specific aging times to provide insight
into key thermally activated mechanisms and processes, with an emphasis in the region where
abnormal recrystallization occurs. By combining crystallographic and elemental characterization,
a thorough analysis of the grain evolution and precipitate formation was conducted. Overall, this
work demonstrates a route to generate and tailor a unique sputtered superalloy that can further
expand the current design space of HNMs.
A 7.6 cm Inconel 725 target (Plasmaterials, Inc.) was magnetron sputtered onto 2.5 cm Corning
Eagle 2000 glass substrate using a DC power supply. The nominal composition of the sputtered
film is provided in Table 1, which is in good agreement with the industry standard for Inconel 725
[182]. The sputtering conditions were set to a power of 1500 W, a working distance of 7.6 cm, and
a working pressure of 2 mTorr, in order to achieve a sputtering rate of ~6 nm/s. The resulting film
thickness was 15 µm, which was measured using an XP-2 profilometer. Prior to performing heat
treatments, samples were peeled from the substrate to obtain free-standing films. Individual pieces

104
of film were then aged in a GSL1100X tube furnace (MTI Corporation) at a constant temperature
of 730˚C at intervals of 3, 5, and 8 hours under a vacuum pressure of 5x10
-4
Pa. Upon completion
of the aging sequence, samples were furnace cooled at a rate of ~5˚C/min.
Grain size and morphology evolution as a function of aging time were examined through electron
back scattering diffraction (EBSD), providing a means to examine large cross-sectional regions of
the film. Films were sandwiched in a split specimen mount (Ted Pella, Inc), mechanically polished,
and ion milled using a Leica RES 102 ion polisher, allowing for cross sectional surface preparation.
Scans were carried out in a JEOL IT300HRLV scanning electron microscope (SEM) with an
EDAX Velocity high-speed detector at a working voltage of 30 kV, working distance of 20 mm,
and step size of 20 nm. Furthermore, transmission electron microscopy (TEM) based automated
crystallographic orientation mapping (ACOM) was performed with a LaB6 JEOL 2100 TEM
operated at 200 kV with a step size of 3 nm to resolve nanoscale grains. Preparation of TEM
lamellae was done by the focused ion beam (FIB) liftout method using a JEOL FIB – 4500, where
the region of interest was protected from Ga beam damage through the deposition of a 1 µm-thick
platinum protective layer. ACOM data was constructed with ASTAR NanoMEGAS software, and
both EBSD and ACOM scans were analyzed using OIM software to obtain grain size statistics and
crystallographic orientation information. EBSD data with confidence index lower than 0.1 were
removed and shown as black pixels in the orientation maps. Segregation and precipitation behavior
at the different time intervals were investigated with corresponding electron imaging and elemental
maps gathered from a JEOL 2100F with both STEM and EDS capabilities (Oxford X-MaxN 100
TLE Windowless SDD).

105
6.1 As-Sputtered Microstructure
A detailed analysis of the as-deposited Inconel 725 sample is shown in Figure 1. The cross-
sectional microstructure of the sputtered sample is presented in the bright field STEM image in
Figure 62A showing a fully NT columnar structure where the arrow to the left of the image
indicates the growth direction of the film. It should be noted that the formation of twins is mostly
controlled by the material’s stacking fault energy, which can be reduced with alloying [183]. High
deposition rates can also help to induce favorable twinning conditions in high stacking fault energy
materials such as Ni [43], leading to microstructures resembling the one observed in Figure 62B,
which highlights copious amounts of nanoscale growth twins perpendicular to the growth direction
of the film. The average columnar grain width was determined to be ~60 nm by measuring over
100 grains, where the distribution plot is presented in Figure 62C. A uniform elemental distribution
of key elements is shown in the individual EDS maps scanned from Figure 62A in Figure 62D,
indicating the formation of a solid solution. Further characterization of the microstructure through
ACOM mapping seen in Figure 62E reveals that the film exhibits a strong texture in the (111)
direction as indicated by the IPF triangle to the right of the figure.

106

Figure 62: (A) BF-STEM image showing the nanoscale columnar structure of the as-sputtered
Inconel 725 sample, with (B) bright field TEM images highlighting density of nanotwins. (C)
EDS maps of corresponding STEM image in (A) of Ni, Cr, Ti, Nb, (D) grain width distribution of
columnar grains, and (E) ACOM map showing a strong (111) texture.
In addition to control over microstructural features, magnetron sputtering has the capability to
develop stored energy gradients throughout the film thickness in the form of intrinsic stresses and
defects that can in turn promote heterogeneity [82, 92, 184]. As discussed in the introduction,
certain Ni-superalloys with low-moderate stored energy distribution will undergo abnormal
recrystallization when annealed. This appears to be realized in a specific region of the current

107
sputtered Inconel 725 film in Figure 63 where the transformation from a NT microstructure to one
composed of ALGs is depicted at different aging intervals. Thus, to determine potential sources
for the observed heterogeneous behavior, it is necessary to understand the potential stored energy
landscape of the NT sample in the as-deposited state. A common method in recrystallization
studies to obtain a semi-qualitative representation of this data is through grain orientation spread
(GOS) maps, which examine deviations in local misorientations [88, 89, 116]. Specifically, for
abnormal recrystallization, it has been reported that ALGs emerge in heterogeneous regions with
low-moderate stored energy and GOS, which can promote the formation of a limited number of
stable nuclei that grow at the expense of higher energy grains. [88, 89, 180]. For the purpose of
microstructural visualization, Figure 63A shows a representative cross-section schematic of the
as-sputtered sample, to indicate where GOS analysis was performed. Figure 63B, shows the GOS
map from the ~5 µm representative area outlined in green, which was obtained from the ACOM
data presented in Figure 62E. To the right, distribution plots of the orientation spreads are
displayed along three different regions of the highlighted area. The GOS spread is colored from 0˚
to 5˚, where previous studies have indicated that grains with larger orientation spreads are inclined
to have higher stored energy [185]. The transition in color of the map and shift in distribution plots
point to grains with higher GOS values being located towards the bottom of the region, whereas
grains with lower GOS values are seen towards the top. This gradient would suggest that the stored
energy in the film decreases along the growth direction [109]. The source for the inhomogeneous
distribution of presumed stored energy can be attributed to strain induced by lattice misfit between
the substrate and film, in combination with defects generated by atomic peening that are more
prominent in the early stages of film growth [108, 138, 184]. Furthermore, grain boundary energy
anisotropy induced by grains which deviate from the preferred (111) orientation can generate sharp

108
gradients in stored energy [84, 186]. Overall, it is estimated that the ~4 µm region delineated by
blue lines in Figure 63A exhibits the heterogenous distribution of low-moderate stored energy
required for abnormal recrystallization to occur, as seen in other studies [88]. As such, this
particular region of interest will be referred to as the “AR region” throughout the manuscript.

109
Figure 63: (A) Schematic of a representative cross-section of the as-sputtered film highlighting
the region expected to undergo abnormal recrystallization delineated by blue lines. Growth
direction is shown by arrow to the left. This region includes a green rectangle representing (B)
and a red rectangle representing (C-F). Where (B) shows the GOS map and corresponding
distribution plots, which reveals a gradient in GOS. (C-F) highlight the EBSD scans (top) and
BF-STEM images (bottom) representing the microstructural evolution of the abnormal
recrystallization region in (C) the as-sputtered condition and after aging treatments of (D) 3
hours, (E) 5 hours, and (F) 8 hours. IPF triangle is shown to the bottom right of the scans.
Cross-sectional EBSD and STEM imaging of the AR region are presented for three aging times of
3,5, and 8 hours at 730˚C in Figure 63C-F, which highlight the development of larger grains. While
this conventional EBSD technique does not readily resolve the NT structure or nanoscale features
as indicated by the low-index (black) regions in the scans, it allows for the isolation of ALGs from
the original as-sputtered microstructure. Previous studies have mainly observed abnormal
recrystallization near the γ′ solvus (~1000˚C), however, it was expected that the nanostructured
nature of the sputtered film would cause recrystallization to occur at lower temperatures [84, 187].
Furthermore, the selected temperature would allow for a comparison of the precipitation behavior
with coarse-grained Ni-superalloys, which typically form g' and g'' precipitates under similar
conditions [188], but traditional formation of these phases was not observed. Meanwhile, the
characterization of the microstructures at the selected aging times were guided by structural
changes detected through XRD that is provided in Supplementary Information. The use of interval
aging times allowed for key microstructural transitions to be captured through ACOM and EDS,
enabling the investigation of thermally activated mechanisms and precipitation behavior. Through

110
the combination of these techniques, a detailed analysis of the transformation from a NT alloy to
an HNM is presented in the following sections with an emphasis on the labeled AR region.
6.2 Three hours
After a 3-hour aging treatment, subtle but critical changes to the microstructure are observed. In
comparison to the as-deposited state (Figure 63C), cross-sectional EBSD scans reveal the
emergence of a few larger grains in Fig. 2D. Likewise, the representative STEM image shows a
twin-free grain that has consumed the neighboring columnar structure. The development of these
grains with distinct morphologies was constrained to the AR region that is likely a function of a
stored energy gradient throughout the thickness of the film [84]. Namely, the site-specific changes
in the microstructure are presumed to be driven by the heterogeneous distribution of low-moderate
stored energy in this region, that can drive a nucleation limited recrystallization process [89].
While it is difficult to determine the process responsible for the emergence of ALGs at the current
aging time, the parallels seen with previous studies strongly indicate the initiation of abnormal
recrystallization [88, 89, 180]. Despite the minor changes in the AR region, the overall
microstructure remains predominantly a NT structure as can be seen in the STEM image of the
upper portion of the film in Figure 64A. There are, however, several globular particles that are
decorating the columnar grain boundaries. The EDS maps in Figure 64B, which highlight Ni
(orange) and Cr (green), reveal that these precipitates are rich in Cr. Based on the high Cr content
of this alloy, low ageing conditions and observations from previous studies, it is assumed that the
Cr-rich precipitates dispersed throughout the microstructure are M23C6 carbides [189]. Other
common carbide formers, such as Mo, were also examined, but these elements appear to remain
in solid solution under the current conditions. To further examine the segregation behavior, ACOM

111
was performed to identify the grain boundary character of several columnar grains as solute
segregation has been shown to be highly dependent on boundary energies [190]. The scan
presented in Figure 64C was utilized to identify preferential sites for Cr segregation, where the
grain boundaries in the green bordered insets were analyzed. The red-dotted lines indicate high
angle grain boundaries with misorientations greater that 15°, while the white dashed lines indicate
low angle and coincident site lattice grain boundaries. Cr-rich precipitates were identified and
outlined by dotted yellow lines, which were verified with corresponding STEM images. Over 20
boundaries were examined (inset), where the M23C6 precipitates appear to form preferentially
along high angle grain boundaries, with no identifiable precipitates along low angle or coincident
site lattice grain boundaries. In coarse grained Ni-superalloys, it has been reported that the
enhanced diffusion pathways and increased energy of high angle grain boundaries leads to
preferential segregation of Cr and C [191], which is in agreement with our current observations.
The formation of these Cr-C precipitates could also lead to improved thermal stability by impeding
grain boundary migration, and as a result influence the microstructural evolution at longer aging
times [7]. While it is difficult to predict if the M23C6 precipitates formed prior to the formation of
the ALGs, secondary-phase particles can heavily affect the recrystallization rate [84].
Additionally, given the time and temperature dependency on recrystallization and grain growth, it
can be assumed that there lacked sufficient time for further growth of the newly formed grains.
Overall, the 3-hour aging treatment captures the formation of the early stages of abnormal
recrystallization and Cr segregation.

112

Figure 64: (A) BF-STEM image of the 3-hour aged sample of the upper portion of the film and
(B) corresponding EDS maps highlighting Cr segregation where white scale bars are 1µm. (C)
ACOM scans of representative region, where green insets highlight select GBs with red dotted
lines indicating HAGBs and white dotted lines indicating LAGBs and CSL GBs. Cr precipitates
are encircled with yellow dotted lines. IPF triangle is shown to the right of the scan.
6.3 Five and Eight Hour Heat Treatments: Heterogeneous Nanostructured
Material
Upon aging at 5 and 8 hours, there is a transition to an HNM as illustrated through cross-sectional
EBSD scans of the film shown in Figure 65A. For visualization and comparison purposes, the
areas below and above the AR region are presented in this figure. The three distinct regions shown
for both aging times consists of a (i) NT region containing nanoscale columnar grains with a high

113
density of stabilized nanotwins; (ii) the AR region where ALGs have formed and undergone grain
growth; and (iii) a nanocrystalline region where equiaxed grains have formed. These
microstructural transitions are highlighted in Figure 65B, where the feature sizes of the as-
sputtered, 5 hour, and 8 hour samples are plotted along the thickness of the film. These
observations are consistent with the expectation that the sputtered film exhibits a residual stress
and stored energy gradient that would govern the formation of distinct microstructural regions
[89]. The three regions comprise an overall complex microstructure with each region undergoing
discrete thermal processes, and in this study the microstructural analysis at these select aging times
focuses primarily on the AR region (ii) although regions (i) and (iii) will also be briefly discussed.
For region (i), the lower stored energy towards the upper portion of the film does not sufficiently
drive recrystallization, leading to a similar NT structure as seen after the 3-hour aging condition.
Other NT Ni alloys, such as Ni-Mo-W, have shown similar thermal stability under comparable
temperatures due to the low enthalpy stored in twin boundaries and Zener drag forces generated
from secondary phase precipitates [176]. Hence, the high density of nanotwins and Cr-carbides
along grain boundaries of the aged sputtered Inconel 725 sample similarly stabilize the NT
structure in region (i). Despite the lower stored energy, a few ALGs do appear in this region, as
seen in the EBSD scans in Figure 65, and are likely associated with localized stored energy
variation generated during the sputtering process. In contrast, region (iii), which should exhibit the
highest stored energy due to its vicinity to the bottom of the film, has a sufficient density of stable
nuclei sites that induce complete recrystallization and a smaller grain size [89, 116]. As a result,
nanoscale equiaxed grains are formed with Cr-carbides located along grain boundaries, potentially
minimizing further growth. As can be seen from the insets in Figure 65, the microstructural

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features in region (i) and (iii) showed minimal changes between the aging treatments of 5 and 8
hours.
In the labeled AR region (ii), which as seen in Figure 65 spans ~4 µm and is observed in between
region (i) and (iii), the NT structure is quickly replaced by ALGs. The EBSD scans and grain size
distribution plots in Figure 65A show differences in morphology and an increase in average grain
size from ~460 nm and ~715 nm after 5 and 8 hours, respectively. To investigate the effects of
stored energy and the microstructural changes in the AR region (ii) at these two aging times,
correlative STEM, ACOM, and EDS maps at each interval are presented in the following sections.

115

Figure 65 (A) EBSD maps of cross-section of heterogenous structured achieved at aging times of
5 and 8-hours (top) with respective grain size distributions of grains in the AR region (ii)
(bottom). Insets highlights characteristic microstructures of the region (i).
6.4 AR Region: Five Hour Treatment
In order to examine the microstructural and precipitate evolution in detail, ACOM and EDS
analysis of two different sections in the AR region after the 5-hour aging treatment are presented
in Fig. 5 and Fig. 6, respectively. Specifically, Figure 66 presents a corresponding STEM image

116
(Figure 66A) and ACOM map (Figure 66B) which give insight into the grain morphology and
orientation of the observed microstructure, where the arrow on the left indicates the growth
direction of the film and is also an indication of the direction of decreasing stored energy. The IPF
triangle on the right reveals the crystallographic orientations of the grains seen in the ACOM scan.
In Figure 66A-B, several grains with random orientations and different morphologies compared to
the surrounding (111) oriented grains are outlined with dashed lines. The distinct characteristics
between the highlighted and non-highlighted grains provide evidence of two thermally activated
processes invoking the increased grain size and formation of ALGs. The changes in orientation of
the highlighted grains suggests that a nucleation-based recrystallization is driving the growth of
select grains. As previously indicated by the GOS map in Fig. 2A, this region should exhibit areas
of moderate stored energy, which restricts the number of stable nuclei that can form and grow,
analogous to the hypothesized theory of abnormal recrystallization in other Ni-superalloy studies
[116, 179]. The newly formed low stored energy grains then grow at the expense of surrounding
pre-existing grains with higher stored energy, leading to the eventual emergence of ALGs [116].
Additionally, the new grains generated through recrystallization appear to promote the formation
of annealing twin boundaries, where select boundaries are outlined in red in Figure 66B. Similar
observations have been reported in studies of thermo-mechanically processed Ni alloys, where
twinning occurs during the growth process after recrystallization [117, 119, 192].

117

Figure 66: (A) BF-STEM images and (B) corresponding ACOM scans of the AR region in the 5-
hour aged sample, highlighting ALGs governed by different thermal processes. Randomly
oriented recrystallized grains are highlighted by dashed black and white lines. White scale bar is
1 µm and IPF triangle is shown to the right of the scan.
Concurrently, several ALGs with a slightly rectangular morphology are observed in Figure 66A,
with the ACOM map in Figure 66B revealing an evident (111) orientation, identical to that of the
as-deposited film. The orientation and observed morphology suggest that these grains develop
through the detwinning and coarsening of columnar grains in a direction perpendicular to the
growth direction of the film [48, 193]. Considering that the NT structure in region (i) did not
undergo similar transitions, the source of this increased twin and grain boundary mobility can be
attributed to the low-moderate stored energy specific to the AR region. Thus, the microstructural
changes denote a strain induced boundary migration (SIBM) process, facilitated by variations in
stored energy causing boundaries of low energy grains to move towards high stored energy

118
regions. As a result, low stored energy regions are generated with the same crystallographic
orientation as the parent grain [86]. While the driving force for recrystallization and SIBM are
similar, the differences in resulting microstructures are considerable as seen by the distinct types
of observed grains. Notably, both processes are dependent on the initial stored energy state, where
SIBM typically occurs at lower stored energy conditions than nucleation-based recrystallization
[87].
The heterogeneous distribution of stored energy not only leads to variation in grain morphology
and orientations but also in precipitation behavior. Figure 67 highlights the distribution of
precipitates including corresponding STEM and EDS maps. Throughout the entire region, M23C6
precipitates are observed in addition to rafted structures and δ precipitates, where arrows in Figure
67A mark individual examples of each. The rafted structure as indicated by the purple arrow, is
primarily observed in the upper portion of the film in the square shaped grains the have undergone
SIBM. This type of structure is typically observed in single crystal Ni-superalloys [194, 195] and
is the product of the discontinuous precipitation of g and g' that can be driven by local reductions
in chemical free-energy, decreases in stored energy, or grain boundary migration [196, 197]. Thus,
it can be deduced that the observed rafting was likely promoted during the SIBM process. In
contrast, the recrystallized ALGs towards the bottom of the film contain elongated plate shape
particles across the width of the grain as indicated by the blue arrow. The EDS maps in Figure 67B
show that these particles are rich in Nb and Ti indicating the formation of the Ni3(Nb,Ti) δ phase.
This was further validated with an SAED pattern from the [111] matrix zone axis shown in the
inset, revealing the presence of [100] δ superlattice reflections [198]. While δ phase has been
observed in Ni-based superalloys under similar experimental temperatures, it typically requires
much longer aging times to form. For example, an Inconel 718 sample aged at 750˚C required

119
100h for noticeable precipitation of δ phase [199]. However, it has been demonstrated that
increasing the stored energy through plastic deformation with techniques such as cold rolling can
accelerate precipitation kinetics at lower temperature [200]. Therefore, it is likely that the higher
stored energy that governed recrystallization also promoted the rapid precipitation of δ phase.
From the STEM images it appears that precipitation of δ phase occurs intragranularly, but through
comparing the ACOM scan and EDS maps it is revealed that many of the δ precipitates are located
along annealing twins. Due to the low grain boundary energy associated with twin boundaries,
precipitation at higher energy sites would be expected, but several studies have identified twin
boundaries as nucleation sources for δ precipitates as a function of the good matching between the
twin boundary plane and δ habit plane [198, 199, 201]. Moreover, the formation of δ phase is
typically associated with a transformation from metastable g'' which was not identified in this
study. Direct phase transformation of a similar phase has been observed in a NT Haynes 242
sample upon aging, which was attributed to the presence of HCP-like ordering [183]. Nonetheless,
an HCP structure was not apparent from our observations, thus, it is unclear whether the δ phase
is the byproduct of phase transformation or direct precipitation.

120

Figure 67: (A) BF-STEM image of the 5-hour aged sample illustrating the heterogenous
precipitate and microstructural behavior in the AR region with inset SAED pattern confirming
the presence of δ precipitates. (B) EDS maps of Ni, Cr, Ti, and Nb highlighting location of
different precipitates where white scale bars are 1 µm.
Overall, the 5-hour treatment causes both a SIBM and a nucleation-based recrystallization process
to occur in the AR region of the film, giving rise to the formation of different types of ALGs with
distinct precipitate behavior. The stored energy advantage of ALGs has been shown to overcome
Zener pinning forces and the same is shown in this study, as both types of grains exhibit growth
despite the high density of Cr-C precipitates decorating the grain boundaries [116]. It is still
expected that the pinning precipitates would promote sluggish kinetics that lower the overall grain
growth rates [62]. Furthermore, unconsumed NT columnar grains (black arrow) are seen
throughout this region suggesting that certain columnar grains exhibited initial low stored energy
as they are unaffected by either thermal process. Thus, the stored energy profile appears to heavily

121
influence the AR region, producing a more complex overall microstructure with higher levels of
order.
6.5 AR Region: Eight Hour Treatment
After the 8 hour treatment, further coarsening of recrystallized ALGs is observed throughout the
AR region with a more homogeneous distribution of precipitates as detailed in Figure 68, which
presents similar STEM, EDS, and ACOM analysis as performed in the previous section. In Figure
68A, a STEM image depicts an area between region (i) and the AR region (ii), where the transition
is highlighted by a dashed white line, and the arrow to the left indicates the growth direction of the
film. From the image, it can be seen that the equiaxed grains have increased in size and are more
prominent towards the upper portion of the AR region. Accompanied with the EBSD scan from
Figure 65, it is apparent that this morphology is more pronounced in this region of the film,
indicating that the increased aging time has allowed for further growth and the formation of new
recrystallized grains. NT columnar grains are also still observed in this region as denoted by the
black arrow in Figure 68A. In terms of precipitates, a higher density of δ phase are present as a
result of the increased number of recrystallized ALGs. The EDS maps in Figure 68B show the
elemental distribution of key elements of the corresponding area, where the δ phase can be
identified by a Nb enrichment. M23C6 are still also observed, where the Cr map in Figure 68B
illustrates the distribution of these precipitates along grain boundaries. Furthermore, the rafted
structure is no longer observed after the 8 hour treatment as the SIBM grains appear to have been
consumed. The crystallographic orientation data of the region enclosed by the orange box is
presented in the ACOM map in Figure 68C. Similar to the observations after the 5-hour aging
treatment, the recrystallized ALGs have formed annealing twins, where select twin boundaries are

122
outlined in red. Coincidentally, a large fraction of δ precipitate are observed on or near annealing
twin boundaries.

123
Figure 68: (A) BF-STEM Image, corresponding (B) EDS maps of Ni, Cr, Nb, and (C) ACOM
maps from the outlined (orange) region in (A) with inset of IPF triangle of the 8-hour aged
sample. White scale bars are 500 nm.
The observed microstructure points to the continued evolution and growth of the recrystallized
grains as the primary source for the increased size of ALGs at 8 hours. This shift is likely attributed
to a stored energy difference, as it is anticipated that the recrystallized grains would exhibit lower
energy and preferentially grow at the expense of SIBM grains. Meanwhile, the presence of NT
columnar grains reaffirms that certain grains are deposited at an initial low stored energy state. As
indicated by the continued growth of ALGs, the extra aging time has allowed for sufficient
mobility to overcome sluggish kinetics due to drag forces generated by secondary phases at grain
boundaries, such as M23C6 and δ precipitates [62, 201]. Consequently, this has also led to an
increased density of δ phase that, as previously discussed, is driven by stored energy and the
formation of annealing twins. Through comparing STEM, EDS, and ACOM, several δ precipitates
are seen on or near annealing twin boundaries, providing further evidence of that as a preferential
precipitation site. Effectively, the observed microstructure after the 8 hour treatment presents a
natural progression from the one seen after 5-hours. Previous studies have shown that during
abnormal recrystallization, there is a time dependency until the eventual mutual impingement of
select equal energy grains. This in turn, leads to a more homogeneous distribution of ALGs, as the
lowest stored energy grains will dominate the resultant microstructure [89, 116]. Likewise, the
recrystallized ALGs with δ precipitates comprise the majority of the AR region.

124
Altogether, the observed microstructures of the aged sputtered samples demonstrate a significant
contrast with that of traditional Inconel 725 that could lead to attractive mechanical behaviors [182,
202]. The absence of conventional g′ and g′′ is especially notable. As previously discussed, it is
expected that the increased stored energy in the AR region (ii) facilitates either an accelerated or
a direct formation of δ phase along annealing twins and grain boundaries. The presence of
intergranular δ precipitates have been shown to lower the yield strength of aged Ni-superalloys
while improving ductility [203]. However, given the high density of nanotwins in region (i), which
are known to promote exceptional strength [176, 204, 205], the absence of traditional g' or g''
precipitates is not expected to diminish the properties of the presented alloy. Similarly, the
formation of a heterogenous structure containing three distinct domains with variable grain size,
precipitate formation, and morphologies could lead to a strength-ductility synergy that has been
observed in other HNM studies [103]. The NT structure of region (i) and nanocrystalline structure
in region (iii) would significantly strengthen the material, whereas the AR region (ii) could provide
a ductile interior that can accumulate dislocations during deformation [100]. It is possible that
these attributes would also translate to enhanced tribological properties as the ductile core could
suppress strain localization and crack propagation [206]. In this regard, it would be interesting to
examine the structure-property relationship of the observed heterogeneous microstructure after the
5 and 8 hour heat treatments in future studies.
6.6 Summary
In this study, a novel approach for producing a HNM was developed, inspired from the
heterogeneous behavior observed in CG Ni-based superalloys and translated to the nanoscale by
annealing a sputtered NT structure with heterogeneous distribution of stored energy. The resulting

125
microstructure consisted of three distinct regions that were influenced by the stored energy profile
of the film, presenting a unique combination of grain morphologies and sizes that has not been
previously observed. A graphical summary of the microstructural evolution during the aging
process is depicted in Figure 69. While the three regions are presented, the complex nature and
heterogeneous evolution of the microstructure prompted the current study to mainly focus on the
thermal processes governing the changes in the AR region (ii). The microstructural analysis
revealed that development of ALGs in this region is induced by both a nucleation-based
recrystallization and a SIBM process. Furthermore, different grain sizes and precipitation behavior
are observed after 5 and 8 hours, demonstrating the potential to tailor the AR region and the overall
HNM depending on aging time. Future studies are needed to further elucidate the processes
occurring in region (i) and (iii) as well as to investigate the effects of domains with varying
morphological features and size on the mechanical properties.

126
Figure 69: Schematic overview of the microstructural evolution of the sputtered Inconel 725 film
at different aging times that presents the changes in grain morphology, precipitate formation,
and development of a heterogenous microstructure. The arrow to the left indicates the growth
direction of the film and the red dotted box highlights the AR region.
Overall, this study validates the proposition of combining the flexibility of magnetron sputtering
with a Ni-based superalloy to produce a unique HNM which can be further modified. Through
controlling factors such as the distribution of stored energy during the deposition process, this
approach could be utilized to tailor different domains and generate more complex microstructures.
In addition to the phenomenon of abnormal recrystallization being prevalent in Ni-based
superalloys, it has also been observed in Co-based superalloys [207] which provides opportunities
to explore different microstructural designs and alloy combinations. Thus, the current
methodology provides a potential path to expand the current design space of nanomaterials.

127
7 General conclusions and future research outlook
Despite the tremendous promise of nanostructured materials, unpredictable grain growth behavior
and limited ductility hinders utilization in many advanced engineering applications. However, it is
clear that controlling the interfaces is an encouraging route to resolve these drawbacks. In this
work several tailored nanostructured materials were synthesized by magnetron sputtering and the
role of interfaces, thermal processes, and stabilization mechanisms in relation to microstructural
evolution were assessed. Through extensive characterization of the thermal transformations,
insights into the development of more stable and practical nanostructured materials was obtained.
One such approach in this dissertation work focused on GBE materials, and the importance of
initial twinned microstructure of four Cu alloys was examined. Through isolating the effect of twin
spacing, a threshold to the stabilizing effects of nanotwins was observed. At high densities, the
excess energy generated from twin boundaries was shown to drive the early onset of abnormal
grain growth and recrystallization. This is contrary to previous studies which have theorized that
the relatively low energy associated with twin boundaries should resist twin and grain boundary
mobility at elevated temperatures. Moreover, a random texture induced by tilted twin boundaries
can mitigate abnormal grain growth and promote a stepwise growth behavior, a behavior that was
not observed in films with strong (111) textures. Thus, through tuning the twin density and texture,
it may be possible to design microstructurally stable nanotwinned materials.
In addition to energy and texture, the chemistry of interfaces is equally important for stabilizing
nanostructured material. While it has been assumed that kinetic and thermodynamic mechanisms
are independent, through characterization of the microstructural evolution of a Mo-Au NMM, it
was shown that these mechanisms can occur simultaneously. Specifically, from temperatures of

128
350˚C to 1000˚C, a transition from a primarily thermodynamic contribution to a dominating kinetic
mechanism was observed. It was determined that recrystallization facilitated the segregation of Au
to the interfaces, which dictated the microstructural transformation observed at the different
annealing temperatures. Furthermore, this study demonstrates the potential to use NMMs and
GBSE as a route to design stable nanocrystalline systems.
Through the combined understanding of GBE, GBSE, and initiation of thermal processes, a novel
approach for producing a Ni-superalloy nanostructured material with a unique grain topology. The
resulting microstructure consisted of three distinct regions that were influenced by the stored
energy profile of the film, presenting a combination of grain morphologies and sizes that has the
potential to improve strength-ductility synergy. Many of the observations made from this study
also provided a deeper understanding of abnormal recrystallization and precipitation behavior of
Ni-superalloys. The formation of abnormally large grains was attributed to both a strain induced
boundary migration and a limited nucleation-based recrystallization process. More so, the
unconventional precipitation pathways shown to be dependent on ageing time, stored energy, and
possibly the density of interfaces. Ultimately, an approach for designing very tunable HNMs was
achieved, but the impacts of this study are farther-reaching.
The studies detailed in this dissertation investigate the relationship between interfaces and thermal
processes that directly affect the microstructural evolution of nanostructured materials. However,
there are still many unanswered scientific and engineering questions, which lends to further
exploration building off this work. Of considerable interest is the grain boundary character in the
context of discussing segregation of solute elements. While it has been suggested that CSL grain
boundaries should resist solute formation, current studies do not report any clear trends [208]. It is

129
known that CSL theory takes into consideration three of the five macroscopic degrees of freedom,
and it has become increasingly important to also consider the remaining two, which relate to the
grain boundary plane [22]. Boundaries with similar orientations but with different boundary
planes, can vary significantly in energy and effect segregation behavior [162]. In order to design
specific microstructures with improved properties, knowledge of the segregation behavior at
specific interfaces is crucial. It has already been observed that a sputtered film with primarily high
angle GBs towards the top and low angle GBs towards the bottom, demonstrates the activation of
different thermal processes and mechanisms in the two regions [104]. Further investigation
through the use of advanced complementary techniques such as APT and ACOM/t-EBSD to obtain
atomic precision of composition and character of interfaces. Furthermore, this type of study could
be expanded to understand segregation behavior or cluster evolution along specific interfaces.
Doing so could provide a template for determining the most advantageous boundaries, depending
on the desired stabilization mechanism.
While the sputtered films in this study were annealed and characterized at different intervals, it
would be advantageous to examine the thermally activated microstructural evolutions in real time.
Through in-situ TEM characterization, it would be possible to confirm assumptions made from ex-
situ studies while also identifying intermediate steps involving structural, morphological, or
chemical transitions. Building off of the abnormal thermal phenomena observed in this study,
insights into nucleation sites and characterization of grains that are likely to have a competitive
growth advantage can be drawn.
In addition, several unique precipitate pathways were identified in the processes of developing an
Inconel 725 HNM. In-situ characterization would allow for determination of the key mechanisms

130
controlling complex phase evolution that are still not yet entirely understood [209]. There is
speculation that twin boundaries play a crucial role in accelerated or indirect precipitation of delta
phase, which could be clarified with such characterization methods. In understanding phase
evolution pathways, there are opportunities to tailor the properties of both nanostructured and
coarse-grained superalloys.
Furthermore, the current studies examine how thermally activated transformations can be used to
manipulate microstructures synthesized via magnetron sputtering. Specifically, a nanotwinned Ni-
superalloy was leveraged with abnormal recrystallization to produce of a complex gradient grain
topology, which can be used as a route to further expand the design space of HNMs. To date,
efforts have been placed on understanding the effects of grain size gradients and order of
heterogeneity for producing optimal strength-ductility synergy [167]. Many of the current
synthesis and processing techniques to produce HNMs can be restrictive, but the methodology
proposed in this study provides a flexible means to alter the distribution of grain size, morphology,
and precipitate behavior. Thus, by controlling deposition parameters, a tailored approach to
achieve different domains and generate more complex microstructures could be realized. As such,
it would be possible to produce materials to match designs from computational work to optimize
mechanical properties [210]. While the initial design approach would be limited to materials
susceptible to the abnormal recrystallization, an improved understanding of recrystallization could
expand the material selection.
In conjunction with tailored HNMs, it is also necessary to perform mechanical properties to assess
structure property relationships. Nanoindentation and micro-tensile testing would allow for the
measurement of both local and global mechanical properties of these films, which if compared

131
across different topologies could facilitate tailoring of strain hardening effects. Specifically, the
contributions from localized microstructural features such as grain size, precipitate formation, and
grain morphology across the different regions can be determined through cross-sectional
nanoindentation. Meanwhile, micro-tensile tests would provide a method to evaluate the strength-
ductility synergy. There are still many areas of interest to explore, but the novel approach presented
in this study provides an initial basis for designing nanostructured materials with optimized
mechanical properties.
Overall, the thermal studies conducted in this thesis present 1) the impact of initial twinned
microstructure on grain growth behavior of NT metals 2) a deeper understanding between the
interplay of kinetic and thermodynamic mechanisms in NMMs, and 3) a unique approach to
manipulate interfaces and thermal processes to generate a unique heterogenous nanostructure. The
aforementioned studies have opened exciting scientific areas to explore and provide a clear
direction for future investigation on the structure and behavior of nanostructured materials.

132
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Copyright 2020 Joel Antunez Bahena
Appendix A: Additional APT characterization of annealed NMMs
Additional APT was performed for the Mo-Au system as well as other annealed binary NMMs.
In this appendix, additional analysis for these materials are presented, which provide insights into
the processes and mechanisms involved of the microstructural evolution of these systems.
For the Mo-Au system, a second APT map of the sample annealed at 800˚C, discussed in Section
5, is presented below. APT is a site-specific technique and this additional atomic map reinforces
the phase separation of Mo and Au at these temperatures, which points to a transitions to a kinetic
stabilization contribution.

Figure 70: APT maps of a Mo-Au sample annealed at 800˚C, highlighting the phase separation
occurring of Mo (red) and Au (yellow). Top left to right: Combined APT maps of Mo and Au,
Mo Au
ROI
-20
0
20
40
60
80
100
0 5 10 15 20 25
Concentration (Atomic %)
Distance (nm)
Concentration Profile ROI
Au %
Mo %

146
APT map of Mo, and APT map of Au. Below is a concentration plot extracted from the ROI
indicated by blue line.
Another system that was investigated was a Hf-Ti NMM which has been investigated in previous
studies [72, 82]. An APT map of a tip obtained from this system annealed at 500˚C is shown in
Figure 71. At this temperature, the multilayer structure is still preserved which is seen by Ti layers
in black and supported by the concentration profile along the Z axis. Intermixing is present as ~
30 at.% Hf is detected in the Ti-rich regions. Furthermore, Figure 72 provides a top-view
perspective of the tip in Figure 71, where segregation of Ti to the grain boundaries is visible. The
enrichment is quantified in the corresponding concentration profiles where the values reach around
20 at.% Ti. The segregation of Ti to the grain boundaries is expected to assist in stabilizing the
microstructure thermodynamically. No further APT characterization was performed on the Hf-Ti
system.

0
10
20
30
40
50
60
70
0 20 40 60 80 100
Concentration (Atomic %)
Distance (nm)
1D Concentration Profile - along Z axis
Ti %
Hf %
Ti
Hf
20 nm
6

147
Figure 71: APT map (left) of a Hf-Ti NMM sample annealed at 500˚C, highlights the multilayer
structure and segregation of Hf (black) along interfaces. 1D concentration profile (right) along
the z axis displays the elemental fluctuations.

Figure 72: APT map (left) of a Hf-Ti NMM sample annealed at 500˚C shows a top view
perspective of Figure 71, where 1D concentration profiles (rights) were extrapolated from
different grain boundaries, showing an Hf enrichment.
Lastly the Hf-Ta system [104] was examined at temperatures of 550˚C and 1000˚C through APT.
Figure 73 specifically highlights the atomic mapping of a Hf-Ta NMM annealed at 550˚C.
Visually, the APT maps shows that the multilayer structure is still intact and the concentration
profile demonstrates sharp transitions between the Hf and Ta layers. However, as seen in Figure
74, there are regions where the multilayer structure has broken down either through a diffusion-
based process or, as indicated by large HF rich regions, a recrystallization progress.
20 nm
0
10
20
30
40
50
60
70
80
0 1 2 3 4 5 6 7 8 9
Concentration (Atomic %)
Distance (nm)
Concentration Profile ROI#1
Ti %
Hf %
0
10
20
30
40
50
60
70
0 1 2 3 4 5 6
Concentration (Atomic %)
Distance (nm)
Concentration Profile ROI#2
Ti %
Hf %
ROI#2
Ti
Hf
ROI#1
7

148

Figure 73: APT map (left) of a Ta-Hf NMM sample annealed at 550˚C showing a preserved
multilayer structure of Ta (purple) and Hf (yellow), and is further confirmed by the
concentration profile (right)
20 nm
Ta
Hf
0
10
20
30
40
50
60
70
80
90
100
0 10 20 30 40
Concentration (Atomic %)
Distance (nm)
Concentration Profile ROI
Hf %
Ta %
ROI
10

149

Figure 74: Rotated view of the APT map from Figure 73 of a Ta-Hf NMM sample annealed at
550˚C which highlights the breakdown of the multilayer structure and shows the formations of
HF rich regions.
At 1000˚C, there are regions where the multilayer structure has degraded, and it is presumed that
a phase separation process has occurred as illustrated by Figure 75. Moreover, several Hf clusters
are seen in the Ta rich regions as shown in Figure 76. From the data it is unclear whether these
clusters are within the Ta matrix or along interfaces, but it is likely that the Hf-rich clusters provide
a drag force in stabilizing the microstructure at these elevated and intermediate temperatures.
ROI
20 nm
Ta
Hf
0
10
20
30
40
50
60
70
80
90
100
0 10 20 30
Concentration (Atomic %)
Distance (nm)
Concentration Profile ROI
Hf %
Ta %
11
Specimen A

150

Figure 75: APT map (left) of a Ta-Hf NMM sample annealed at 1000˚C shows the a non-
multilayer structure that has phase-separated as indicated by the sharp transitions in the
concentrations profile (right).

Figure 76: Cross-section of APT map from Figure 75 depicting the formation of Hf clusters
forming in the Ta rich regions.
0
10
20
30
40
50
60
0 2 4 6 8 10 12 14
Concentration (Atomic %)
Distance (nm)
Concentration Profile ROI
Hf %
Ta %
10 nm
ROI
Ta
Hf
13
Specimen A
Hf (isosurface 18 at.%)
10 nm
0
5
10
15
20
25
30
35
40
45
50
0 2 4 6 8 10 12 14
Concentration (Atomic %)
Distance (nm)
Concentration Profile ROI
Hf %
Ta %
Ta
ROI
14

151
Appendix B: Complementary characterization of sputtered Inconel
725
This appendix contains XRD, ACOM, STEM, and in-Situ TEM characterization of an as-sputtered
Incone1 725 alloys used in this study in addition to some preliminary work.
Table 3: Sputtering conditions for Nanotwinned Inconel 725
Sample
Name
Power
Density
(W/cm
2
)
Working
Pressure
(mTorr)
Working
Distance
(inches)
Deposition
Rate
(nm/s)
Thickness
(μm)
Holder
Temp.
Inconel725#4 33 2 3 6 15 16˚C

In order to strengthen Ni-based alloys, a three-step aging treatment is performed as seen in Figure
77. The initial heat-treatment is a solution anneal to dissolve all elements into the matrix followed
by a rapid quench. The following annealing steps initiates the precipitation and subsequent aging
of gamma prime and gamma double prime secondary phases, which typically provide the best
balance of strength and ductility. For the sputtered films, the initial solution anneal was skipped as
sputtering has been shown to force solid solutions. Furthermore, the XRD in Figure 78 guided the
aging times, where aging was only performed after 3, 5 and 8 hours.

152

Figure 77: The heat-treatment process for conventional Inconel 725. For sputtered films, solid
solution step was skipped, and samples were annealed at 3, 5 and 8 hours.
Several XRD scans were performed on sputtered films on Corning glass substrates where peaks at
4, 5, and 8 hours indicated the presence of structural changes, such as the formation of secondary
phases.

153

Figure 78: XRD data of samples annealed at 730˚C at various aging times, collected on a
Rigaku Ultima IV diffractometer. The boxes highlight structural changes.

In addition to the ex-situ TEM and ACOM performed in Chapter 6, preliminary in-situ heating
experiments were conducted. The lamellae were annealed to 700˚C using an in-situ TEM heating

154
holder. The initial testing reveals that this abnormal recrystallization occurs even when thin film
effects are present as snapshots were taking of the AR region and towards the top of the film (NT
region). ACOM was performed upon completion of the in-situ heat treatment which highlight
some of the microstructural transitions. Ultimately, this type of methodology could elucidate the
mechanisms involved in abnormal recrystallization.

Figure 79: TEM snapshots of the beginning and end of an in-situ heating experiment performed
on sputtered Inconel 725 along with ACOM showing some of the grains exhibiting a texture
change.
Preliminary mechanical testing was performed on the cross-section of the sputtered films in the
form of nanoindentation. Tests were only performed on the as-sputtered film and after being
AR Region
2 Step HT
NT Region
500 nm

155
annealed at 3 hours, but interesting initial observations are seen. For the as-sputtered film an
increase in hardness is seen from the top of the film towards the bottom. This can be explained by
the gradient in residual stresses and stored energy. At 3 hours, a drop-in hardness is seen in the AR
region which, from observation in Chapter 6, should be populated with smaller recrystallized
grains.

Figure 80: (Left) SEM image of cross-section of as-sputtered Inconel 725 with the various
indentations caused by nanoindentation. (right) Hardness profile from the top of the film towards
the bottom in microns.

9.8
10
10.2
10.4
10.6
10.8
11
0 2 4 6 8 10 12
Hardness (GPa)
Distance from top of film
Average Hardnesss
9.5
10
10.5
11
11.5
12
0 2 4 6 8 10 12
Hardness (GPa)
Distance from top of film
Average Hardness
3 Hr
AS

156
Appendix C: Summary of sputtered samples
Table 4: Sputtering conditions of various samples

Sample Name
Power
Density
(W/cm
2
)
Working
Pressure
(mTorr)
Working
Distance
(inches)
Deposition
Rate
(nm/s)
Thickness
(μm)
On/Off
Cycle
(s)
Holder
Temp
Al5456 – High 33 2 3 7 15 10/100 16˚C
Al5083 – Mid 16 2 3 3.5 15 10/100 16˚C
Al5083 – Low 7.7 2 3 1 15 10/100 16˚C
Inconel600#26 33 2 3 8 15 100/0 16˚C
Inconel 718#2 33 2 3 6 15 100/0 16˚C

Abstract (if available)

Abstract Nanostructured alloys have many attractive mechanical, chemical, and thermal properties that are largely attributed to a high density of interfaces. However, these same interfaces can drive microstructural instability at relatively low temperatures, resulting in the annihilation of the nanostructure and subsequent loss of associated properties. Thus, there is significant interest in investigating routes to prevent or mitigate thermal processes such as grain growth or recrystallization. To date, two routes have demonstrated some level of effectiveness: (i) introducing a high density of low energy boundaries that resist grain boundary mobility and (ii) utilizing solutes to decorate interface that restrict grain boundary motion through either kinetic or thermodynamic mechanisms. Despite current progress, the thermal processes and mechanisms are not well understood, particularly the intermediate microstructural transitions that can provide key insights. Thus, this dissertation examines the thermal evolution of several tailored nanostructured alloys in order to examine the effects of interfaces and chemical composition on thermal stability, and also as a means to develop new nanocrystalline systems. Overall the results from these investigations culminate in three important findings: 1) An interplay between kinetic and thermodynamics is shown in a Mo-Au nanometallic multilayer system which dictates microstructural transformations at different temperatures, 2) twin boundaries at high densities, were shown to initiate the early onset of recrystallization and abnormal grain growth in nanotwinned Cu alloys, however, abnormal grain growth can be mitigated by the presence of tilted twins that can induce a random texture, and 3) stored energy gradients in a nanotwinned Ni superalloy, in the form of defects and intrinsic stresses, can lead to the heterogenous activation of thermal processes that results in a unique combination of grain size and morphology. Overall, this dissertation contributes to the understanding of the thermal processes and mechanisms that influence the evolution of interfaces and microstructures and provides insight into developing stable and unique nanostructured materials.

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Creator Bahena, Joel Antunez (author)

Core Title Understanding the formation and evolution of boundaries and interfaces in nanostructured metallic alloys

School Viterbi School of Engineering

Degree Doctor of Philosophy

Degree Program Mechanical Engineering

Publication Date 07/27/2020

Defense Date 06/22/2020

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

Tag nanostructured materials,nanotwinned,OAI-PMH Harvest,sputtering,thermal evolution

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Advisor Hodge, Andrea M. (committee chair), Branicio, Paulo (committee member), Luhar, Mitul (committee member)

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