Microstructural Evolution during Heat
Treatment and High Strain Rate
Deformation of an Fe-10Ni-0.1C Steel
By Ian Harding
Master of Science, Brown University, Providence, RI, 2015
Bachelor of Science, Temple University, Philadelphia, PA, 2013
A dissertation submitted to the School of Engineering in partial fulfillment
of the requirements for the degree of Doctor of Philosophy
Brown University
Providence, Rhode Island
May, 2019
© Copyright 2019 by Ian Harding
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This dissertation by Ian Harding is accepted in its present form by
the School of Engineering as satisfying the dissertation
requirement for the degree of Doctor of Philosophy.
Date________________
_________________________
Prof. K. Sharvan Kumar, Advisor
Recommended to the Graduate Council
Date________________
_________________________
Prof. Clyde Briant, Reader
Date________________
_________________________
Prof. C. Cem Taşan, Reader
Approved by the Graduate Council
Date________________
_________________________
Andrew Campbell, Dean of the
Graduate school
iv
Curriculum Vitae
Education
- Ph.D. in Materials Science and Engineering 2019
Brown University, Providence, RI, USA
Advisor: Prof. Sharvan Kumar
- Sc.M. in Materials Science and Engineering 2015
Brown University, Providence, RI, USA
- B.S. in Mathematics and Physics 2013
Temple University, Philadelphia, PA, USA
Conference Presentations
Harding I, Mouton I, Gault B, Raabe D, and Kumar S, “The Stability of Precipitated
Austenite in Fe-10Ni-0.5Mn-0.1C Steel”. TMS 148th
Annual Meeting and Exhibition,
March 2019.
Harding I, Kumar S, “Dynamic Deformation Behavior of an Fe-Ni-C High Strength,
High Toughness Steel”. TMS 148th
Annual Meeting and Exhibition, March 2019.
Harding I, Kumar S, “Thermal Stability of Precipitated Austenite in Fe-10Ni-0.1C
Steel”. TMS 147th
Annual Meeting and Exhibition, March 2018.
Harding I, Kumar S, “Microstructural Evolution in Fe-10Ni-0.1C Steel during Dynamic
Deformation” (Poster). TMS 147th
Annual Meeting and Exhibition, March 2018.
Teaching Experience
Introduction to Materials Science, Lab TA (Falls 2015, ’16, ’17)
iv
Acknowledgements
I would like to begin by thanking my family, who have always fostered my curiosity and
firmly believed in the value of education. They instilled in me a love of learning, and I am
grateful for the sacrifices they made for me to have this opportunity.
To Jenn- thank you for your patience, companionship, and endless support. It takes
significant dedication to get a PhD, but comparable devotion to help someone get through a PhD.
Thank you for your sacrifices for my dream, for helping me stay focused, and for looking out for
me when I forget to look after myself.
I would like to acknowledge the help from my labmates, who have taught me the
experimental methods I have used, who have helped me think through problems, who have
shared their extensive knowledge of the literature, and who have helped guide my research. In
particular, thank you to postdocs Hyokyung Sung, Hyunmin Kim, and Hyung Soo Lee. I am
grateful to the many technicians, without whom all research in the university would quickly
grind to a halt. Specifically, I would like to thank Tony McCormick for his help in all things
electron microscopy and John Shilko, Brian Corkum, and Chris Bull for their assistance in Prince
Lab. Additionally, I’d like to thank the Max-Planck-Institut für Eisenforschung for hosting me
and allowing me access to their facilities. I am grateful to the help of Isabelle Mouton, Bat Gault,
and Prof. Dierk Raabe.
I’d like to acknowledge the many friends I have made at Brown. Mariami Bekauri and
Leah Nation for getting through Materials courses together; the ‘Brown Engineering Mafia’ of
Stelios Siontas, Steve Racca, Peter Sun, Gerardo Pradillo, and others; the mentorship of Steve
Ahn, Jay Sheth, and Max Monn; and my good friends Jon Estrada, Hadley Witt, Alex Landauer,
v
Rana Ozdeslik, Mrityunjay Kothari, and Mohak Patel. Research can be arduous and isolating,
and I cherish the time we spent together.
Lastly, I would like to thank the many academic mentors that I have had. I owe my
scientific and mathematical foundation to Mark Hammond and Eric Kemer. They believed the
foundation of science lies in rigorous understanding of fundamental principles, which in turn I
have relied heavily on throughout my higher education. I am forever grateful to my
undergraduate mentors, Dr. Ruth Ost and Prof. Dieter Forster, for believing in me and working
so hard on my behalf. Thank you both for your fierce advocacy and for your help in finding my
way. I’d like to thank Prof. Ke Chen for patiently teaching me the fundamentals of experimental
research; without the exposure I had in his lab, I undoubtedly would not have pursued this career.
Lastly, thank you to my advisor, Prof. Sharvan Kumar. I did not have a background in
engineering, but you had the patience for me to learn and change the trajectory of my career. I
appreciate your infectious dedication to learning and love of science, but more so your ability to
understand the big picture- life, family, happiness, and fulfillment.
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Table of Contents
Curriculum Vitae ............................................................................................................................ 4
Education ..................................................................................................................................... 4
Conference Presentations ............................................................................................................ 4
Teaching Experience ................................................................................................................... 4
Acknowledgements ......................................................................................................................... 4
Chapter 1: Introduction ................................................................................................................... 1
Chapter 2: Technical Background .................................................................................................. 4
2.1: Outline .................................................................................................................................. 4
2.2: Fe-Ni Alloys and Fe-Ni low C Steels .................................................................................. 5
2.3: Quench and Partition, TRIP, and other Advanced High Strength Steels ........................... 14
2.3.1: Autotempering ............................................................................................................. 14
2.3.2: Thermal and Mechanical Stability of Austenite .......................................................... 16
2.4: Fe-10Ni-0.1C steels and the QLT Heat Treatment ............................................................ 21
2.5: Dynamic Deformation of Steels and Other Alloys ............................................................ 26
2.6: Scope of this Effort ............................................................................................................ 36
Chapter 3: Experimental Procedure .............................................................................................. 37
3.1: Materials and Heat Treatment Schedules ........................................................................... 37
3.2: High Strain Rate Deformation ........................................................................................... 39
3.3: Microstructure Characterization ......................................................................................... 41
3.3.1: Optical Microscopy: .................................................................................................... 41
3.3.2: Scanning Electron Microscopy (SEM):....................................................................... 42
3.3.3: Transmission Electron Microscopy (TEM): ................................................................ 43
3.3.4: Atom Probe Tomography (APT) ................................................................................. 50
3.3.5: Nano-Indentation ......................................................................................................... 52
Chapter 4: The Partitioning of Carbon During the Heat Treatment of Quenched Fe-10Ni-0.1C
Steel............................................................................................................................................... 53
4.1: Introduction ........................................................................................................................ 53
4.2: Results: L Temper .............................................................................................................. 54
4.3: Results: T’ Temper and QLT ............................................................................................. 61
4.4: Discussion .......................................................................................................................... 62
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Chapter 5: The Partitioning of Ni During the Heat Treatment of Quenched Fe-10Ni-0.1C Steel 64
5.1: Introduction ........................................................................................................................ 64
5.2: The As-quenched Microstructure (AQ) ............................................................................. 64
5.3: The L Tempers Microstructure (QL and Q25L) ................................................................ 65
5.4: Discussion of Isothermal Tempering ................................................................................. 78
5.5: The QLT Treatment – Results and Discussion .................................................................. 82
Chapter 6: Microstructural Evolution in an Fe-10Ni-0.1C Steel During Dynamic Deformation 88
6.1: Introduction ........................................................................................................................ 88
6.2: Kolsky-Bar Calibration using 4140 Steel .......................................................................... 88
6.3: Kolsky-Bar Testing of 10Ni-QLT ...................................................................................... 89
6.4: Nanoindentation across the ASB ....................................................................................... 91
6.5: Microstructural Analysis of the dynamically deformed Specimen .................................... 93
6.5.1: Location far from the shear band: ............................................................................... 96
6.5.2: Location ahead of shear band: ..................................................................................... 97
6.5.3: Location adjacent to the shear band: ......................................................................... 102
6.5.4: Location inside shear band: ....................................................................................... 103
6.6: Discussion ........................................................................................................................ 112
Chapter 7: Conclusions ............................................................................................................... 115
7.1: Microstructural Evolution during Heat Treatment of an Fe-10Ni-0.1C Steel ................. 115
7.2: Microstructural Evolution during High Strain-Rate Deformation of an Fe-10Ni-0.1C Steel
Subjected to a Two-Stage Heat Treatment .............................................................................. 116
Chapter 8: Recommendations for Future Work .......................................................................... 118
References ................................................................................................................................... 121
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List of Tables
Table 3.1: Nominal alloy composition .......................................................................................... 37
Table 3.2: Example table of calibrated diffraction radii ............................................................... 44
Table 4.1: C content in austenite and ferrite in Q25T’, Q125T’, and QLT as measured by APT
(at.%) ............................................................................................................................................. 61
Table 5.1: Composition of austenite in the four isothermal heat treatments as measured by APT
(at.%) ............................................................................................................................................. 77
Table 5.2: Size and composition of precipitates in isothermal treatments ................................... 78
ix
List of Figures
Figure 2.1: Fe-rich end of the binary Fe-Ni equilibrium phase diagram [30]. Note that as
temperature decreases, the equilibrium Ni content in austenite increases rapidly. ........................ 5
Figure 2.2: Schematic representation of the QLT treatment (adapted from [12]). Three single-
stage tempered samples (QT2, QT100, and QL) were also produced to compare to the full QLT
treatment. ...................................................................................................................................... 12
Figure 2.3: a) The QLT heat treatment- the steel is first austenitized at 800°C for 60 minutes and
water quenched; tempered at 650°C for 40 minutes and water quenched; and then tempered at
590°C and quenched. b) The Fe-Ni binary equilibrium phase diagram with each QLT treatment
temperature marked [30]. c) The Fe-Ni Ms temperature diagram with the equilibrium Ni content
in austenite at each stage of QLT marked [43]. ............................................................................ 22
Figure 2.4: Overview of the QLT treatment. At the austenitizing temperature, the steel is fully
austenitic (with some undissolved fine carbides). When quenched, it forms a lath martensite
structure. The L-temper produces a high volume of precipitated austenite, though the thermal
stability of the austenite is not fully understood. The T-temper produces additional austenite
precipitation that is more Ni-rich than L-generation, and the final result is the presence of an
appreciable fraction of thermally stable austenite. ....................................................................... 23
Figure 2.5: Ballistic field test results that led to the combination of a 10Ni alloy and the QLT
treatment, as reported by Zhang [25]. ........................................................................................... 24
Figure 2.6: DICTRA model results compared to APT results, taken from [28], confirm austenite
precipitation in Ni-rich fresh martensite pockets that are formed by the transformation of some of
the L-generation austenite to martensite during the quench from the L temperature. .................. 26
Figure 3.1: Heat treatment schedules considered. All treatments begin with austenitization and
quench (AQ). Two single-stage tempering temperatures (650°C-L and 540°C-T’) with various
lengths are analyzed, as well as the two-stage temper QLT ......................................................... 38
Figure 3.2: Kolsky-bar setup. The sample is sandwiched between two bars; a projectile is fired at
the end of one bar, and the strain pulses through the incident bar, reflected back through the
x
incident bar, and transmitted through the transmission bar are recorded with strain gauges. A
stress/strain curve is thus derived from these pulses. ................................................................... 39
Figure 3.3: a) an example of EDS analysis of QLT using a manually condensed beam (rough size
marked by pink circles) in the CM20. Austenite precipitates include roughly one datum point
each. b) STEM EDS on the JEOL 2100F allows for a high enough data point density to create
linear gradient plots within grains................................................................................................. 46
Figure 3.4: STEM mode images are mapped to BF using manually selected correlative point
pairs- i.e. a location is selected on the STEM image, and then the same point is selected in BF.
Several of these are used to create an affine map. Data is collected in small sets to avoid
significant drift during collection; these sets are uploaded in STEM space, mapped to BF mode,
and color coded by composition. Dozens of individual sets create dense composition maps. ..... 49
Figure 3.5: Sample EDS Ni gradient on a Q336T’ austenite precipitate ...................................... 50
Figure 3.6: A pictorial summary of the experimental process utilized in this effort to characterize
the undeformed microstructure as a function of heat treatment. .................................................. 52
Figure 4.1: a) TEM brightfield image of an austenitized and quenched (AQ) specimen showing a
lath martensite microstructure; some fine carbides (~20nm) are highlighted with dashed circles.
b) APT composition measurement of a carbide shows predominantly Mo but includes significant
amounts of V and Ti. c) C segregation along interfaces in martensite. ........................................ 56
Figure 4.2: a) TEM EDS/MDP shows Ni-depleted ferrite of about 8 wt.% Ni and Ni-rich (15-17
wt.%), thermally stable austenite precipitates. b) APT composition measurement of a carbide
particle in the QL specimen. A Ti- and V-rich core is surrounded by a Mo-rich but Ti- and V-
depleted outer shell (delineated by dashed lines). ........................................................................ 58
Figure 4.3: a) TEM EDS/MDP show the austenite precipitates are moderately Ni-rich (15 wt.%)
while the ferrite is close to equilibrium value of 5 wt.%. The precipitates are thermally unstable.
b) Interface segregation of C in fresh martensite in the Q25L specimen. c) APT measurement of
the composition across a carbide particle shows a V-rich, Mo-poor carbide is sandwiched by Mo-
rich, V-depleted carbide. ............................................................................................................... 60
Figure 5.1: TEM brightfield image with SADP of the AQ treatment. Note the lath martensite
structure......................................................................................................................................... 65
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Figure 5.2: a) SEM images of QL and Q25L. After only 40 minutes of tempering, austenite has
precipitated out.; after 25 hours, the precipitates have evolved and occupy more than 50% of the
total area. b) TEM-EDS arrays start around bulk composition (10 wt.%) for short tempering, but
a bi-modal distribution is visible after 25 hours. .......................................................................... 66
Figure 5.3: a) TEM BF/MDP with corresponding EDS measurements (b) of the QL treatment.
Note that the center of the precipitate is Ni-rich, around 17 wt.% Ni, and radially decreases in
composition. This is an artifact of overlap between the austenite and ferrite, which causes a
dilution effect where the austenite is thinnest. This is confirmed from interfacial composition
measurements using APT (c), where the composition measured near the phase boundary matches
the composition measured by EDS at the center of the austenite particles. ................................. 68
Figure 5.4: a) TEM BF/MDP with corresponding EDS (b) of Q25L shows Ni content varies
throughout, with some pockets of ~17 wt.% Ni and large swaths of 12-15 wt.% Ni. Some of the
moderate Ni values are believed to be true, as they are from the thickest parts of the austenite,
and thus are unlikely to be an average of a higher Ni content averaged with low-Ni ferrite. b)
APT confirms that some of the low-Ni readings are not an artifact from overlapping phases, as
the boundaries can be as low as 12 wt.% Ni. Together, we can conclude that some regions within
Q25L austenite are quite Ni-rich, while others are very Ni-lean. MDP shows the austenite of
Q25L is thermally unstable. .......................................................................................................... 71
Figure 5.5: a) SEM comparison of T’ treatments. Ni diffusion at the T’ temperature is much
slower than at L, and so even after 5 hours there is not a significant volume of precipitates. After
125 hours, precipitates now line most lath/packet/block/grain boundaries, and after 336 hours,
coarsened globules are seen along high-angle boundaries. b) EDS measurements of the Ni
content in ferrite show a progressive shifts towards its equilibrium value with longer tempering
times .............................................................................................................................................. 73
Figure 5.6: a) TEM BF with corresponding EDS (b) of Q25T’ shows the core of precipitates
have a composition between 24-26 wt.% Ni with radially decreasing values. As in QL, APT (c)
confirms that this decrease is an artifact of phases overlapping through the foil thickness – APT
composition measurements near austenite/ferrite interface closely matches those measured with
EDS at the austenite core. As precipitates grow with increasing tempering time (Q125T’, d), the
overlap effect in EDS diminishes while the core value remains constant (e). APT again
xii
corroborates EDS measurements(f). All precipitates in Q5T’-125T’ were found to be thermally
stable (FCC). ................................................................................................................................. 75
Figure 5.7: Q336T’ has large, coarsened, equiaxed austenite precipitates on the scale of 1-2µm.
They have a leaner composition than smaller T’ treatment precipitates (20-22 wt.% Ni vs 24-26
wt.%), and are sometimes thermally stable (a, c) or have transformed to martensite during the
quench (b, d). Note the thermally stable austenite (c) has a low defect structure, while the
thermally unstable austenite (d) has a lath-like internal structure. ............................................... 76
Figure 5.8: The Ms temperature diagram for a binary Fe-Ni alloy [43] predicts both L and T’
austenite to be thermally unstable at room temperature. The observed austenite stability in short
tempering times is thought to be due to a size-effect.................................................................... 80
Figure 5.9: For both L and T’ tempers, the small austenite from short tempers is thermally stable.
However, even the very Ni-rich T’ austenite becomes unstable after 336 hours, suggesting both
small L and T’ tempers are only stable due to their small size. If let coarsen long enough,
presumably all T’ austenite should be unstable. ........................................................................... 80
Figure 5.10: SEM of QL (shown here again for convenience) and QLT. Note that there is
significantly more austenite than QL, which suggests that the T temper contributes heavily to
additional austenite nucleation and growth................................................................................... 82
Figure 5.11: a) TEM BF of QLT with corresponding EDS (b) shows mixed composition of 15-20
wt.% at the center of precipitates. This is not an artifact of phase overlap, but rather is a result of
two distinct generations of austenite with different growth compositions. c) APT tips confirm
that both L-generation and T-generation austenite are present in QLT. ....................................... 83
Figure 5.12: QLT process produces thermally stable austenite around 15-17 wt.% Ni during the
L treatment. During T, growth resumes, resulting in additional growth in the range of 20-22
wt.% Ni. ........................................................................................................................................ 85
Figure 5.13: QLT Process, understood in terms of austenite size and composition. .................... 85
Figure 6.1: Strain/strain-rate pairs for Kolsky experiments on 4140 tempered martensite. Each
datum point correlates to a specific firing pressure with the noted bar. ....................................... 89
Figure 6.2: A diffuse and a well-developed adiabatic shear band ................................................ 90
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Figure 6.3: Sample chosen for microstructural analysis. A well-developed shear band initiates at
the bottom left, propagates towards the center of the sample (top right) and eventually dissipates.
Another band initiated in the opposite corner (not visible in this image) that also propagated
towards the center of the sample, though it was smaller and less developed than the one seen in
the montage above. The total height strain was about 20% and the strain-rate was about 1900s-1
.
Specific TEM specimen lift-out locations are marked using numbers from 1 to 10. The relative
displacements of the vertical etching bands (banding in the rolled plate) provide a sense of the
large shear experienced within the ASB. ...................................................................................... 91
Figure 6.4: Nanoindentation across the shear band (marked in purple). There is an increase in
hardness (~5GPa to ~7.5GPa) that extends beyond the boundary of the unetched portion of the
shear band. .................................................................................................................................... 92
Figure 6.5: SEM micrographs of QLT at increasing magnification (top to bottom), undeformed
(left) and near the shear band (right). Note the lack of features in the band, which gives it the
unetched appearance. Outside of this region, the austenite precipitates can be seen to be sheared
parallel to the shear plane. ............................................................................................................ 94
Figure 6.6: SADP in undeformed QLT compared to in the band using the same aperture. On the
left, the lath martensite structure produces a highly textured SADP; on the right, the highly
misoriented, equiaxed, mechanically recrystallized grains in the band produce a ring-like SADP.
....................................................................................................................................................... 95
Figure 6.7: TEM lift-out from a location that is far from the shear band (circled in blue dashed
line). .............................................................................................................................................. 96
Figure 6.8: Far from the shear localized region, the lath martensite seen in the undeformed
microstructure is still intact. Here, the composition of two austenite precipitates are measured
and found to be between 15-20 wt.% Ni. Thus these precipitates are a mixture of L- and T-
generation austenite, and MDP shows that they are mechanically stable. .................................... 97
Figure 6.9: a) TEM specimen lift-out from a location ahead of the shear band. Notice the stark
white band on the left, which dissipates to the right. b) SADP from the location ahead of the
band microstructure (aperture marked in orange). Ring pattern suggests significant grain
xiv
refinement and grain rotation and recognized from the relatively low magnification bright field
image. ............................................................................................................................................ 98
Figure 6.10: a) Two adjacent subgrains, which are brought into contrast by tilting. b) EDS on the
first subgrain gives a composition of about 22 wt.% Ni (T-generation), and MDP shows it is
mechanically stable FCC. c) Its neighbor is also from T-generation austenite (20 wt.% Ni), and is
also mechanically stable austenite. Together, they show a T-generation austenite precipitate has
mechanically broken down into subgrains and these are mechanically stable. .......................... 100
Figure 6.11: An L- and T-generation austenite precipitate has mechanically broken down into
two subgrains (b and c) and can be brought into contrast through tilting (a). The subgrain before
tilting (b) is mechanically stable and has Ni content varying between L- and T-generation, and
was likely part of a core/shell structure before deformation. Its neighboring subgrain (c) is also
mechanically stable, but its composition shows that it is entirely growth from the T temper. ... 101
Figure 6.12: a) Lift-out location with ASB demarcated by the two orange dashed lines and axes
defined (shear plane normal is X axis, and shear band is propagating in the Z direction). b) TEM
BF image of area about 20µm from the band with rotated axes marked. Microstructure is
elongated in the shear plane (Y-Z).............................................................................................. 102
Figure 6.13: Both, the ferrite and the austenite precipitates adjacent to the band are aligned
parallel to the shear plane (axes marked) and have broken up into subgrains. The subgrain that is
in contrast here has Ni content ranging from 15-22 wt.%, suggesting it originates from a mixture
of L- and T-generation austenite. MDP shows the subgrain to have the FCC structure. ........... 103
Figure 6.14: Exaggerated schematic of TEM specimen lift-out geometries. The austenite
precipitates are stretched parallel to the shear band- when lift-outs are taken along this plane,
there is an averaging effect in the EDS measurements. Lift-outs are taken normal to the band to
mitigate this effect....................................................................................................................... 104
Figure 6.15: Left: Location of in-band lift-out prior to FIB milling. The unetched portion of the
band is approximately 4µm from the lift-out edge, while the other boundary of the band is about
10µm into the lift-out. Right: The measurements taken on the SEM are overlaid on a BF image
of the lift-out specimen (which extends beyond what is shown, as marked by the black arrows).
..................................................................................................................................................... 104
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Figure 6.16: Normalized EDS Ni distribution of undeformed QLT (Green) in comparison to data
obtained from a TEM specimen that was lifted out parallel to the band (Blue) and another that
was oriented normal to the shear band (Red). The parallel geometry has a strong, narrow peak
around the nominal Ni content due to an averaging effect rising from the ratio of grain depth to
sample thickness; the normal-to-band sample has a distribution that closely matches the
undeformed material, suggesting it largely has single-grain depth. ........................................... 105
Figure 6.17: BF and corresponding DF image of the region of the deformed specimen that is
estimated to include the white band and the region immediately outside it. There is no observable
difference in microstructure, but note that the demarcation of the boundary is a best estimate
based on markings in Figure 6.15. .............................................................................................. 105
Figure 6.18: A small grain is in contrast. It is low in Ni (boundaries marked in orange on EDS
gradient plot, 5-7 wt.%Ni). Inside, a band-like substructure suggests twinning. ....................... 107
Figure 6.19: The high-Ni region, formerly an austenite precipitate, is spread over a very narrow
strip (about 20nm wide by 200nm in length). The portion that is in contrast is marked in the plot
in orange and has a Ni content of 20 wt.%, suggesting it is T-generation austenite; an MDP from
this location confirms it to be FCC and additionally reveals twinning- it is mechanically stable
austenite. ..................................................................................................................................... 108
Figure 6.20: The region presented in this image falls along the border of the white band. On the
left, the narrow strand of high-Ni (running left to right) is shown to be over 100nm long with
only a small grain of about 20nm in diameter in contrast. On the right, the composition of the
grain is isolated and highlighted in orange: it is about 20 wt.% Ni, and therefore came from T-
generation austenite. The MDP is BCC [111]- this grain is very fine, high in Ni but was
mechanically unstable. ................................................................................................................ 110
Figure 6.21: Another region close to the border of the white band; here, a small grain that
dynamically recrystallized from L-generation austenite (15-17 wt.%, marked in orange) has
mechanically transformed to martensite. .................................................................................... 111
1
Chapter 1: Introduction
There is considerable interest in the U.S. Navy to improve the performance of high strength
low alloy (HSLA) steels used in structural naval applications, due to their cost, mechanical
properties, and fabricability [1]. Ballistic resistance of steels used in hulls and ship decks is a key
property that needs to be improved to meet the requirements of the next generation military
vessels.
Ni-alloyed, low-C steels are promising candidates for the next generation of naval ships
due to their high static strength and toughness, especially at low temperatures [2–18]. This effort
began within the US Navy with off-the-shelf Ni-steels with promising static properties [2–18],
that were then incrementally adjusted in composition and heat treatment through a series of
studies including microstructural analysis, quasi-static mechanical testing, and ballistic field-
testing [20–29]. At the time we began the current study, one particular alloy composition and a
specific heat treatment schedule were isolated as having superior ballistic resistance while
maintaining good static strength and toughness. Specifically, an Fe-10Ni-1.0Mo-0.6Mn-0.6Cr-
0.08V-0.1C (wt.%) steel heat-treated according to a two-stage temper called the QLT treatment
(hereinafter referred to as 10Ni QLT) was deemed optimized [25]. The QLT treatment produces
a dispersion of fine, Ni-rich, thermally stable austenite precipitates in a ferritic matrix [25,28]; it
is thus thought that the high quasi-static toughness and resistance to shear localization during
ballistic impact are due to the mechanical instability of this austenite. Specifically, by
mechanically transforming to martensite, the austenite increases the work hardening capacity of
the steel, which delays the onset of fracture in quasi-static loading and perhaps in high-strain rate
deformation as well [13,25]. However, further improvement in ballistic response is limited by the
2
lack of a thorough understanding of microstructure/heat-treatment relationships and
microstructural evolution during dynamic deformation. These are the goals of this effort.
The thermal and mechanical stability of austenite are both largely governed by the
austenite precipitate size and composition. Therefore, in order to further improve 10Ni QLT, it is
critical to understand the relationship between heat treatment, austenite size and composition,
and its thermal stability. While some work has been done to partially address these questions, it
is not adequate or expansive enough to suggest improvements to 10Ni QLT [20,25,27,28].
Next, how do the various phases in the microstructure (i.e. austenite, ferrite, and carbides)
evolve during high strain rate deformation? Specifically, how does austenite dynamically evolve
as a function of its size, composition, and proximity to shear localization? It is broadly
understood that mechanical transformation does occur during ballistic field tests, but it is unclear
when during the localization phenomenon the transformation occurs, and if it is the main reason
that 10Ni QLT has superior ballistic properties over other treatments and alloys or if the stress-
induced austenite to martensite transformation is simply a side effect of high strain rate
deformation [25]. Understanding the evolution of QLT microstructure during dynamic
deformation will allow us to identify microstructural aspects that could be further tweaked and
could conceivably result in further enhanced ballistic response.
This effort fills some of these gaps in understanding. This thesis is organized in the
following manner. Chapter 2 (Technical Background) discusses the literature relevant to this
work. The areas covered in this review include studies pertaining to: heat treatment and
mechanical properties of Ni-containing alloys and steels for cryogenic applications; mechanical
properties, design, and characterization of advanced high strength steels including
Transformation-Induced Plasticity (TRIP) steels and Quench and Partition (Q and P) steels;
3
microstructure characterization of 10Ni QLT and its predecessors; microstructural evolution of
steels and non-ferrous alloys during high strain rate deformation; and characterization of the
microstructural evolution of 10Ni QLT and its predecessors after high strain rate deformation.
The experimental methods used in this study are explained in Chapter 3. Chapter 4
discusses the partitioning of C during tempering in this alloy and its role on the stability of
precipitated austenite and precipitation and growth of carbides. In particular, the partitioning of
C during isothermal tempers of 40 minutes and 25 hours at 650°C is compared to the as-
quenched material to determine the partitioning behavior as a function of tempering time;
isothermal treatments at 540°C and the QLT treatment are also analyzed to compare the role of
tempering temperature on C partitioning and carbide formation. Chapter 5 details the results of
the analysis of Ni partitioning as a function of tempering time and temperature. Multiple
characterization techniques are used to determine the relationship between tempering time and
temperature on precipitated austenite chemical composition and size, which are then related to
austenite thermal stability. Analysis of isothermal heat treatments additionally provides possible
microstructure-based explanations for the superior mechanical properties of the two-stage QLT
temper. Chapter 6 includes the results and discussion of the microstructural evolution in 10Ni
QLT during high strain rate deformation. Microstructure from various locations relative to an
adiabatic shear band (ASB) is characterized to understand the microstructural evolution during
shear localization. The last chapters are the Conclusions chapter, which summarizes all results,
and a Chapter discussing future considerations.
4
Chapter 2: Technical Background
2.1: Outline
In this Chapter, the literature is reviewed in the following four areas that are relevant to
the research reported in this thesis.
1) Fe-Ni Alloys and Ni-containing, Low-C Steels: The current state of knowledge
regarding austenite precipitation and mechanical properties of Fe-Ni alloys and Ni-
containing steels with relation to heat treatment.
2) Transformation Induced Plasticity (TRIP), Quench and Partition (Q+P), and other
Advanced High Strength Steels (AHSS): Although TRIP/Q+P steels are often
primarily Mn- or C-alloyed to produce thermally stable austenite, the research on
novel heat treatment schedules, austenite precipitation kinetics, and austenite thermal
and mechanical stability may be extended to 10Ni QLT, which is the subject of this
study.
3) Fe-(2.5-10)Ni-0.1C (wt.%) Alloy Development: Summary of research done on Ni-
containing steels for ballistic resistance, associated heat treatment optimization
studies leading to the 10Ni QLT material, and current understanding of the
microstructure evolution in this material.
4) High Strain Rate Deformation and Adiabatic Shear Banding (ASB): Current
understanding of high strain rate deformation in steels, including the 10Ni steel, as
well as in some non-ferrous alloys. The review attempts to connect prior research in
dynamic deformation mechanisms and microstructural evolution in other materials to
post-deformed microstructure observations in ballistic field-tested 10Ni QLT.
5
2.2: Fe-Ni Alloys and Fe-Ni low C Steels
Nickel-containing steels were developed in the 1940s for use in cryogenic applications
due to the strong suppression of the ductile-to-brittle transition temperature (DBTT) of the steel
[2]. The microstructural basis for this increased toughness at low temperature, especially in
relation to heat treatment of the steel, was extensively studied in the 1970’s and 80’s [2–18].
While these studies did not have the spatial or compositional resolution of modern techniques
(e.g. APT), they nevertheless provide a foundation for our understanding of deformation
processes and a framework for our research. The studies considered in this section represent a
steady, incremental progression in understanding of the Fe-Ni system, and key results are
described.
Figure 2.1: Fe-rich end of the binary Fe-Ni equilibrium phase diagram [30]. Note that as temperature decreases, the equilibrium
Ni content in austenite increases rapidly.
The binary Fe-Ni phase diagram is shown in Figure 2.1. Ni is a strong austenite stabilizer
and upon cooling a binary Fe-rich Fe-Ni alloy from the single-phase austenite region, it
decomposes to yield relatively Ni-poor ferrite and Ni-rich austenite. Conversely, if a low Ni
alloy (for example Fe-(6-10) wt.% Ni) is suitably quenched from the austenite region to produce
6
a lath martensite structure, upon tempering in the two-phase region, the martensite will
decompose into Ni-poor ferrite and Ni rich austenite; that is, Ni-rich austenite precipitates out
during tempering. If tempering time is adequate, the volume fraction and the composition of the
austenite can be simply determined from the phase diagram.
However, Ni is a substitutional element in Fe, and so at lower tempering temperatures
(i.e. below 550°C), Ni diffusion can be quite sluggish. For this reason, equilibrium
austenite/ferrite phase fractions may not be reached in practical tempering times, as noted by
Romig and Goldstein [31]. Nevertheless, Ni-alloyed steel (e.g. 5-15 wt.% Ni) can be
austenitized, quenched, and tempered (e.g. between 550-650°C) to produce Ni-rich austenite
[2,3,6,12,13], and if rich enough in Ni, the austenite may be thermally stable even after a
subsequent quench from the tempering temperature [32]. The presence of thermally stable
austenite in the steel has been shown to result in significant improvements in DBTT suppression
and low-temperature fracture toughness [2,4–18]. Understanding the relationship between heat
treatment, the low-temperature mechanical properties of the material, and the thermal and
mechanical stability of austenite were the focus of these studies.
To understand the relevant mechanisms that govern the relationship between heat
treatment and mechanical properties in Ni-alloyed steel, it is first important to describe the
morphology of precipitated austenite in Fe-Ni alloys and how it can be affected by tempering
time and temperature. The Ni-rich austenite precipitates along lath boundaries [2,3,6,13] and
obeys the Kurdjumov-Sachs relationship, as observed by Kim et al. [13] and Fultz et al. [2] in
tempered Fe-6Ni-0.1C (wt.%) steel and tempered Fe-9Ni-0.1C (wt.%) steel respectively.
Furthermore, Kim et al. found that the austenite in Fe-6Ni-0.1C (wt.%) steel that was thermally
unstable reverted to the orientation of the surrounding laths, while mechanically unstable
7
austenite transformed to a different orientation [13]. The morphology of precipitated austenite
also depends on tempering temperature. Kim et al. found that when an Fe-6Ni-0.1C (wt.%) steel
was tempered at 670°C, the precipitated austenite had a lenticular morphology along interlath
interfaces, while the austenite formed at 600°C was blocky and equiaxed [13].
Low-C steels with nominal Ni content under 10 wt.% have been reported to peak in the
volume fraction of thermally stable precipitated austenite for a given tempering temperature.
That is to say, with sufficiently long tempering times, some of the precipitated austenite becomes
thermally unstable in this type of steel. This was observed by Marschall et al. in Fe-9Ni-0.1C
(wt.%) tempered between 590-650°C [33], by Hwang et al. in Fe-12Ni-0.25Ti (wt.%) alloy
tempered between 575-600°C [6], and by Fultz et al. in Fe-9Ni-0.1C (wt.%) steel tempered at
590°C [2]. The onset of thermal instability for extended isothermal tempering was not fully
understood, but Fultz et al. suggested that both chemical and microstructural factors contribute to
the stability of austenite [2,17].
The partitioning of Ni, C, and other minor alloying elements during tempering
contributes to austenite stability [17], and was directly measured using STEM EDS, Mössbauer
spectroscopy, and wet chemical analysis on chemically extracted austenite particles [2,12,17].
The most reliable composition measurements were made by Fultz et al. on an Fe-9Ni-0.1C steel
tempered at 590°C for a range of tempering times [2,16,17]. STEM EDS and Mössbauer
spectroscopy were both used to measure the composition of the precipitated austenite, but the C
content could not be measured using these techniques. Instead, the C content in the austenite was
calculated as a function of austenite volume fraction and ferrite lattice parameter as measured by
Mössbauer spectroscopy. The precipitated austenite was found to be enriched in Ni, Mn, Cr, Si,
and C, as predicted in [10,11]. Additionally, the composition of the austenite was found to be
8
richest in Ni on the outer edges of the austenite particles, which was attributed to its slower
diffusion rate relative to other elements. The C content in the austenite held at 0.7 wt.% for the
first 10 hours of tempering and then decreased, an effect that was attributed to the ferrite
becoming depleted of C after 10 hours, thus diluting the C content in the austenite as it continued
to grow. While the error in measuring the minor alloying elements using these techniques was
relatively high, the measured solute enrichment was estimated to strongly lower the martensite
start (Ms) temperature of the precipitated austenite. However, the estimated Ms of the different
austenite particles was not enough to fully explain their thermal stability. Thus, they concluded
another factor must also contribute to the thermal stability of the austenite.
Hayzelden and Cantor related austenite grain size to thermal stability of the austenite
phase by measurement of the Ms temperature for an Fe-26Ni-0.4C (wt.%) steel [18]. They
showed that for austenite grains around 150µm there was no change in Ms temperature as grain
size decreased until grain size approached 10µm, where there was a suppression of Ms by
approximately 50°C. They did not measure grain sizes smaller than 10µm, but the trends
suggested that this suppression would increase with further grain refinement. Thus, combined
with the results of Fultz et al. [16,16,17], the thermal stability of precipitated austenite is strongly
affected by composition, but other microstructural factors including precipitate size also play a
significant role in controlling the thermal stability of austenite and therefore the deformation
mechanics and mechanical properties of the steel.
In addition to the thermal stability relationship of Hayzelden and Cantor [18], Fultz et al.
demonstrated that the mechanical stability of austenite was also a function of grain size in an Fe-
9Ni-0.1C steel [2,16,17]. The steel was tempered for various times at 590°C to produce a range
in austenite grain sizes, and then the sheets were cold-rolled to induce mechanical
9
transformation. The defect structure surrounding transformed particles was used to qualitatively
estimate the transformation strain energy. Based on their observations, they concluded that larger
particles needed less strain energy to transform, and thus were more mechanically unstable.
Several microstructural factors have been observed to have beneficial effects on low-
temperature mechanical properties such as Charpy fracture energy, fracture toughness, and
DBTT: a reduction in grain size improves DBTT [3,4]; the precipitation of austenite, thermally
stable or not, improves Charpy upper shelf energy because of a gettering effect on interstitial
elements in the martensite [9,10,13]; increasing the volume of thermally stable austenite can
improve the low temperature Charpy fracture energy, fracture toughness, and suppress the DBTT
[5–7,11]; increasing mechanical stability of the austenite tends to lead to increases in fracture
toughness because it delays the onset of stress-induced austenite-to-martensite transformation
during loading [8,11,13]; and lastly the morphology of the austenite plays an important role in
the mechanical properties, with a fine distribution along lath boundaries reducing the effective
grain size by dissuading trans-packet cleavage [12–14]. Each of these findings is briefly further
examined in the following paragraphs.
Jin et al. found that the reduction in grain size in a ferritic Fe-12Ni-0.25Ti alloy from 40-
60µm down to 0.5-2µm through a cyclic annealing heat treatment lowered the DBTT to below
4.2K [3,4]. They argued that by decreasing the grain size, the critical stress for cleavage fracture
is increased above yield for lower temperatures, and thus DBTT is suppressed. Additionally, they
found that the precipitation of austenite through tempering (between 3-5 vol%, as measured by
XRD) further suppressed the DBTT beyond the effects of grain refinement [5]. The
microstructural mechanism behind the improved DBTT from thermally stable austenite was not
conclusively determined, but they suggested two possible hypotheses. First, there may be a
10
beneficial TRIP effect wherein the austenite mechanically transforms to martensite, also
enhancing the overall ductility. Second, the austenite could serve as a sink for deleterious
elements in the ferrite (‘gettering’), thus increasing its ductility and suppressing the DBTT.
Deformation-induced martensite transformation of precipitated austenite and its influence
on Charpy fracture energy has been examined by several researchers. Syn et al. found that the 8
vol% austenite in a tempered Fe-9Ni-0.1C steel had fully transformed to martensite as far as
1mm from the fracture surface in Charpy tests at 77K [8]. Additionally, Fultz and Morris found
that for Charpy tests on an Fe-9Ni-0.1C steel, the transformation depth was similar to the plastic
zone size and thus the austenite had already transformed before it could interact with the
propagating crack front [15]. Lastly, Kim and Schwartz also investigated TRIP effects during
Charpy tests in an Fe-9Ni-0.1C steel [9]. They noticed that an increased volume of thermally
stable austenite improved Charpy fracture energy and found evidence for a TRIP effect as deep
as 300µm from the fracture surface. However, they estimated that the increased toughness from
the additional austenite mechanically transforming via a TRIP effect was insufficient to explain
the observed increase in Charpy impact energy. Thus, while a TRIP effect can contribute to
improved Charpy fracture energy, additional factors enabled by austenite precipitation likely
play a role in improved mechanical properties.
Gettering was directly shown to have a positive effect on Charpy upper shelf energy
[9,10,13,14]. The clearest demonstration was by Kim et al. on an Fe-6Ni-0.1C (wt.%) steel that
was treated with a two-stage temper [14]. The composition of the thermally stable precipitated
austenite and ferrite in the heat treated Fe-6Ni-0.1C was each measured in one of their previous
studies [12], and so, two steels with nominal composition of each phase (austenite and ferrite)
were produced to compare to the two-phase material. The grain sizes of the austenitic steel and
11
the ferritic steel were reduced to be similar to that in the Fe-6Ni-0.1C heat treated steel, and then
the Charpy fracture energies at 77K and DBTT of all three were compared. The ferritic steel had
a similar upper shelf energy to the heat-treated Fe-6Ni-0.1C steel, and both the heat-treated Fe-
6Ni-0.1C steel and ferritic steel had higher upper shelf energy than the as-quenched Fe-6Ni-0.1C
steel. Thus, the removal of C and other elements from the martensite during tempering was
primarily responsible for the increased Charpy upper shelf energy seen in tempered Fe-6Ni-0.1C
steel. However, the single-phase steels had lower Charpy fracture energy than the tempered Fe-
6Ni-0.1C steel, suggesting the DBTT suppression was related to the two-phase morphology
rather than because of the properties of either phase individually.
Strife and Passoja [11] found that the mechanical stability of precipitated austenite played
a role in the low-temperature mechanical properties of low-C, Ni-containing steels. Two steels
were considered, an Fe-9Ni-0.1C steel with a single stage temper and an Fe-5Ni-0.1C steel with
a two-stage temper. While the 9Ni steel had improved fracture toughness with increased
austenite volume fraction, the two-stage tempered 5Ni steel was found to be more brittle with
increased austenite volume fraction. The difference in fracture behavior for austenite-containing
9Ni versus 5Ni suggested that improved mechanical stability in the precipitated austenite (which
can delay the stress-induced phase transformation and extend the strain hardening characteristic),
correlated to improved fracture properties, rather than simply austenite volume fraction.
Kim et al. proposed a morphology-based hypothesis to explain the effect of thermally
stable precipitated austenite on DBTT temperature that took into account the volume, thermal
stability, and mechanical stability of the precipitated austenite [12–14]. While precipitated
austenite does beneficially serve as a sink for deleterious elements like C, the presence of
thermally stable austenite that is resistant to mechanical transformation along lath boundaries is
12
additionally helpful because it decreases the effective grain size by the following mechanism. At
low temperatures, the martensitic material is susceptible to trans-packet cleavage because the
laths are similarly oriented, and so fracture can travel across the similarly-oriented cleavage
planes. By adding mechanically stable austenite along these lath boundaries, the cleavage planes
across laths are broken up, and the effective grain size is decreased. Therefore, increasing the
volume fraction of austenite that resists transformation should suppress the DBTT and improve
the fracture toughness.
Figure 2.2: Schematic representation of the QLT treatment (adapted from [12]). Three single-stage tempered samples (QT2,
QT100, and QL) were also produced to compare to the full QLT treatment.
13
To test their hypothesis, Kim et al. considered an Fe-6Ni-0.1C (wt.%) steel that was
austenitized at 800°C (Q) and then given various heat treatments (Figure 2.2) to produce
different austenite morphologies: one tempered at 670°C (QL), two tempered at 600°C for
different times (QT2 and QT100), and one given combined two-stage temper at 670°C and
600°C (QLT) [12–14]. The isothermal QT (600°C) tempers produced austenite far slower than
the L (670°C) temper and the austenite had a blocky, spherical morphology. The L temper
produced lenticular austenite, but it was thermally unstable due to a lower solute content. The
QLT sample was found to have a lenticular austenite, despite being tempered at T (which
produced blocky austenite in QT2 and QT100), and so it was determined that its lenticular
morphology was due to how it formed. The QL step resulted in lenticular pockets of Ni-rich
martensite, and the additional T step caused an austenite with even richer solute content to
precipitate in these pockets and inherit their morphology [13].
All four tempers had improved Charpy upper shelf energy over the as-quenched
microstructure (Q), which was attributed to the gettering of C from the martensite matrix. This
effect occurs even if the austenite is thermally unstable, as in QL, and thus QL had better Charpy
upper shelf energy than Q. However, the QLT treated sample had a higher Charpy fracture
energy and DBTT than the sample with the long T temper, despite having austenite with similar
composition and volume fraction. Thus, the difference between the DBTT of QT100 and QLT
was attributed to morphology: the lenticular austenite in QLT had better resistance to trans-
packet cleavage than the blocky-austenite in QT100. Additionally, Kim et al. noted that
thermally unstable austenite reverted to the orientation of neighboring laths, while mechanically
unstable austenite (produced during Charpy tests) transformed to a different orientation (as
mentioned previously) [13].
14
In summary, the volume fraction, thermal stability (composition and size dependence),
and morphology of austenite precipitated during tempering of low carbon-(5-10) wt.% Ni steels
is intimately tied to the low-temperature mechanical properties of the steels and these
microstructural aspects rely on tempering schedules (temperature, time, and multi-stage).
However, the role of minor alloying elements, most notably C, has not been adequately
identified and so the extent of their role on austenite thermal and mechanical stability or
mechanical properties of the steels has not yet been rigorously determined.
2.3: Quench and Partition, TRIP, and other Advanced High Strength Steels
While the work on cryogenic Ni-alloyed steels from the 1970’s and 80’s was limited by
the inability to accurately quantify light and minor alloying elements in precipitated austenite,
measure concentration gradients on the nanometer scale, or characterize nano-scale features,
these types of experiments have recently been enabled by advanced experimental capabilities and
performed in more contemporary work on steels with similar compositions and mechanical
properties. A few relevant studies have been isolated, and results from these are discussed in the
following section.
2.3.1: Autotempering
Through the use of high-resolution TEM and atom probe tomography (APT), it has been
observed that C can partition to small amounts of austenite along lath boundaries as martensite
forms during quenching, even in low-C steels [34–36]. This C-enrichment can be high enough
that these austenite films are thermally stable after the quench. This phenomena is called
‘autotempering,’ and depends on quenching rate [34].
Sherman et al. measured the C enrichment in these films as a function of quench rate in
Fe-1.3C-3.2Si-3.2 (Mn, Ni, Cr, Mo, Cu, Al, Ti, V) (at.%) using APT [34]. For samples quenched
15
at a rate of 55K/s, there was an enrichment with a peak of 10 at.% C and a width of 4nm in small
C-rich clusters, and austenite films of 12-18nm were observed using TEM. For samples
quenched at a rate of 560K/s, the films were too fine to identify in TEM, but APT showed 4nm
wide C-rich planes with a peak of 6 at.% C. 90% of the C that remained in the martensite in the
560K/s quenched steel was estimated to be trapped in Cottrell atmospheres at dislocation cores.
Morito et al. quantified C-enrichment from autotempering in an Fe-2Mn-0.2C (wt.%) steel
which was austenitized at 1200°C and water quenched to room temperature [35]. Despite the low
nominal C concentration, APT showed a local C enrichment in 3nm austenite films to the extent
of 4.5 at.% (~1 wt.%). The C content in the martensite matrix was measured to be 0.08 wt.%,
which was lower than the 0.2 wt.% initial C content in the alloy. The decrease in the C
concentration in the bulk martensite was attributed to short-range diffusion of the C to retained
austenite and dislocation structures in the martensite over six months at room temperature.
Morsdorf et al. noted that the extent of autotempering in individual martensite laths
depended on when during the quench the martensite transformed [36]. They analyzed an Fe-5Ni-
0.1C (wt.%) steel that was austenitized at 900°C and water quenched. Two types of martensite
were found: a coarse-lath martensite that formed early in the quench and fine-lath martensite that
formed closer to the martensite finish temperature. The coarse-lath martensite had less interstitial
C than the fine-lath martensite (~0.15 at.% vs 0.3-0.7 at.% respectively) because the early-
forming martensite had more time at high temperature than later-forming martensite, thus
allowing for additional C diffusion during the quench. Additionally, the C in the late-forming,
fine-lath martensite was found to segregate to Cottrell atmospheres. They estimated that even if
the alloy was quenched at 1000K/s, the C in the early-forming laths could diffuse as far as 1.5µm
16
(on the scale of the coarsest lath widths), while late-forming laths were confined to nanometer-
scale local diffusion.
2.3.2: Thermal and Mechanical Stability of Austenite
The idea of Quench and Partition (Q+P) steels was first proposed by Speer et al. [37] as a
novel heat treatment schedule for C-containing steels to produce a large, tunable volume of
thermally stable austenite. The process involves austenitizing and quenching the steel to a
quench temperature between the Ms and Mf temperatures, which retains a controlled volume of
austenite. The steel is then tempered at or above the quench temperature to allow the partitioning
of C from the martensite to the austenite, thus thermally stabilizing it during a final room-
temperature quench. Specifically, Speer et al. [37] noted that there was an optimum austenite
volume fraction given the total available C in the system; further, the volume of austenite during
the temper and the C content in the austenite after the temper are inversely related and depend on
the quench temperature. If the austenite after the temper is rich enough in C, it will be stable after
the final quench. Therefore, there is an optimum quench temperature in which the amount of C in
the retained austenite is just enough to thermally stabilize it at room temperature.
Takaki et al. studied the relationship between austenite grain size and the austenite to
martensite transformation in an Fe-16Cr-10Ni steel [38]. The composition was chosen such that
the Ms temperature was around room temperature. The steels were initially austenitic, cold-
drawn to mechanically fully transform the austenite to martensite, and then tempered at different
temperatures around 630°C (the austenite reversion temperature) for 10 minutes to create initial
austenite grain sizes between 0.8-80µm. The samples were held at room temperature and at 77K
each for up to 700 hours to allow the austenite to transform to martensite. The transformation
behavior could be classified by three grain sizes: coarse grain (~80µm), fine grain (~10µm), and
17
ultra-fine grain (~0.8µm). The coarse grain austenite had 20% transformation to martensite at
room temperature and 50% at 77K, the fine grain austenite had about 4% transformation to
martensite at room temperature and 35% at 77K, while the ultra-fine grain had almost no
transformation at either temperature. Using TEM, they showed that coarse-grain austenite
formed several packets each with multiple blocks of laths about 1µm wide (multi-variant
transformation); the fine-grain austenite formed martensite with a single-variant or single habit
plane; and what little ultra-fine grain austenite that did transform did not have a lath structure and
was single-variant.
The reason for the suppression of the austenite-to-martensite transformation in ultra-fine
grain austenite was rationalized through an estimation of the elastic strain energy barrier for
transformation in multi-variant transformation versus single-variant transformation. The elastic
strain energy for multi-variant transformation can be thought of as isotropic, as there are many
lath orientations present in the original austenite grain. In contrast, single-variant transformation
is highly anisotropic. The difference in these energies is substantial: approximately 1840MJ/m3
for single-variant versus 70MJ/m3 in multi-variant. Therefore, they concluded that a substantial
chemical driving force is necessary for sub-micron austenite to transform to martensite in
thermally metastable austenite.
Size also plays a role in the mechanical stability of austenite, as studied by Wang et al. in
manganese-containing TRIP steels [39,40]. In their first study, an Fe-9Mn-3Ni-1.4Al-0.01C
(wt.%) steel was austenitized, quenched, and then tempered at 600°C for 8 hours to precipitate
Mn- and Ni-rich austenite along lath boundaries. The precipitated austenite was thermally stable
at room temperature, and while the austenite precipitates varied in size, the precipitates had
similar composition. Furthermore, the former austenite grain size was large enough to form
18
multiple martensite packets, and so the austenite that precipitated within a particular packet had
similar crystalline and geometric orientation to each other with respect to the loading axis during
deformation experiments. Thus, the effect of size on deformation mechanisms was isolated from
effects of composition and orientation and their findings are described below.
The mechanical stability of the austenite was analyzed using EBSD in-situ as well as
after tensile and three-point bend tests, and the fine-scale microstructure was analyzed using
TEM [39]. The deformation behavior of the austenite was categorized by their size (surface area)
as small austenite precipitates (0.1-0.3µm2) and large austenite precipitates (0.3-4µm
2). At
relatively low global strains, the large precipitates showed little signs of plastic deformation,
while the small precipitates primarily deformed through stacking fault multiplication and
transformed to martensite. At higher global strains, the larger austenite precipitates were seen to
form subgrains separated by mechanical twins; the small subgrains then deformed through the
multiplication of stacking faults and eventually transformed to martensite. They concluded that
while smaller austenite is thought to be mechanically more stable than large grains, there is also
a size-dependent TWIP effect that can affect the TRIP mechanism. In TRIP steels with rather
homogenously sized and composed austenite, the austenite transforms at low strains and roughly
at the same global strains, which could actually promote crack propagation and failure [41],
suggesting delaying the transformation and spreading it out over a strain spectrum would be
beneficial. Specifically, in their experiment, the austenite-martensite transformation was spread
over a larger range of global strains because of the difference in deformation mechanisms for
small austenite precipitates versus large austenite precipitates. Thus, the mechanical properties of
the material were improved through a ‘spectral TRIP effect’.
19
Wang et al. then designed a thermo-mechanical treatment to produce a wider distribution
of precipitated austenite sizes in Fe-9Mn-3Ni-1.4Al-0.01C [28] than was available in their
previous study [39]. Before tempering, the steel was cold-rolled to increase the density of
nucleation sites for austenite. The associated nucleation energies for defects introduced by cold-
rolling vary; for example, high-angle boundaries have a higher interfacial energy than low-angle
boundaries, and thus lower the energy barrier for austenite nucleation during tempering. By
creating more nucleation sites with varied associated nucleation energies, not only was austenite
precipitation accelerated in comparison to the as-quenched material and tempered material, but
the size of austenite varied as well due to staggered nucleation rates. The austenite that
precipitated early had more time to grow than later nuclei, and so there was a wider range of
austenite size present after tempering. As in the as-quenched material, the cold-rolled steel also
experienced a spectral TRIP effect; furthermore, the mechanical properties were further
improved as a result of the enhanced spectral TRIP compared to as-quenched material that was
tempered to produce a similar volume fraction of austenite.
The spectral TRIP effect was also observed in austenite with varied composition. Yuan et
al. studied an Fe-13.6Cr-0.44C (wt.%) steel that was Q+P treated by austenitizing at 1150°C,
water quenching, and tempering at 300°C, 400°C, and 500°C for 1-30 minutes [41]. Some
austenite remained after the initial water quench (8-20 vol%), and as seen by Sherman et al.,
Morito et al., and Morsdorf et al. [34–36], the C segregated out of the martensite to
austenite/martensite and martensite/martensite interfaces during the quench. With subsequent
tempering, additional C segregated to this new C-rich layer, which reverted to austenite with a
different C content than the retained austenite. After quenching from the tempering temperature,
the final microstructure contained thermally stable austenite with varied sizes and C
20
concentrations. Because the composition of the austenite varied in C, the mechanical stability of
the austenite was also varied. Overall, the composition-aided spectral TRIP effect in these steels
improved the toughness of these steels.
Compositionally distinct layers in austenite have also been produced in low-C Mn
maraging steels. Dmitrieva et al. austenitized and water-quenched an Fe-12Mn-2Ni-0.05C steel
to produce a sample containing a mixture of retained austenite and martensite, both with the
same nominal composition [42]. The steel was then tempered for 48 hours at 450°C to produce
additional Mn-rich austenite. The austenite composition profiles were measured using APT, and
the Mn diffusion kinetics modelled using DICTRA. They found that the austenite in the final
material had two distinct compositions: the nominal composition (retained austenite) and a Mn-
rich region between the retained austenite and the ferrite matrix. The diffusion of Mn at 450°C in
austenite is several orders of magnitude slower than the diffusion of Mn in martensite, which led
to the formation of a Mn-rich boundary at the austenite/martensite interface. This region
continues to enrich until it reaches a local equilibrium, and then the interface grows into the
martensite at the composition dictated by the local equilibrium. After quenching, the final
microstructure consists of two compositionally distinct layers of austenite, the austenite retained
from the first quench and the austenite that formed during the temper.
It is thus seen that in low-C steels there are nanometer-scale features that arise during
quenching and tempering that can have a visible and substantial effect on mechanical properties.
A spectral TRIP effect can be produced in steels with austenite of varied size and composition
that can increase the toughness of the steel. Nanoscale composition variation has been measured
in tempered C-alloyed and Mn-alloyed steels (with advanced experimental capabilities available
in the past two decades), that give rise to this composition and size-based spectral TRIP effect
21
and enable enhanced understanding of the TRIP and TWIP phenomena. Together with improved
computational capabilities, such understanding enables pathways to improved advanced high
strength-high toughness steel design.
2.4: Fe-10Ni-0.1C steels and the QLT Heat Treatment
In this section, we focus on the characterization and optimization that led to the 10Ni
QLT heat treatment and the current understanding of microstructure evolution during the QLT
heat treatment, all of which is directly relevant to the scope of the current effort. Alloy design,
alloy production, and heat-treatment optimization within the family of Ni-containing steels
targeted for improved ballistic resistance for ship structure has resided largely within the U.S.
Navy research lab(s) and has historically relied on a combination of laboratory-scale testing for
static properties and performance in field tests for dynamic properties [25]. Microstructural
analysis has largely been on the optical and scanning electron microscopy level, rather than on
detailed fine structures, although recent US Navy-sponsored research at universities and the
Naval Research Laboratory (NRL) has been filling this gap in knowledge [22,23,27–29].
However, the microstructural understanding has lagged the alloy development and heat treatment
optimization processes.
The nominal alloy composition of the 10Ni steel is Fe-10Ni-1.0Mo-0.08V-0.6Mn-0.6Cr,
and the QLT heat treatment is shown schematically in Figure 2.3a. The QLT heat treatment
begins with an austenitization at 800°C for one hour and water quenching (Q). It is given an
initial ‘lamellarizing’ (L) temper at 650°C for 40 minutes and is then water quenched. It is then
given a final temper (T) at 590°C for an hour and again water quenched. As discussed above, the
lower the tempering temperature, the more Ni-rich is the precipitated austenite (Figure 2.3b).
22
Thus, this treatment significantly decreases the Ms temperature of the austenite through Ni-
enriched austenite (Figure 2.3c).
Figure 2.3: a) The QLT heat treatment- the steel is first austenitized at 800°C for 60 minutes and water quenched; tempered at
650°C for 40 minutes and water quenched; and then tempered at 590°C and quenched. b) The Fe-Ni binary equilibrium phase
diagram with each QLT treatment temperature marked [30]. c) The Fe-Ni Ms temperature diagram with the equilibrium Ni
content in austenite at each stage of QLT marked [43].
The microstructural evolution during this QLT heat treatment is schematically illustrated
in Figure 2.4. The austenitization treatment is sufficient to convert the steel into polycrystalline
austenite with a fraction of undissolved MC and M2C carbides left in it; upon water quenching to
room temperature, a lath martensite microstructure is observed. The L-stage temper produces
moderately Ni-rich austenite which may be thermally unstable upon water quenching. The T-
stage temper produces additional austenite that is more Ni-rich than L because of the lower
23
tempering temperature, and so this generation of austenite is more likely to be thermally stable
than that resulting from the L temper. The end result of this multi-step process is a tempered steel
with a large volume fraction of thermally stable austenite (around 18 vol% as measured by Jain
et al. [28] and Zhang [25]).
Figure 2.4: Overview of the QLT treatment. At the austenitizing temperature, the steel is fully austenitic (with some undissolved
fine carbides). When quenched, it forms a lath martensite structure. The L-temper produces a high volume of precipitated
austenite, though the thermal stability of the austenite is not fully understood. The T-temper produces additional austenite
precipitation that is more Ni-rich than L-generation, and the final result is the presence of an appreciable fraction of thermally
stable austenite.
It is worth noting that prior to heat treating, these steels are hot rolled into plates and
macro-segregation resulting from solidification manifests in the rolled plate as a banded structure
of solute-rich and solute-poor bands [19] that are on a scale too large to be eliminated by any
form of solid-state heat treatment. Therefore, these regions remain distinct through
austenitization and subsequent tempering and will influence austenite precipitation and have
been shown to have an effect on failure mechanisms (e.g. Charpy fracture surfaces [19]).
The 10Ni QLT alloy composition and heat treatment combination was isolated through
ballistic field tests that were conducted on several alloys ranging from 2.5 to 10 wt.% Ni in
composition that had been subjected to heat treatments ranging from single isothermal tempers at
24
different tempering temperatures through various multi-temper heat treatments [25]. Ballistic
resistance in these field tests was quantified by a ‘V50’ metric. Namely, given a particular
caliber projectile, plate specification, and firing distance, the V50 is the velocity at which half of
the projectiles fully penetrate the plate. The results of these ballistic tests showed that the
combination of a 10Ni alloy with the QLT treatment as described above showed superior
ballistic resistance while still maintaining high static tensile strength, elongation, and Charpy
impact energy (Figure 2.5).
Figure 2.5: Ballistic field test results that led to the combination of a 10Ni alloy and the QLT treatment, as reported by Zhang
[25].
Isheim et al. [27] performed microstructural analyses on of some of the precursors to the
10Ni QLT alloy, including a 4.5Ni and a 6.5Ni alloy subjected to a three-temper treatment and a
10Ni alloy subjected to heat-treatment with different L and T temperatures. Comparing APT and
TEM results with DICTRA simulations, they determined that the higher temperature L-tempers
were largely responsible for Ni partitioning and most of the austenite volume in the final state;
25
the lower temperature T-temper merely allowed the precipitation of a new layer of austenite
around L-austenite, but was not long enough to equilibrate the Ni content throughout the
austenite. This correlated with the SEM results of Zhang [25] where the precipitated austenite
volume fraction in the optimized 10Ni QLT treatment was compared to just the microstructure
after QL and after a QT treatment (QLT without the intermediate L). The QT treatment had very
little austenite in comparison to QL and QLT, and so Zhang suggested that the T-temper was
primarily responsible for austenite precipitation within any fresh martensite after L (i.e. in
thermally unstable L precipitates) and their additional growth.
The effects of each stage of the optimized QLT treatment on the microstructure evolution
in the 10Ni alloy was recently examined by Jain et al. by using APT and synchrotron techniques
to compare austenite precipitation in material subjected to the QT, QL, and QLT schedules [28].
Using dilatometry, they determined that some of the austenite precipitated during the L stage was
thermally unstable during the quench and that the final volume fraction was about 8% austenite
by volume. The QT sample had only 3 vol% austenite, and the QLT treatment resulted in 18
vol% austenite. Thus, the increase in austenite volume seen between QL and QLT could not be
ascribed simply to nucleation and growth of austenite precipitates by long-range diffusion of Ni
out of the ferrite during the T treatment, as QT alone had far less austenite than the difference
between QL and QLT. APT showed the composition of the L-generation austenite to be between
12 and 16 wt.% Ni. Two chemically distinct generations of austenite were seen in QLT
specimens, one matching the composition of the L-generation believed to be retained L-
austenite, and a second generation about 21 wt.% Ni that formed during T. Therefore, the
majority of the increased volume in thermally stable austenite between QL and QLT was through
26
the precipitation of new austenite in the L-generation fresh martensite pockets. DICTRA models
supported this result, as seen in Figure 2.6.
Figure 2.6: DICTRA model results compared to APT results, taken from [28], confirm austenite precipitation in Ni-rich fresh
martensite pockets that are formed by the transformation of some of the L-generation austenite to martensite during the quench
from the L temperature.
Additionally, Gupta and Kumar [29], Wang and Kumar [22], Isheim et al. [27], and Jain
et al. [28] each reported the presence of Mo-rich MC and M2C carbides in tempered samples in
these alloys. The steels they considered were compositionally the same but varied in Ni content
and tempering temperature (ranging from 450-650°C). Thus, some of the C in these steels is
removed from the matrix via carbide precipitation and growth. The relationship between C
partitioning and carbide formation kinetics as a function of heat treatment and hence its role in
affecting austenite stability (in particular in the L temper) has not been clarified.
2.5: Dynamic Deformation of Steels and Other Alloys
While the quasi-static mechanical properties of QLT-treated 10Ni steel are relevant, more
important to this effort is the alloy’s response to high strain-rate deformation because of the
intended application as hull and ship deck armor. In particular, we consider the phenomenon of
Adiabatic Shear Banding (ASB), an acute form of shear localization often seen in ballistic
27
applications that implies loss of structural integrity. In high strain-rate deformation, the
temperature that arises from the localized plastic deformation can be quite significant (on the
order of 500°C [44–48]), as there is not adequate time for the heat generated by deformation to
dissipate into the bulk (thus, adiabatic). If the work-hardening from plastic deformation is
overridden by the thermal softening due to this temperature increase, the localization becomes
unstable, concentrating the deformation zone to an increasingly smaller, hotter, and softer band.
Inside this band, the presence of extreme conditions exemplified by a combination of large strain
rate, large strain, and high temperature, enables deformation mechanisms that are otherwise not
active elsewhere in the material and certainly not present in quasi-static deformation. These
include temperature- and/or strain-induced phase transformations (usually diffusionless, as the
duration is of the order of milliseconds), and deformation modes like micro-twinning, carbide
plastic deformation and dissolution, and mechanical recrystallization. Eventually, the material
fractures by cracks forming and linking within this adiabatic shear band. Thus, understanding the
microstructural characteristics that discourage shear localization is central to designing materials
that can better resist ballistic impact. While only limited analysis exists on the microstructural
evolution in 10Ni steels during high strain-rate deformation in general, and even less in 10Ni
steel subjected to the QLT treatment, we can use the extensive literature on high strain rate
deformation in other steels and non-ferrous alloys to guide our thinking in this effort.
Adiabatic shear banding was first described in steel by Zener and Hollomon in 1944 [49],
but it has since been observed and studied in a wide range of materials including various steels,
pure iron, titanium and titanium alloys, aluminum and aluminum alloys, copper, brass, tantalum,
uranium, magnesium, and zirconium [49–57]. Generally, ASBs are seen in materials exhibiting a
low strain-hardening rate, low strain-rate sensitivity, low thermal conductivity, and high thermal
28
softening rates; the easier it is for thermal softening to overcome work hardening for a given
material, the more prone the sample is to localization. In terms of application, shear banding can
occur in ballistic plate armor; self-sharpening projectiles; high-velocity machining, shaping, and
forming; and extrusion and punching [52–54]. For a given alloy, there are several microstructural
factors that can affect shear localization, including defect structure, grain size, and prior strain
hardening [58–62]. ASB formation also depends heavily on specimen geometry, surface defects
(e.g. scratches), initial temperature, stress state, and strain rate [46,53,54,61–63].
Strain rates in the range of 103 (similar to ballistic impact) can be achieved with a
selection of setups, including the Kolsky (Split-Hopkinson) pressure/torsion bars, Taylor impact
experiments, explosive tubes, and ballistic field tests [52,64]. Furthermore, various sample
geometries can be used (e.g. top-hat [60]) to induce shear in predictable regions with
predetermined total strain and strain rates. Each can be paired with high-speed cameras and
thermal sensors to determine global strain, local strain, local temperature, and localization
progression as a function of time. For further reading on experimental setups, refer to Ramesh
[64], Walley [52], and Lindholm [65].
Accurately characterizing local strain, temperature, and microstructural evolution during
high strain-rate deformation is difficult due to the small scale and the rapid nature of shear
localization. Shear bands propagate between 250-1200m/s [44,46,48], and the most developed
portion of the band can be as narrow as 5µm [53]. Experimental work to characterize the
sequence of events and prevailing conditions in the shear band frequently utilizes high-speed
photography and in-situ infrared imaging but is limited by small time-windows and spatial
resolution. Nevertheless, the information obtained from such experiments provides insights into
rationalizing the post-deformation microstructure in the ASBs, for understanding microstructure
29
evolution, and is useful as experimental input for modelling high strain rate deformation and
shear localization.
High-speed cameras have been used to estimate local strain and ASB propagation speed.
Using painted horizontal stripes on a torsional Kolsky sample of HY-100 steel, Marchand and
Duffy [44] were able to measure local shear strains ranging from 500-1900%; they noted that the
strain was inhomogeneous during band formation and during propagation, likely due to
perturbations of the thickness of the cylinder walls [44]. Therefore, they concluded, there is a
strong relationship between geometry and shear band propagation. In addition, they estimated the
propagation speed to be around 500 m/s, though possibly 250m/s, if it propagated along two
fronts. Zhou et al. [46] and Guduru et al. [48] measured with higher confidence a propagation
speed up to 1100 m/s in C300 steel.
Local temperatures as a function of local strain within shear bands have also been
estimated with increasing accuracy through the use of high-speed cameras in conjunction with
arrays of infrared detectors. Marchand and Duffy and Giovanola each measured the increase in
temperature in the shear band in HY100 steel and found it to be between 550-1000°C, although
they were limited by the spatial resolution of their detectors [44,45]. With slightly improved
resolution, Zhou et al. showed that in a C300 maraging steel, the temperature was elevated in a
region between 200-300µm, much larger than the band itself, with a 1400°C maximum increase
in the most concentrated portion of the band [46,47]. Guduru et al. performed similar
experiments on C300 and measured a lower bound of 600°C, but importantly noted that ASBs
are not thermally homogeneous: shear bands can form periodic ‘hot-spots,’ and therefore may
have quite varied internal structure [48]. From a theoretical approach, temperature increases can
30
be estimated using the constitutive equations of the Zerilli-Armstrong model [66], for example,
as Kad et al. have done with ultra-fine grain zirconium [67].
Historically, shear bands have been categorized into two types: transformed bands and
deformed bands [52,53,63,68]. Transformed bands were observed in materials that were capable
of phase transformation upon heating and rapid cooling and were thought to have undergone a
cycle of phase changes during the rapid temperature spike during ASB formation and the rapid
heat extraction that follows as the heat is dissipated to the surrounding material (e.g. ferritic
steels up-quenching to austenite in the band from the temperature spike and then transforming to
martensite following the rapid temperature drop). However, actual evidence for temperature-
induced phase transformation was largely circumstantial and relied on macroscopic observations;
transformed bands were resistant to etching and have been referred to as ‘white bands’ and they
had a measurable hardness increase across the band, which was used as evidence for a phase
change. On the other hand, deformed bands did not etch a different color, and thus were thought
to have had no phase transformation. This historical perspective has been demonstrated to be not
entirely correct: ‘deformed’ bands can in fact have phase transformation, and ‘transformed’
bands may have none at all [51,52,57,61,68]. The increased hardness previously used as an
identifier of a transformed band is now believed to be due to substantial grain refinement
processes [51,57]. As localization continues, shear bands continue to narrow [44,45,58,69]; it is
therefore thought that ‘deformed’ bands may be those that have not progressed so far as to
mechanically recrystallize [57,70].
The mechanical properties and microstructures of Cu, Ta, Al-Li, Ti, and steel are all quite
different, but share a common microstructural feature within the shear band in that they all have
equiaxed grains between 100-300nm in diameter [51,71]. It is thought that each undergoes a
31
similar microstructural process during shear localization: first, a plastic elongation in the shear
direction; next, a breakup of these elongated grains into sub-grains separated by dislocation cell-
walls; lastly, a rotational deformation mechanism that results in a fine-grained equiaxed structure
with low dislocation density [51,57,67,71,72]. This final step of ‘recrystallization’ is thought to
be facilitated by the high temperatures in the shear band, which allows for the small-scale
diffusion necessary for the subgrains to rotate. Note that this model does not require phase
transformations; indeed, in some steels there is strong evidence that no up-quenching/re-
quenching of martensite occurs within these bands [57,58,69,73]. The increased hardness and
white etching seen in these bands are thus attributed to a well-defined region in which the
average grain size is significantly smaller than the edges of localization where grains have not
refined but are elongated by intense shear [57,58,69]. Deformed bands, therefore, are those that
have not undergone this dynamic recrystallization process. However, it is worth noting that
mechanically-induced phase transformations (TRIP) may still occur within the shear band (e.g.
austenite to martensite) [51,57].
Two studies in particular clearly demonstrate this dynamic recrystallization process. Xue
and Gray observed a grain-refinement process in austenitic 316L stainless steel [58,69]. The steel
was deformed using compression Kolsky-bars on top-hat specimens, which were loaded with
strain pulses of varied duration but equal magnitude to produce various stages of shear band
propagation and growth under similar loading conditions. TEM specimens were cut and thinned
from the sheared area for further analysis. The initial stages of shear localization were marked by
dislocation cell structures and a lath-like microstructure elongated along the shear axis. In a well-
developed band sample, there were three distinct microstructural regions with respect to the
band. At the furthest edges of localization, there was a ‘dislocation-avalanche’ structure, where
32
the dislocation cell structures were annihilated parallel to the elongation but grew into thick
tangles at the long ends. Closer to the band, an elongated subgrain region was observed, where
the grain showed a very high aspect ratio along the shear axis (~30). This elongation mechanism
was believed to be facilitated through the activation and multiplication of twins that were along
preferential orientations. Lastly, the center of the shear band was marked by a fine equiaxed
subgrain structure with diameters ranging from 20-100nm. This region was about 10µm in width.
Thus, the full microstructural evolution during shear localization was characterized: first, grains
break into dislocation cells; the cells elongate and further refine; and finally break down into
equiaxed subgrains.
This model was theoretically rationalized by Kad et al. [67] in ultrafine-grain zirconium.
Specimens were cut in the top-hat geometry and deformed using compression Kolsky-bar
experiments; global plastic strain was constrained by changing the protruding hat height. As in
the work of Xue and Gray, they observed a sharply contrasted microstructure within the center of
the shear band denoted by fine, equiaxed grains. To determine if such a dynamic recrystallization
process was feasible (specifically, if equiaxed subgrains could rotate during deformation to form
the high misorientation boundaries seen using SADP and darkfield), they used the Zerelli-
Armstrong model [66] and experimental stress/strain curves to determine the temperature rise in
the band. The temperature was predicted to increase by over 600°C to become about 40% of the
melting temperature. They then estimated the time it would take for enough grain boundary
diffusion at that temperature to rotate hexagonal subgrains of diameter 100-300nm by 30°, and
found it to be between 50-100µs. This time falls within the range of shear band development,
thus reinforcing the hypothesis for dynamic recrystallization.
33
The loss of load-bearing capacity of a material has been shown to be due to void
coalescence that often occurs during ASB propagation and growth [56,57,74]. Xu et al. [56] used
SEM analysis on low-C steel that was deformed using a torsional Kolsky-bar; by interrupting the
test at different stages, they were able to correlate the global stress-strain curve with the
microstructure at different stages of ASB propagation and growth. The sharp drop in load-
bearing capacity more strongly correlated with the formation and growth of micro-voids than
with ASB formation. Cracks were seen to initiate at interfaces, such as grain boundaries,
inclusions, and precipitates. A similar observation was made in Ti-6Al-4V in radially-collapsed
thick-walled cylinders [57]. In this material, ellipsoidal voids were seen to form in the thermally
softened shear band and grow to the edges of the band, creating a perforated structure. The
nucleation and growth of these voids correlated strongly with the loss of load-bearing capacity.
Rigorous microstructural analysis at various stages of deformation has not been
performed on 10Ni QLT and the work on its previous iterations is limited. In total, there are
three principle studies to consider.
Wang and Kumar characterized the microstructural evolution of the Fe-10Ni-1.0Mo-
0.6Mn-0.6Cr-0.08V-0.1C (wt.%) steel subjected to an unspecified heat treatment during quasi-
static and dynamic deformation [22]. The initial microstructure was composed of lath martensite
with a high dislocation density, no discernable retained austenite, and a dispersion of MC and
M2C carbides. High strain rate deformation with strain rates between 1.0x103s
-1 and 4.5x10
3s
-
1were performed on cuboidal specimens using a compression Kolsky-bar. A series of strains and
strain rates were produced, and shear bands were found to occur in samples deformed over 40%
and at strain rates over 2000s-1
. Nano-indentation arrays across the well-developed portion of the
shear band showed that the hardness increased from about 6.5GPa to about 7.5GPa across the
34
width of the band. TEM liftouts from the localized region showed that outside of the band, the
laths were oriented along the shear direction, while the microstructure in the band itself was
comprised of fine, equiaxed grains. Coupled EDS and MDP showed that some of these equiaxed
grains had Ni content close to the nominal concentration yet were austenitic. If these grains had
been austenite in the undeformed microstructure, they would be predicted to be much more Ni-
rich by the Fe-Ni phase diagram. Therefore, their austenitic structure in the band suggests that
they were martensite that reverted to austenite from the high temperature spike during shear
localization, while their fine size was hypothesized as preventing them from transforming back
to martensite during the rapid cooling that followed. Also present were ferritic grains that
appeared to be highly twinned. There were no carbides found inside the band, either because
they were plastically sheared into very thin ligaments that were not recognizable or they re-
dissolved. Another specimen consisted only of ferritic subgrains (i.e. no up-quenched austenite),
suggesting that the temperature in the band was heterogeneous.
Gupta and Kumar examined the microstructural evolution during high strain rate
deformation of the Fe-10Ni-1.0Mo-0.6Mn-0.6Cr-0.08V-0.1C steel that was subjected to two
different heat treatments [29]. In both instances, the steel was austenitized at 840°C for 1 hour
and water-quenched. The first material was then tempered at 450°C for 5 hours, and the second
was tempered at 620°C for three hours, quenched, and then tempered at 540°C for one hour. The
samples were deformed using a compression Kolsky-bar at ~1x103 s
-1 and characterized using
TEM. Both steels initially had a lath martensite structure with MC carbides; the two-stage-
temper steel additionally had a lower dislocation density, M2C carbides, and thermally stable
precipitated austenite. The dynamic stress/strain curves for the single-temper specimen had a
higher yield than the two-stage-temper specimen but exhibited strain-softening indicative of
35
severe shear localization. In contrast, the stress/strain curve from the two-temper specimen did
not have strain-softening, suggesting resistance to shear localization. The single-temper
specimen had well-developed shear bands and an SADP pattern with FCC spots, suggesting that
some of the lightly tempered martensite up-quenched to austenite within the band, similar to that
reported by Wang and Kumar [22]. The shear bands in the two-temper specimen were more
diffuse and so they introduced a notch in the specimen to generate a narrow, well-defined shear
band. TEM analysis using SADP showed no FCC spots suggesting that the retained austenite in
the initial structure had dynamically transformed to martensite during the process of shear
localization. Thus, it was concluded that a TRIP effect provided from precipitated austenite may
help dissuade shear localization.
As discussed earlier in this chapter, Zhang [25] performed numerous field tests to
determine the ballistic resistance of a series of alloy compositions and heat treatment
combinations and those results were presented in Figure 2.5. The 10Ni steel with the QLT
treatment that is the focus of this work was shown to have superior ballistic resistance as
compared to alloys with less Ni but subjected to the same QLT treatment, and other 10Ni alloys
subjected to various single and other multi-stage tempers. Inferior steels were found to form
ASB and fail via plugging; the 10Ni QLT steel was found to resist ASB formation and instead
deformation was distributed by plate bulging. The retained austenite volumes after heat treatment
and in the impact zone were measured using a vibrating sample magnetometer. While the
undeformed material had around 18 vol% austenite, the impact zone was found to have
negligible austenite. Zhang suggested that the mechanical transformation of this large volume of
thermally stable austenite served to suppress shear localization and instead dissipated the impact
energy over a much larger volume of material.
36
2.6: Scope of this Effort
While the ballistic response of an Fe-10Ni-1.0Mo-0.6Mn-0.6Cr-0.08V-0.1C plate
subjected to the QLT heat treatment was clearly superior compared to the other variants tested,
the following fundamental questions remain inadequately unanswered:
1) What is the fine microstructure of 10Ni QLT? How did it evolve in the context of
composition, size, and thermal stability of the precipitated austenite and how do these
relate to the two-stage tempering process?
2) How does the austenite in 10Ni QLT mechanically evolve during high strain-rate
deformation? Specifically, does deformation-induced austenite to martensite
transformation occur in the dynamically deformed material and if it does, how
prevalent is it?
Thus, our experimental effort can be viewed as being composed of two sequential parts:
first, a study of the relationship between heat treatment and microstructure, focusing on
how precipitated austenite size, composition, and morphology relate to thermal stability;
and second, a comparative analysis of the microstructural evolution of 10Ni QLT during
high strain-rate deformation to determine if the thermally stable austenite is mechanically
unstable and if so, to what extent or under what circumstances.
37
Chapter 3: Experimental Procedure
3.1: Materials and Heat Treatment Schedules
Table 3.1: Nominal alloy composition
Fe Ni Mo Mn Cr C V
wt.%
(Given)
87.6 10.0 1.0 0.6 0.6 0.1 0.08
at.% (Calc) 86.7 10.6 0.7 0.7 0.7 0.5 0.1
The nominal composition of the alloy examined in this study was Fe-10Ni-1.0Mo-
0.6Mn-0.6Cr-0.08V-0.1C (wt.%) (Table 3.1). This steel (referred to as the “10Ni steel” later)
was first melted in vacuum using induction heating, cast into billets that were then hot rolled into
plates, and air cooled; samples were then cut from the plate and heat treated. The focus is on the
QLT heat treatment that consists of three stages:
1) An austenitizing step at 800°C for 1 hour followed by a room-temperature water
Quench (the Q step)
2) A high temperature ‘Lamellarizing’ temper at 650°C for 40 minutes followed by a
water quench (the L step)
3) A low temperature Temper at 590°C for 60 minutes followed by a water quench.
To understand the effects of tempering time and temperature on microstructure evolution,
specifically austenite composition, size, dispersion, and thermal stability, we additionally
examined a series of single-stage heat treatments at the L temperature (650°C) and at an
exaggerated low temperature designated T’ (540°C). Specifically, as shown in Figure 3.1, we
38
consider a 40 minute and 25 hour L treatment with water quench (QL and Q25L respectively)
and 5, 25, 125, and 336 hour T’ treatments with water quench (Q5T’, Q25T’, Q125T’, and
Q336T’, respectively). This range of single-temper treatments allows us to examine and
understand the kinetics of austenite precipitation at high and low tempering temperatures. This
also allows us to cumulatively piece the individual stages of QLT together (in the AQ, QL, and
QLT samples) to understand directly the effects of each stage on the final QLT microstructure.
Figure 3.1: Heat treatment schedules considered. All treatments begin with austenitization and quench (AQ). Two single-stage
tempering temperatures (650°C-L and 540°C-T’) with various lengths are analyzed, as well as the two-stage temper QLT
A large piece of QLT-treated plate was supplied by our collaborators at NSWC-CD
(Naval Surface Warfare Center-Carderock Division); in addition, fractured Charpy test
specimens of Q25L and Q5T’, Q25T’, Q125T’, and Q336T’were also provided. AQ and QL
samples were important to our understanding of heat treatment and therefore pieces were cut
from the QLT plate and re-heat-treated at Brown University to obtain the AQ and QL specimens.
39
3.2: High Strain Rate Deformation
The experiments that focused on understanding microstructural evolution during high
strain rate deformation were limited to the material subjected to the QLT heat treatment. A
Kolsky (Split-Hopkinson) compression bar setup (Figure 3.2) was used to generate dynamic
stress-strain data, high strain-rate deformed microstructure, shear localization, and adiabatic
shear banding in the QLT specimens. Such experiments produce strain-rates that are comparable
to ballistic impact and can reliably produce predetermined strain and strain-rates. A brief
description of the experiment follows: the sample is placed between two long steel bars (the
incident bar and the transmitted bar); a variable length projectile bar loaded into a pressurized air
gun with controlled firing pressure can be impacted on one end of the incident bar. The resulting
strain pulse travels through the incident bar to the sample, and the transmitted and reflected
pulses are each recorded using an oscilloscope. These strain pulses, in combination with initial
and final sample dimensions, are then used to calculate the stress/strain response of the
specimen. Detailed explanations of Kolsky-bar setups, including underlying theory and
derivations, can be found in Ramesh [64], Walley [52], and Lindholm [65].
Figure 3.2: Kolsky-bar setup. The sample is sandwiched between two bars; a projectile is fired at the end of one bar, and the
strain pulses through the incident bar, reflected back through the incident bar, and transmitted through the transmission bar are
recorded with strain gauges. A stress/strain curve is thus derived from these pulses.
40
Kolsky-bar experiments have only two input variables, projectile firing pressure and
projectile length, but can achieve a wide range of total strains and strain rates. Therefore, an
extensive amount of calibration is necessary to produce a deformed sample with a desired strain
and strain rate. A wide range of projectile and pressure combinations were thus initially
examined using 4140 tempered martensite to calibrate the Kolsky-bar and to hone in on the
range of parameters that was considered appropriate to achieve the desired strain rate and strain
in the QLT-treated 10Ni-steel samples. 4140 steel was chosen due to its availability, ease of
machinability, and the similar range of static mechanical properties typical of the QLT-treated
10Ni steel. To determine the effects of bar length on strain and strain rate, 15.3cm, 20.3cm,
30.5cm, and 40.6cm projectile bars were used. A wide range of pressures (20-100psi) was used
for each bar that yielded velocities ranging from 8-36m/s. The strain and strain rates for each run
were plotted as functions of the bar and loading pressure used, thus providing a rough estimate of
experimental parameters needed for testing the QLT-heat treated steel.
A few sample geometries were considered, including cuboidal samples, cylindrical
samples, and notched samples. Samples with a cylindrical geometry were selected due to their
radial symmetry that helps maintain experimental consistency. Notched samples were briefly
used because the notch served as an initiation point for shear localization; however, experiments
on un-notched samples were also able to produce shear banding, and so notches were determined
to be an unnecessary complication in sample uniformity.
Cylindrical samples with a diameter of 6.35mm and a height of 9.52mm and subjected to
the QLT heat treatment were cut from bar stock using electro discharge machining (EDM). The
anvils and bars were both of hardened steel and had a diameter of 2.54mm; the projectile bar was
also a steel bar and 30.5cm long. Projectile velocities ranged from 13m/s to 27m/s (which
41
equates to chamber pressures ranging from 47-95psi). Strain and strain-rate were varied on the
QLT specimens to capture different stages of shear localization, including: homogeneous
deformation (i.e. no localization), localization but no developed shear bands, well-developed
shear bands that decayed within the sample, shear bands that ran through the sample, and
fractured samples. Deformed samples were cut in half vertically with EDM or slow-speed
diamond saw and further prepared for characterization.
3.3: Microstructure Characterization
3.3.1:Optical Microscopy:
The microstructure of both undeformed and deformed samples was too fine to fully
define with optical microscopy, though optical microscopy was useful to each of these two
categories for specific reasons. Metallographic samples were prepared by grinding with SiC
paper, polishing with 1.0µm, 0.3µm, and 0.05µm alumina solution, and etching with a solution
of 2-5% nitric acid in methanol. One use for optical microscopy of the undeformed samples was
for analyzing banded structural features, such as those observed by Zhang [25]. During hot
rolling and subsequent treatment, alloying elements may segregate to alloy-rich and alloy-poor
bands, which influences the precipitation of austenite during tempering, and while these banded
features were also observed in SEM, low magnification light microscopy was also used to assess
these features on the millimeter scale. This was most evident in Q5T’, the sample with the least
amount of austenite precipitation. The primary purpose for using optical microscopy to
characterize the deformed specimens was to determine the presence and development of
adiabatic shear bands in deformed QLT samples. Shear bands on etched samples appear as white
bands, and so they can be clearly seen running through the sample if present. Multiple deformed
samples were cut in half to determine the progression of shear bands within, and this information
42
was used to iterate and tune the Kolsky-bar experiments. Furthermore, the presence of a banded
structure in the millimeter scale (described above) in these materials enable an estimation of
localized strain, as they serve as markers and their displacement across a shear band are can be
readily observed. Lastly, optical microscopy acted as a screening tool to isolate suitable
specimens for further analysis using electron microscopy techniques.
3.3.2: Scanning Electron Microscopy (SEM):
SEM was used to qualitatively compare the size and dispersion of austenite as a function
of heat treatment schedules as well as to guide smaller-scale analysis with TEM and APT. SEM
samples were prepared in the same fashion as optical microscopy samples: ground with SiC
paper; polished with 1.0, 0.3, and 0.05µm alumina; and etched with a solution of 2-5% nitric acid
in methanol. Secondary electron and back scattered SEM images were obtained for all heat
treatments to compare differences in: precipitation across the rolling bands, morphology of
precipitated austenite (e.g. spherical, lamellar, etc.), and variations in austenite size. In deformed
material, SEM was used to examine austenite morphology as a function of proximity to the shear
band, as well as for targeted TEM liftouts, which is discussed later in this section in more detail.
The two microscopes used were a LEO 1530VP and a FEI Helios equipped with electron
backscatter diffraction (EBSD) and focused ion beam (FIB) systems.
Electron backscatter diffraction was attempted to quantify the amount of thermally stable
austenite in given heat treatments. Several sample preparation techniques were utilized,
including mechanical polishing with alumina down to 0.05µm; polishing with 1µm alumina
followed by electropolishing in a 10% solution of perchloric acid in acetic acid; and polishing
with a 1µm alumina polish followed by a polycrystalline diamond polish, but the outcome was
43
not satisfactory. TEM provided results with a higher point density and reliability than we could
have hoped for using EBSD, so this aspect was not pursued.
3.3.3: Transmission Electron Microscopy (TEM):
Transmission electron microscopy was extensively used to follow the composition,
structure, and size of the phases present as a function of the various heat treatments examined in
this effort, as well as to examine the microstructure of the dynamically deformed specimens. For
undeformed materials, samples were cut from the bulk into small plates using EDM and
mechanically thinned to widths between 50-100µm by grinding with SiC paper. They were then
polished using 1µm alumina and electrochemically thinned to perforation using a Tenupol 5
twin-jet polisher at 20-25V at -30°C with a solution of 20% nitric acid in methanol. For
deformed QLT specimens, FIB liftouts were made with specific geometries and at distinct
locations with respect to the adiabatic shear band including: far from the band (homogeneous
deformation zone), just ahead of the band (some localization without well-defined ASB), parallel
to the well-defined band, at an angle to (but in) the band, and across the band normal to the shear
band plane. Collectively, these various lift-outs provided snapshots of the microstructure at
various stages of localization.
Initial TEM analysis was done using a Philips CM20 equipped with an Oxford energy-
dispersive X-ray spectrometry (EDS) detector; finer-scale analysis was done using a JEOL
2100F in scanning mode (STEM) with an equipped Oxford EDS detector
Electron diffraction was used to determine the structure/identity of the phases in these
multiphase microstructure (e.g. if a precipitate is thermally stable austenite versus martensite) as
well as to examine texture (or lack thereof) across larger polycrystalline regions (e.g. randomly
oriented). For most treatments, even the smallest selected-area apertures were much larger than
44
individual austenite precipitates, and so selected area diffraction (SADP) was not useful to
determine their thermal stability. However, SADP in conjunction with darkfield (DF) imaging
was useful for comparing the degree of misorientation in undeformed microstructure versus
various stages of shear localization. The primary diffraction technique used to characterize
individual austenite precipitates and martensitic laths was microdiffraction (MDP) - condenser
apertures and spot sizes were reduced until the converged beam was smaller than the feature size,
and a diffraction pattern was taken.
Table 3.2: Example table of calibrated diffraction radii
Ring Calibrated Ratio Lattice Index
1 2.70 1.00 FCC 111 2 2.77 1.00 BCC 110 3 3.15 1.17 FCC 200 4 3.88 1.41 BCC 200 5 4.44 1.65 FCC 220 6 4.73 1.71 BCC 112 7 5.18 1.92 FCC 311
8 5.40 2.00 FCC 222 9 5.51 2.00 BCC 220
10 6.13 2.22 BCC 310 11 6.24 2.31 FCC 400 12 6.69 2.43 BCC 222 13 6.79 2.47 FCC 331 14 6.98 2.58 FCC 420 15 7.26 2.63 BCC 321 16 7.65 2.83 FCC 422 17 7.76 2.82 BCC 400 18 8.10 3.00 FCC 511, 333
19 8.10 3.00 BCC 330 20 8.21 2.98 BCC 411
Patterns were indexed, including measuring diffraction radii and angles. A calibration
table was created for FCC and BCC Fe (austenite versus ferrite and martensite), which was then
used to calculate all diffraction radii for both crystal structures (an example given in Table 3.2).
45
When a crystal was on a zone axis that had similar symmetries for both FCC and BCC (e.g. BCC
and FCC [100] both have a square pattern,), the calibrated table was used to isolate the phase
through the difference in lattice spacing. The table was recalibrated for each individual session to
ensure no mistakes were made.
Lastly, these steels are highly magnetic, which makes tilting the specimen difficult as
significant image shifting occurs. Therefore, most regions of interest were found by looking for
grains already on a zone axis, rather than through using diffraction to tilt to a particular grain's
zone axis. Usually, this meant doing characterization at 20°, the tilt required for EDS analysis.
Very few tilting experiments were performed.
EDS was used extensively in this effort. In undeformed material, single-point EDS was
used to measure Ni content in ferrite and precipitated austenite, allowing for tracking the Ni
content in the precipitated austenite as a function of tempering temperature and time, as well as
understanding the thermal stability of this phase as a function of size and Ni level. EDS was also
used to measure the distribution of Ni by compounding hundreds of single-point measurements
over a large area. These distributions showed on a larger scale the Ni depletion in ferrite as a
function of tempering time and the Ni content approaching the equilibrium phase diagram values
for longer times.
In deformed QLT samples, Ni content served as the link to the undeformed
microstructure. As shear localization during dynamic deformation is essentially a diffusionless
process, a feature with 15 wt.% Ni in the deformed microstructure must have evolved from a 15
wt.% Ni entity in the undeformed state. In the case of QLT, this meant that the information
obtained on the undeformed microstructure could be mapped onto the deformed microstructure
through its Ni content. Thus, by examining multiple stages of shear localization, we could
46
understand the microstructural evolution in the baseline QLT material during dynamic
deformation.
Figure 3.3: a) an example of EDS analysis of QLT using a manually condensed beam (rough size marked by pink circles) in the
CM20. Austenite precipitates include roughly one datum point each. b) STEM EDS on the JEOL 2100F allows for a high enough
data point density to create linear gradient plots within grains
In the early stages of this study, the CM20 was used to measure Ni content with
converged beam EDS. However, the smallest possible spot size was sometimes larger than the
features we sought to characterize- this was especially the case for shear band samples, where
features could be as small as 20nm. Therefore, EDS on the JEOL in the STEM mode was
employed instead. STEM mode has a rastering 1.5nm spot that creates an image, far smaller than
the features of interest, and by selecting points on the interface, the microscope focuses the beam
47
automatically to the desired location. This allowed the number of data points taken, specifically,
the density of data points, to increase from ~20 points per session to ~1000 points per session
(Figure 3.3). STEM EDS thus allowed for two particularly useful analyses: arrays of ~200 data
points across ~2µm2 used to create composition distributions (as mentioned above), as well as a
high enough point density within individual austenite precipitates to create linear gradient plots.
It is worth noting that while automated rastered compositional mapping is possible using STEM
EDS, it was not used in this study due to significant sample drift issues.
STEM EDS analysis however does have its own set of challenges. While it was possible
to create data sets with hundreds of data points, it was difficult to perform the analysis manually.
STEM images needed to be overlaid on BF images, the data points had to be manually placed
(including compensating for drift during the EDS collection), and then the data points had to be
colored according to composition and a binning scheme. Single images would take days to
produce, and it did not allow for secondary analysis like linear gradient plots; adjusting color
schemes from different data sets to maximize contrast was also not feasible. Therefore, a Matlab
script was written and developed to help analyze the large EDS data sets.
The main purpose of the Matlab script (Figure 3.4) was to produce useful figures from
the EDS data sets, including performing secondary analysis such as linear gradient plots.
Experimentally, each individual collection of points was related to three STEM images: an image
taken before data collection, the EDS interface overlay of the ‘before’ image with the data points
locations marked, and an image taken after the data were collected (used to calculate drift during
data collection). First, points were manually selected in order on the second image, which saves
the location of the points in STEM space; simultaneously, the script tags each data point with the
corresponding composition data. Next, the script calculates the total drift (in STEM space) using
48
built-in Matlab image correlation functions comparing the before and after images. The total drift
is linearly interpolated and then applied to each datum point; accuracy in the linear estimate was
mitigated through keeping the individual data sets small enough that they can be collected before
significant drift has occurred.
Next, the STEM image is mapped to the BF image using manual feature mapping- a point
in STEM mode is selected, and then the corresponding point in BF is selected. After several pairs
have been created, a built-in Matlab script determines a best-fit linear transformation from
STEM to BF using the pairs. In order to ensure the map meets a threshold of quality, a
quantifiable metric is assigned to it that describes the deviation between the manual pair
selection and calculated best fit map. The manually selected STEM points are mapped to BF
using the calculated affine transformation and then subtracted from their corresponding manual
BF points; the standard deviations of both X and Y components of these vectors are given, and if
unacceptable, the mapping process can be repeated from feature point selection.
Lastly, the EDS data points are mapped from STEM to BF and the final figure is created
by creating colored boxes on the locations of EDS measurements. The boxes are colored using a
preset color scheme based on their corresponding Ni content in order to visually compare Ni
distribution- for example, low-Ni points may be blue while high-Ni points may be red. The
dimensions of the boxes are the error bars for the point’s exact location, and are a function of
total drift during its particular set’s collection and the standard deviation metric from its
STEM/BF map. As the primary focus of the EDS work related to austenite, ferrite and
martensite, the Mo content was used to filter out data points from carbides as most carbides were
Mo-rich - a Mo threshold was entered, and all data points with Mo above that threshold are not
displayed.
49
Figure 3.4: STEM mode images are mapped to BF using manually selected correlative point pairs- i.e. a location is selected on
the STEM image, and then the same point is selected in BF. Several of these are used to create an affine map. Data is collected in
small sets to avoid significant drift during collection; these sets are uploaded in STEM space, mapped to BF mode, and color
coded by composition. Dozens of individual sets create dense composition maps.
These figures can be made of dozens of data collection sets, each with dozens of
individual EDS measurements; collectively, hundreds of data points can be represented in a final
figure. Because the data are all stored after entry, including position, mapping, and composition,
these figures can be quickly reconfigured to new color binning if so desired (for example, to
maximize contrast between two different treatments). Lastly, linear gradients can thus be
calculated: a gradient box is drawn by manually drawing a line (which will be the plot axis) and
50
providing a width. All EDS data points that intersect the gradient box are then plotted along the
length of the line (with corresponding horizontal error bars related to the EDS point error). An
example of this analysis can be seen in Figure 3.5. By creating a script to analyze raw STEM
EDS data, not only were the data more quickly synthesized, but secondary analysis could also be
done.
Figure 3.5: Sample EDS Ni gradient on a Q336T’ austenite precipitate
3.3.4: Atom Probe Tomography (APT)
In addition to EDS, atom probe tomography (APT) was used to measure composition
adjacent to interfaces as well as track the partitioning of minor alloying elements (Mn, Cr) and
light elements like C as a function of heat treatment. APT volumes are relatively small (e.g.
200nm x 50nm x 50nm) but are reliable down to a fraction of an atomic percent. Therefore APT
is a powerful tool when measuring trace element composition and composition at locations of
51
interest that are tens of nanometers in size such as fine carbides and locations adjacent to
interfaces. By combining APT results with extensive SEM and TEM analysis, a full-scale picture
of microstructure as a function of heat treatment was thus developed.
Our access to APT was limited (APT was performed at the Max Planck Institute für
Eisenforschung in Düsseldorf, Germany), and so the samples examined were carefully chosen.
The QLT steps (AQ, QL, and QLT) were analyzed to provide a comprehensive picture of the
effects of each heat treatment step on the final microstructure. For single stage treatments, Q25L
was chosen to contrast with QL, and Q25T’ and Q125T’ were chosen to represent the T’
treatment.
Tips were extracted from samples ground with SiC, polished with 1µm alumina, and
vibratory polished with 0.06µm colloidal silica. Tips were sharpened using standard focused-ion
beam techniques in a plasma FIB [75–77]. APT was performed using a LEAP 3000 in laser
mode (wavelength = 532 nm) at 75K with 200 kHz 0.5 nJ pulses, and tips were reconstructed
using IVAS software (version 3.8). Analysis of APT reconstructions included composition
measurements within selected volumes, compositional gradients within single-grain regions, and
composition gradients across interfaces. Phase identification was by correlating Ni content with
those determined through TEM EDS+MDP (e.g. regions above 10 wt.% Ni in tempered samples
are precipitated austenite, while regions below 10 wt.% are ferrite). Due to the difficulty with
phase identification with APT techniques, austenite stability was assumed to be what was seen in
TEM (i.e. TEM MDP results were not confirmed or refuted using APT). Importantly, APT
allowed the rest of the chemical makeup of austenite and ferrite, including C and trace elements,
to be determined. Secondarily, it also allowed study of carbides evolution as a function of heat
treatment, as carbides were too small to be characterized with TEM techniques. APT results were
52
complementarily contrasted with EDS results to create a complete picture of composition of
different phases as a function of heat treatment.
Figure 3.6: A pictorial summary of the experimental process utilized in this effort to characterize the undeformed microstructure
as a function of heat treatment.
3.3.5:Nano-Indentation
Nano-indentation experiments were performed on deformed QLT specimens to measure
hardness changes across the shear band. Indents were performed using a Hysitron nanoindenter
with a Berkovich tip and analyzed with Triboscan software; a maximum force of 5000µN (load-
controlled) was applied with a loading time of 250s, held at max load for 5s, and unloaded for
250s. Linear arrays of indentations spaced 2µm apart were obtained perpendicular to the band on
well-polished samples. Exact distance from the band for each indent location was calculated by
etching the sample to reveal the band, imaging the band using SEM, and physically measuring
the shortest distance from the indent to the center of the band.
53
Chapter 4: The Partitioning of Carbon During the Heat Treatment
of Quenched Fe-10Ni-0.1C Steel
4.1: Introduction
C and Ni are both austenite stabilizers and both play a role in affecting the thermal
stability of austenite in Fe-Ni-C steels subjected to quench-and-temper type heat treatments.
While the role of Ni on austenite stability has been examined to some extent [25,27,28], the
influence of C partitioning on austenite formation and stabilization in this alloy or the associated
kinetics have not been examined and may be important. Indeed, C is the primary austenite
stabilizing partitioning element in many TRIP steels [37,41,78–85].
The quantification of C in austenite and ferrite in this steel as a function of tempering is
complicated by the fact that C cannot be quantified using the TEM-EDS techniques that can be
used for Ni (due to the characteristic X-ray energies of C overlapping with the detector
background) and that there are two competing mechanisms for C during tempering- the
partitioning to austenite versus the formation and growth of refractory carbides such as MC and
M2C. Additionally, the solid-state transport mechanism for C is very different from Ni, as C is
interstitial while Ni is substitutional. Therefore, the influence of C on austenite formation and
thermal stability during tempering is likely temporally decoupled from the effect of Ni. Due to
the rapid diffusion of C in Fe relative to Ni, we focus on the redistribution of C from martensite
in the early stages of tempering and its likely influence on austenite precipitation rather than on
subsequent growth [81,85,86].
Quantification of C in the phases present (austenite, ferrite, and carbides) as a function of
heat treatment was performed using APT. The AQ, QL, Q25L, QLT, Q25T’, and Q125T’
54
specimens were all examined by APT, and the results from these specimens are presented in this
Chapter and their implications are discussed. (Lastly, the reader is reminded that in this alloy
system, Fe and Ni content numbers are similar when reported either in atomic percent or weight
percent whereas in the case of C, the atomic percent value is roughly five times that in weight
percent (see Table 3.1). The reason this is mentioned is because EDS data are in weight percent
whereas APT results are in atomic percent and C is only detected using APT and not by EDS).
4.2: Results: L Temper
The samples that best represent C partitioning as a function of tempering time are AQ,
QL, and Q25L. AQ is the original state or the reference state; QL includes the shortest temper
time of all heat treatments and is relevant not only because C transport is rapid at the L
temperature (650°C), but also because it is the first step of the multi-step QLT treatment. Q25L
provides a longer-term isothermal comparison to QL and is anticipated to be long enough that C
partitioning is effectively complete. Therefore, for the purposes of analyzing C behavior as a
function of tempering time, we consider AQ, QL, and Q25L treatments.
Bright field TEM imaging of the specimen following the AQ treatment confirmed a lath
martensite microstructure within which were dispersed a few fine second phase particles, ~20 nm
in diameter. These are highlighted in Figure 4.1a using dashed circles. The lath martensite
structure was verified by bright field TEM imaging coupled with microdiffraction, and EDS
analysis provided an average Ni content of 8.91.3 at.% which corresponds to the nominal alloy
composition in terms of Ni content. The highlighted second phase particles were too small to
reliably assess their chemistry by this technique and are believed to be undissolved carbides.
Microdiffraction confirms that some of these particles have a cubic structure and are thought to
be MC carbides that are known to be resistant to dissolution at high temperatures; MC carbides
55
and M2C carbides have been previously reported in such steels [22,27–29]. The second phase
particle morphology and chemistry were further characterized by APT. Figure 4.1b shows a part
of an APT volume after reconstruction: the levels of Mo, V, and Cr in the alloy are low (less than
1 at.% of each element) and therefore, iso-concentration surfaces with a threshold of 2-3 at.% of
these individual elements are sufficient to highlight the carbides in the matrix. Thus, a 2 at.% Mo
iso-concentration surface in Figure 4.1b shows the location of discrete Mo-rich regions believed
to be carbide particles. Composition measurements of these carbides confirm they are Mo-rich,
although they are alloyed with significant levels of V and some Cr. Some carbides are seen to
additionally have moderate levels of Ti (~10 at.%), although Ti was not specified as an alloying
element and therefore is thought to be an impurity.
56
Figure 4.1: a) TEM brightfield image of an austenitized and quenched (AQ) specimen showing a lath martensite microstructure;
some fine carbides (~20nm) are highlighted with dashed circles. b) APT composition measurement of a carbide shows
predominantly Mo but includes significant amounts of V and Ti. c) C segregation along interfaces in martensite.
57
Retained austenite was not immediately evident in the microstructure although it is
entirely possible that it is indeed present in small quantities as films along prior austenite grain
boundaries, between packets and blocks of martensite as well as between laths, as recently noted
[34,35,87]. In the martensite, APT permits measurement of C in solid solution and any C
segregation that might have occurred during the quench. Thus, the measured average C
composition in the martensitic matrix was 0.230.10 at.% whereas the bulk C level in this alloy
was 0.5 at.%; this implies that roughly half of the C in the alloy is tied up in the carbides that
remained undissolved during the austenitizing treatment. In addition, C segregation/enrichment
is observed in Figure 4.1c (highlighted using a dashed rectangular box) at what is likely an
interface (prior austenite grain boundaries and/or martensite packet/block/lath boundaries) with
peak values of around 2-2.5 at.%, which is an order of magnitude higher than the average C level
in the martensitic matrix. Similar levels of interfacial segregation of C was observed in multiple
locations.
The L tempering of a quenched 9.6 at.% Ni alloy (40 mins at 650°C - the QL treatment)
places it in the two-phase region of the Fe-Ni phase diagram, initiating austenite precipitation
and the transformation of martensite to ferrite. While the solubility of C in ferrite is negligible,
the precipitating austenite can serve as a reservoir for C partitioning out of the martensite
together with additional carbide precipitation. Bright field TEM imaging coupled with
microdiffraction (Figure 4.2a) and chemical analysis by EDS confirmed noticeable Ni
partitioning during tempering into relatively Ni-poor (7.1-9.5 at.%) ferrite and Ni-rich austenite
(16-17 at.%). This new generation of austenite particles is small (50-250nm) and Ni-rich, both of
which contribute to its thermal stability after a water quench [28]. Synchrotron measurements by
Jain et al. [28] estimated the austenite volume fraction to be 8.1% following this heat treatment.
58
The measured Ni level in ferrite following this L temper is higher than what the phase diagram
predicts (~5 at.%), implying that the ferrite is still supersaturated in Ni due to the relatively short
tempering time and the slow diffusion of Ni.
Figure 4.2: a) TEM EDS/MDP shows Ni-depleted ferrite of about 8 wt.% Ni and Ni-rich (15-17 wt.%), thermally stable
austenite precipitates. b) APT composition measurement of a carbide particle in the QL specimen. A Ti- and V-rich core is
surrounded by a Mo-rich but Ti- and V-depleted outer shell (delineated by dashed lines).
APT confirms the ferrite phase regions far from, as well as immediately adjacent to the
austenite interface to be almost completely depleted of C (0.010.01 at.% and 0.040.01 at.%,
59
respectively). The austenite, however, has a C level of 0.290.06 at.% - the solubility of C in
austenite is quite high, and the C is adequately mobile to migrate to it. Unlike the AQ sample, no
C-enriched planes were observed in the APT specimens in this condition; however, there are
many carbides observed in both austenite and ferrite and at the interphase interfaces. Some of the
carbides observed have a core composition that is different from the shell that forms around the
core; in the example shown in Figure 4.2b, a Ti-rich carbide is encased in a Mo-rich outer layer.
There is a noticeable decrease in Mo content in both the ferrite and austenite phases versus the
corresponding value measured in the martensite in the AQ condition (0.260.06 at.% and
0.260.06 at.% versus 0.520.10 at.%), implying that some of the Mo has being depleted from
solid solution to form new carbides.
The Ni distribution profile in the Q25L specimen measured by EDS is also bimodal with
Ni-poor regions (4.80.6 at.%) confirmed by microdiffraction to have a bcc structure (ferrite)
and relatively Ni-rich regions (12.81.5 at.%) that also yielded bcc microdiffraction patterns
(Figure 4.2a), indicative of freshly formed martensite following the water quench from 650°C.
Evidently, the Ni level in these Ni-rich regions is inadequate to stabilize the precipitated
austenite. The somewhat lower Ni level in the Q25L fresh martensite (that was austenite prior to
the quench) relative to the stable austenite in the QL treatment is thought to be related to the fine
austenite particle size in the latter and the related interface curvature-enhanced Ni solubility.
60
Figure 4.3: a) TEM EDS/MDP show the austenite precipitates are moderately Ni-rich (15 wt.%) while the ferrite is close to
equilibrium value of 5 wt.%. The precipitates are thermally unstable. b) Interface segregation of C in fresh martensite in the
Q25L specimen. c) APT measurement of the composition across a carbide particle shows a V-rich, Mo-poor carbide is
sandwiched by Mo-rich, V-depleted carbide.
APT shows that after 25 hours of heat treatment, C is depleted in ferrite and is negligibly
low in the fresh martensite (0.040.01 at.%) as well. Nevertheless, within this fresh martensite,
C-enriched regions (0.30.01 at.%) were present, suggesting some C-segregation to interfaces
(highlighted by the rectangular dashed box in Figure 4.3b), although considerably lower in
concentration than that observed in the AQ condition in Figure 4.1c. The freshly formed
martensite and the ferrite have Mo levels (0.330.06 at.% and 0.330.06 at.%) that are similar to
those observed in QL (0.260.06 at.% and 0.260.06 at.%), suggesting that the carbides have
completely precipitated out during the first 40 minutes of tempering. This would imply that the
61
decrease in the average C level in the freshly formed martensite in the Q25L specimen is
primarily a dilution effect resulting from the increase in the precipitated austenite fraction due to
the long tempering time of 25h. Compositionally, the carbides present in Q25L are similar to
those seen in QL: primarily Mo-based, with some showing a ‘core’ of V or Ti. In Figure 4.3c, a
large V-carbide (with almost no Mo) surrounded by patches of Mo-rich carbide is observed. In
this core region where the V level is ~50 at.%, the C level is around 33 at.%. Carbon is known to
be challenging to quantify by APT [88–91]. The level we report are likely an underestimation of
the actual C level in the carbide as such a V-rich carbide is most likely an MC carbide (V4C3)
rather than an M2C carbide as suggested by the chemistry. Thus, APT-based chemistry alone is
not necessarily adequate to distinguish between MC and M2C carbides as both are expected to be
present. The outer patches of the Mo-rich carbide are more likely M2C carbide type that
precipitates during tempering.
4.3: Results: T’ Temper and QLT
Table 4.1: C content in austenite and ferrite in Q25T’, Q125T’, and QLT as measured by APT (at.%)
Q25T' Q125T' QLT
Ferrite 0.01% 0.00% 0.00%
Austenite 0.01% 0.00% 0.05%
The C content in the austenite and ferrite in Q25T’, Q125T’, and QLT as measured by
APT are shown in Table 4.1. Of the three treatments, QLT has the shortest total tempering time
of 1 hour and 40 minutes, and while some C remains in solid solution in the austenite in this
state, it is significantly lower than that observed in QL austenite (0.290.06 at.%). While one
reason for this decrease is a dilution effect from additional austenite growth during the T temper,
some C is also likely stripped out of the austenite to form additional carbides. Thus, the one hour
at the T temperature of 590oC is adequate to remove all C from austenite and therefore, we
62
conclude that C does not contribute to the austenite stability when quenched from the T
temperature to room temperature. The Q25T’ and Q125T’ samples had much longer tempering
times, and no measurable C remains in austenite.
4.4: Discussion
We now discuss and summarize these findings in the context of the role of C in
influencing the stability of the precipitated austenite. Following water quenching from 800°C,
roughly half the total C in the alloy remains tied up in undissolved carbides while the rest is
present in martensite. The trapped C in martensite shows a tendency to segregate at interfaces
where its local concentration is almost ten times as high as within the martensite laths. Upon
tempering for 40 minutes at 650°C, fine austenite particles precipitate at lath interfaces while
Mo-rich carbides precipitate on undissolved carbides and at lath interfaces. In addition, the
austenite particles are Ni-rich, the Ni content in them being enhanced by particle size and
interface curvature. In contrast, the adjoining ferrite (martensite gradually decomposes to ferrite)
is C-depleted but still Ni-supersaturated relative to phase diagram prediction. The C content in
the austenite, the high Ni content, and the fine austenite particle size together stabilize the
austenite at room temperature following water quenching from 650°C. It appears that 40 minutes
at 650°C is adequate to fully redistribute the previously trapped C in martensite. When the
tempering time is increased to 25 hours, austenite fraction increases primarily by the partitioning
of Ni between ferrite and austenite while the C level in this austenite decreases substantially as
no additional C is available. Meanwhile, the average Ni level in the austenite is not as high, as
the interlath austenite size increases and curvature decreases. The austenite particle size increases
as well. Together, these factors reduce the austenite stability and make it more susceptible to
63
fresh martensite transformation on quenching following the 25 hours exposure to the tempering
temperature.
From the foregoing, it is evident that C does not play a major role in influencing the
thermal stability of precipitated austenite in longer tempering treatments. However, the migration
of C to the cores of dislocations that compose the lath boundaries during quenching and
enrichment of the austenite in early stages of tempering suggest that C could be important for the
nucleation and early stage growth of precipitated austenite during tempering. This early-stage
austenite precipitation can be austenite nucleation, or growth of nano-scale films of residual
austenite (which while we did not see but has been noted by other researchers on similar
materials [34,35,87]).
Thus, it is possible to envision a multi-stage austenite precipitation process from a
composition perspective. An initial stage at short tempering times (say 5 minutes to 30 minutes
at the tempering temperature of the QL material) where C diffusion stabilizes very fine austenite
particles while Ni diffusion lags and the Ni content in the austenite is less than what is predicted
by the phase diagram. In this early stage, as Ni diffusion in austenite is significantly slower than
in martensite, internal Ni concentration gradients may be present within the austenite [42]. A
second mixed stage is possible where Ni and C both stabilize austenite but the role of C
continually decreases as all available C is consumed in the first stage in the austenite and by
precipitation of refractory carbides. In this stage, austenite growth continues due to Ni diffusion
into the austenite, but the increasing volume fraction of austenite leads to C dilution in it. Beyond
this stage, Ni partitioning is all that matters.
64
Chapter 5: The Partitioning of Ni During the Heat Treatment of
Quenched Fe-10Ni-0.1C Steel
5.1: Introduction
In the previous Chapter, it was shown that while C partitioning may be relevant in the
very early stages of the L temper in influencing austenite nucleation, its role diminishes rapidly
within the first one hour at that temperature. In this Chapter, we similarly examine the
partitioning of Ni during the L and T tempers in detail.
The L-temper precipitated austenite is not expected to be particularly Ni-rich according to
the binary Fe-Ni phase diagram (Figure 2.3b) and thus there remains the question whether this
austenite is thermally stable upon quenching after this temper. Our observations in the previous
Chapter showed that austenite in the QL condition is in fact stable whereas that is not the case in
the Q25L condition suggesting that size may play an effect in the thermal stability of austenite.
5.2: The As-quenched Microstructure (AQ)
TEM brightfield images of the AQ treatment reveal a predominantly lath martensite
structure (as also mentioned in the previous Chapter), and selected area diffraction and MDP
confirm that it is indeed martensite (Figure 5.1). Arrays of EDS measurements show a narrow Ni
distribution of 9.3±1.2 wt.% Ni, which is close to the bulk composition of the alloy as did APT
results (10.6±1.9 wt.%Ni). It is possible that there are nano-sized films of retained austenite at
the lath boundary [34–36], but they are not readily observed in BF mode.
65
Figure 5.1: TEM brightfield image with SADP of the AQ treatment. Note the lath martensite structure.
5.3:The L Tempers Microstructure (QL and Q25L)
The L-tempering temperature of 650oC places the 10 Ni alloy in the two-phase region of
the Fe-Ni binary phase diagram implying the decomposition of martensite into moderately Ni-
rich austenite and a Ni-poor ferrite phase (refer to Figure 2.3b). An SEM image of a QL
specimen (Figure 5.2a) shows that 40 minutes at 650°C is long enough to enable widespread
precipitation of austenite, though it is much less than the ~0.5 volume fraction predicted by the
equilibrium Fe-Ni phase diagram. After 25 hours of tempering at 650oC, the austenite
precipitates have grown substantially and the small austenite particles have become micron-long
lamellae that cover just about half of the micrograph, consistent with phase diagram predictions
(Figure 5.2a). The austenite morphology in Figure 5.2a after a 25h exposure to 650oC suggests
austenite precipitation along interlath interfaces, and possibly along packet and block boundaries.
66
Figure 5.2: a) SEM images of QL and Q25L. After only 40 minutes of tempering, austenite has precipitated out.; after 25 hours,
the precipitates have evolved and occupy more than 50% of the total area. b) TEM-EDS arrays start around bulk composition (10
wt.%) for short tempering, but a bi-modal distribution is visible after 25 hours.
Arrays of EDS measurements were obtained in the TEM for the QL and Q25L specimens
and the results are compared in Figure 5.2b. In the QL condition, a Ni content distribution in the
microstructure with a peak centered at 9.3±0.9 wt.% Ni is noted although Ni levels as high as
17% and as low as 6.5% were observed. The peak is believed to be ferrite whose composition for
the most part is close to the Ni content of the alloy rather than the ferrite composition of ~5% Ni
predicted by the phase diagram. Indeed, EDS coupled with microdiffraction measurements show
that the ferrite for the most part is still supersaturated in Ni and has a composition ranging from
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7.5-10 wt.% Ni. Locations exhibiting high Ni content are associated with precipitated austenite
and discussed further below. The observed response clearly reflects the sluggish diffusion of
nickel at this temperature. In contrast, in the Q25L specimen, TEM/EDS measurements show
two distinct peaks in Ni content distribution in the microstructure: a ferrite peak at 5.1±0.6 wt.%
Ni and a wider austenite peak centered at 13.3±1.5 wt.% Ni, reflecting the evolution of the
tempering process and the gradual partitioning of Ni between ferrite and austenite.
Returning to the QL specimen, a TEM BF image (Figure 5.3a) shows that the precipitates
are small, around 100-200nm in size, and form at the boundaries between martensite laths. MDP
shows that the precipitate is thermally stable. EDS analysis of the austenite particle shown in
Figure 5.3a is presented in Figure 5.3b confirming its composition. The austenite precipitates
show a moderately high-Ni core between 15-18 wt.%, consistent with the austenite composition
expected from the phase diagram at this temperature (Figure 2.3b). The Ni content in the
austenite however appears to decrease gradually across the austenite/ferrite interface rather than
abruptly (Figure 5.3b), but this is attributed to an overlap through the foil thickness of the two
co-existing phases, one Ni-rich and one Ni-poor; the highest values in the center of the
precipitate are believed to be accurate.
68
Figure 5.3: a) TEM BF/MDP with corresponding EDS measurements (b) of the QL treatment. Note that the center of the
precipitate is Ni-rich, around 17 wt.% Ni, and radially decreases in composition. This is an artifact of overlap between the
austenite and ferrite, which causes a dilution effect where the austenite is thinnest. This is confirmed from interfacial composition
measurements using APT (c), where the composition measured near the phase boundary matches the composition measured by
EDS at the center of the austenite particles.
69
APT measurements (Figure 5.3c) confirm that the variation seen in TEM is indeed
artifact from overlapping phases: within nanometers of the boundary, the austenite composition
is measured to be 17.5±3.0 wt.% Ni, consistent with the values measured in the center of the
austenite precipitates using EDS. Adjacent to the austenite/ferrite interface, the Ni content in the
ferrite is somewhat higher (~6.5% Ni) than that predicted by the phase diagram (~5% Ni)
although it is higher in the bulk ferrite as previously noted.
In summary, APT and EDS together determine that QL produces small 100-200 nm
austenite precipitates with Ni content of about 15-17 wt.%, though the ferrite is still
supersaturated in Ni; this is consistent with the notion that the volume fraction of austenite is less
than that predicted by the phase diagram and the tempering kinetics are limited by Ni diffusion.
MDP show that these precipitates are thermally stable austenite.
After 25 hours of tempering at 650oC, the austenite precipitates have grown substantially,
predominantly along the martensite lath interfaces (Figure 5.2a). Ferrite and austenite were
distinguished using MDP and their Ni content was measured by EDS (Figure 5.4a and b).
Results show that the ferrite is uniformly depleted in Ni to the predicted equilibrium value of 5
wt.% Ni, and this is matched well by APT measurements (4.8±1.0 wt.% Ni). However, the Ni
content as measured by EDS within the austenite varied quite substantially, between 12 and 17
wt.% (Figure 5.4b). Unlike the EDS measurements of QL precipitates, the highest values
measured do not appear at the center of the precipitates, but at rather the ends and edges of the
austenite lamellae, and thus, this variation cannot entirely be attributed to overlapping phases.
APT (Figure 5.4c) measurements on the other hand fall in the lower end of the range obtained by
EDS (12.0±2.2 wt.% Ni) and showed no measurable composition gradient. It is worth noting that
EDS measurements were carried out along the length of the lamellar austenite precipitates that
70
decorate the prior martensite lath interfaces and over a distance of ~1 m, whereas the FIB lift-
out APT specimen likely traversed the austenite lamella and the measurement is typically only
over 100-200 nm. The origin of Ni content fluctuation in austenite in the EDS measurements is
not clear. As we know that in even in the early stages such as the QL treatment (40 minutes at
650oC), it is possible to get fine austenite particles with 17% Ni, and thus Ni pile-up at the
advancing austenite/ ferrite interface (due to reduced diffusivity of Ni in austenite as compared to
ferrite/martensite) appears to be an inadequate explanation. Combined, EDS indicates that there
are pockets of higher Ni similar to QL composition (17 wt.%), while APT and EDS show that
the composition is in fact not uniform and drops as low as 12 wt.% in many regions. Importantly,
MDP along these lamellae confirm that they are thermally unstable and have transformed to
martensite during the quench following this tempering step.
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Figure 5.4: a) TEM BF/MDP with corresponding EDS (b) of Q25L shows Ni content varies throughout, with some pockets of
~17 wt.% Ni and large swaths of 12-15 wt.% Ni. Some of the moderate Ni values are believed to be true, as they are from the
thickest parts of the austenite, and thus are unlikely to be an average of a higher Ni content averaged with low-Ni ferrite. b) APT
confirms that some of the low-Ni readings are not an artifact from overlapping phases, as the boundaries can be as low as 12
wt.% Ni. Together, we can conclude that some regions within Q25L austenite are quite Ni-rich, while others are very Ni-lean.
MDP shows the austenite of Q25L is thermally unstable.
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5.3: The T’ Temper Microstructures (Q5T’, Q25T’, Q125T’ and Q336T’)
The T’ tempering temperature is 110°C lower than L (650oC), and thus results in
significantly slower Ni diffusion. After 5 hours of tempering, the precipitates seen by SEM are
smaller and sparser than just 40 minutes of L (compare Figure 5.5a with Figure 5.2a); only after
125-336 hours of tempering do we begin to see the austenite take a coarser, lamellae-like
morphology. From the Fe-Ni phase diagram (Figure 2.3b), we expect the T’ to have ferrite
composition of ~6 wt.% Ni and a significantly richer Ni content of ~25 wt.% in austenite. TEM
EDS arrays show the slow depletion of Ni in ferrite, starting from 9.5±1.3 wt.% Ni after 5 hours
(Q5T’), going to 8.7±1.1 wt.% Ni in Q25T’, and then to 6.9±0.7 wt.% in Q125T’, and finally to
6.3±0.6 wt.% Ni in Q336T’ (Figure 5.5b); this progression is illustrative of the sluggish Ni
diffusion kinetics and the final value is in excellent agreement with the Fe-Ni phase diagram
(Figure 5.5b).
73
Figure 5.5: a) SEM comparison of T’ treatments. Ni diffusion at the T’ temperature is much slower than at L, and so even after 5
hours there is not a significant volume of precipitates. After 125 hours, precipitates now line most lath/packet/block/grain
boundaries, and after 336 hours, coarsened globules are seen along high-angle boundaries. b) EDS measurements of the Ni
content in ferrite show a progressive shifts towards its equilibrium value with longer tempering times
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The austenite precipitates in Q5T’ were examined in TEM BF mode and measured
between 50-100nm in diameter; after 25 hours, they had grown to around 200nm in diameter;
after 125 hours, they begin to show directional growth along interfaces and approached lengths
of 0.5µm. Q336T’ included two austenite morphologies: lamellae of the order of 1µm along lath
boundaries and large, equiaxed precipitates with radius between 0.5-2µm at high-angle
boundaries.
Austenite composition measurements using EDS on the Q5T’ specimen were not reliable
due to the fineness of the precipitate size, and hence the ratio of precipitate size to TEM foil
thickness. While some measurements were around 22 wt.% Ni, the Ni content often varied from
10-15 wt.%, likely due to overlap of ferrite and austenite in the foil thickness. As the precipitates
grew in size through further tempering (Figure 5.6a and d), EDS data became more reliable, and
a consistent austenite composition of 24-26 wt.% Ni emerged as the ceiling for Ni content in the
core (Figure 5.6b and e). As in QL, these precipitates showed a Ni gradient (radially for equiaxed
precipitates, lengthwise for high aspect-ratio precipitates), due to an overlap with the lesser Ni-
containing ferrite and hence an averaging effect. In contrast, APT confirmed that these austenite
precipitates were in fact uniformly high in Ni, with austenite compositions near the interface of
24.1±3.9 wt.% in Q25T’ and 24.8±4.0 wt.% Ni in Q125T’, both closely matching the EDS
values obtained from the core of the austenite precipitates (Figure 5.6c and f). We therefore
conclude that Q25T’ and Q125T’ precipitates have uniform Ni content between 24-26 wt.% Ni;
it is reasonable to assume that the precipitates in Q5T’ are similarly composed. MDP of these
relatively small and Ni-rich T' precipitates confirmed them to possess an FCC structure (Figure
5.6a and d), implying that the austenite particles were thermally stable and did not transform to
martensite on cooling.
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Figure 5.6: a) TEM BF with corresponding EDS (b) of Q25T’ shows the core of precipitates have a composition between 24-26
wt.% Ni with radially decreasing values. As in QL, APT (c) confirms that this decrease is an artifact of phases overlapping
through the foil thickness – APT composition measurements near austenite/ferrite interface closely matches those measured with
EDS at the austenite core. As precipitates grow with increasing tempering time (Q125T’, d), the overlap effect in EDS diminishes
while the core value remains constant (e). APT again corroborates EDS measurements(f). All precipitates in Q5T’-125T’ were
found to be thermally stable (FCC).
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Figure 5.7: Q336T’ has large, coarsened, equiaxed austenite precipitates on the scale of 1-2µm. They have a leaner composition
than smaller T’ treatment precipitates (20-22 wt.% Ni vs 24-26 wt.%), and are sometimes thermally stable (a, c) or have
transformed to martensite during the quench (b, d). Note the thermally stable austenite (c) has a low defect structure, while the
thermally unstable austenite (d) has a lath-like internal structure.
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In the Q336T’ condition, small, thermally stable austenite particles with 24-26 wt.% Ni
are also observed. However, the lengthy tempering time also allows some precipitates to
significantly grow and coarsen; equiaxed austenite precipitates with diameter up to 2µm can be
observed (Figure 5.7). In these relatively coarse precipitates, overlap effect if present is
negligible even within tens of nanometers of the interface, giving reliable, uniformly high EDS
measurements throughout of 20-22 wt.% Ni. Some of these globular austenite precipitates are
thermally stable, and show a very low-defect structure within them; however, other austenite
precipitates with a similar Ni content are unstable (MDP confirms bcc structure), with internal
lath-like sub-structures and a high dislocation density.
Table 5.1: Composition of austenite in the four isothermal heat treatments as measured by APT (at.%)
AQ QL Q25L Q25T' Q125T' Avg. Err.
Fe 87.1% 79.5% 86.3% 72.6% 73.7% 3.7%
Ni 10.1% 16.7% 11.5% 23.7% 23.0% 2.8%
Mn 0.67% 1.77% 0.92% 2.46% 2.06% 0.30%
Cr 0.66% 0.79% 0.63% 0.60% 0.66% 0.14%
Mo 0.52% 0.26% 0.33% 0.30% 0.28% 0.07%
Si 0.48% 0.29% 0.15% 0.19% 0.12% 0.06%
C 0.26% 0.29% 0.05% 0.01% 0.00% 0.04%
Cu 0.17% 0.29% 0.07% 0.18% 0.04% 0.03%
V 0.07% 0.03% 0.05% 0.02% 0.04% 0.01%
Al 0.02% 0.01% 0.03% 0.02% 0.06% 0.01%
In addition to Ni, it is relevant to consider the possible role of other minor alloying
elements present in the alloy in affecting the thermal stability of austenite. The composition of
austenite measured in several heat treatments using APT is summarized in Table 5.1. However,
while many of these elements also partition during tempering (e.g. Mn), the difference in their
content in austenite between the L and T’ tempering treatments is not as significant as the
78
partitioning of Ni. As discussed in the previous chapter, C is notably present in QL austenite;
however, it is still relatively low (0.3 at.%, or 0.05 wt.%) and is therefore not as significant in
relation to Ni partitioning; no other treatment has measurable C due to carbide precipitation and
growth in the early stages of tempering.
5.4: Discussion of Isothermal Tempering
Table 5.2: Size and composition of precipitates in isothermal treatments
Treatment Time (h) Morphology Size (µm) Austenite Comp
(wt.% Ni) Stability
L (650°C) 0.67 Equiaxed 0.1 17 Stable
25 Lamellar ~1 12-17 Unstable
T' (540°C)
5 Equiaxed 0.05 10-22* Stable
25 Equiaxed 0.25 25 Stable
125 Lamellar 0.5 25 Stable
336 Lamellar ~1 25 Stable
Equiaxed 0.5-2 22 Mixed
* Q5T’ austenite composition measurements are believed to be unreliable due to their size, true composition is assumed to be
25 wt.% Ni.
A summary of the austenite precipitate shape, size, composition in terms of Ni content,
and stability following quenching from the different isothermal tempering conditions examined
in this chapter are presented in Table 5.2. For both tempering temperatures (L and T’), short
tempers produce small, equiaxed, thermally stable austenite with Ni content that closely matches
the Fe-Ni phase diagram. However, as the austenite grows and coarsens during extended
tempering, both isothermal treatments produce thermally unstable austenite. There are two
possible explanations for this onset of instability: their relative decrease in Ni content and
increase in size.
Larger austenite precipitates in both treatments appear slightly leaner in Ni than the
precipitates seen in shorter treatments, which should lessen thermal stability. The austenite in
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Q25L is on average roughly 3 wt.% Ni poorer than QL, although it is highly inhomogeneous; the
homogenous, globular Q336T’ precipitates are also about 3 wt.% poorer than lamellar or smaller
equiaxed T’ austenite. While this slight decrease in Ni undoubtedly decreases the thermal
stability of the austenite precipitates relative to the richer, short-temper ones, the stability of
small, moderate-Ni QL austenite juxtaposed with the instability of globular, Ni-rich Q336T’
austenite suggests that size is likely the dominant factor in thermal stability.
Takaki et al. showed that there is a strong Ms temperature suppression for sub-micron-
sized austenite, as an increasingly significant chemical driving force is required to overcome the
relatively massive elastic energy barrier for transformation [38]. While Q336T’ is significantly
more Ni rich than Q25L, both have micron-scale dimensions and are thermally unstable; despite
being much more Ni-lean than unstable Q336T’ austenite, the 100nm QL austenite precipitates
are thermally stable. Therefore, we can conclude then that within the range of L and T’
tempering temperatures, that is to say, between 15-25 wt.% Ni, austenite precipitates are only
thermally stable if sufficiently small. Moreover, if allowed to coarsen to micron-scale sizes
through extended tempering, austenite in all isothermal treatments within this temperature range
should transform to martensite upon a room temperature quench. This size effect is schematically
superimposed on the Ms temperature versus Ni plot in Figure 5.8, while Figure 5.9 summarizes
the above discussion relating to size and Ni content on the stability of the austenite precipitates.
80
Figure 5.8: The Ms temperature diagram for a binary Fe-Ni alloy [43] predicts both L and T’ austenite to be thermally unstable
at room temperature. The observed austenite stability in short tempering times is thought to be due to a size-effect.
Figure 5.9: For both L and T’ tempers, the small austenite from short tempers is thermally stable. However, even the very Ni-
rich T’ austenite becomes unstable after 336 hours, suggesting both small L and T’ tempers are only stable due to their small size.
If let coarsen long enough, presumably all T’ austenite should be unstable.
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However, effect of composition on thermal stability may still be relevant in certain
circumstances. That is, a sufficiently large, moderate-Ni L austenite may be thermally unstable
while a similarly sized, Ni-rich T’ austenite would not be, as the chemical driving force for the
transformation would be higher for L than T’. In principle then, lower temperature tempering in
this alloy system should be able to produce thermally stable austenite of a larger size than high
temperature treatments due to a composition-enhanced stability.
There are other considerations that go into selecting a tempering schedule, even within a
range capable of producing thermally stable austenite of a desired size. The kinetics of austenite
nucleation and growth are substantially different between high and low temperature tempers as
evidenced by comparing QL with Q25T’. The Q25T’ needed to be tempered roughly 35 times
longer than QL to produce austenite precipitates of similar sizes. In addition, there appear to be
morphological differences between L-austenite and T’-austenite, the former appearing more
lamellar, and forming at inter-lath interfaces whereas T’ austenite appears globular. As shown by
Kim et al. [13], a lamellar austenite morphology is more effective at breaking up trans-packet
cleavage. But beyond these considerations, desired mechanical instability is another factor that
can affect selection of tempering temperature. TRIP mechanisms are also a function of
composition, size, and morphology, and so choosing a composition that produces austenite that
can be readily rendered mechanically unstable may be desirable. So while precipitated austenite
from a wide range of short tempering temperatures would be stable due to its small size, there are
additional considerations that influence selecting a tempering schedule.
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5.5: The QLT Treatment – Results and Discussion
Figure 5.10: SEM of QL (shown here again for convenience) and QLT. Note that there is significantly more austenite than QL,
which suggests that the T temper contributes heavily to additional austenite nucleation and growth.
The microstructure of the multistep QLT temper is the result of a summation of the
microstructure resulting from the QL temper and a temper in between T’ and L; by comparing
QLT to AQ, QL, and the QT’ samples, we can isolate the effect of each tempering step on the
final microstructure. SEM of QLT (Figure 5.10) shows noticeably more precipitation than QL,
which suggests that additional austenite continues to form during the T treatment. The Fe-Ni
phase diagram (Figure 2.3b) predicts this second generation austenite to have a higher Ni content
than austenite formed during L; however, it is unclear if the L-generation of austenite that was
thermally stable will have adequate time to acquire Ni to attain the new equilibrium value at the
T temperature, as Ni diffusion within austenite is slower than Ni diffusion to the austenite/ferrite
interface through ferrite. Brightfield images show a mixture of lamellar and equiaxed
precipitates; lamellar austenite can have a length as much as 1µm, though narrower than Q25L
and Q336T’ lamella (100nm), while equiaxed precipitates are roughly 100nm in diameter
(Figure 5.11a). TEM EDS (Figure 5.11b) shows a scattered range of Ni level between 15-20
wt.%, even at the core of the austenite. Like in QL and T’ treatments, this may be simply due to
83
the Ni-rich austenite and Ni-poor ferrite phases overlapping through the foil thickness, but it also
may be attributed in part to two distinct generations of austenite with different Ni content. APT
confirms that some of this variation is in fact real, as austenite as low as 14.8±2.6 wt.% Ni and as
high as 22.1±3.6 wt.% Ni is present (Figure 5.11c). MDP of QLT precipitates confirm they are
thermally stable.
Figure 5.11: a) TEM BF of QLT with corresponding EDS (b) shows mixed composition of 15-20 wt.% at the center of
precipitates. This is not an artifact of phase overlap, but rather is a result of two distinct generations of austenite with different
growth compositions. c) APT tips confirm that both L-generation and T-generation austenite are present in QLT.
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SEM images corroborate the austenite volume increases seen by Zhang [25] and by Jain
et al. [28]. The T treatment measurably increases the volume fraction of austenite precipitates.
However, the reason for this increase appears to be somewhat different from the theory they
proposed; they argue that many of the precipitates seen in the SEM image following the QL
temper are in fact fresh martensite, and thus the thermally stable austenite content is less than
what it appears to be visually in such images. They postulate that much of the stable austenite
found in QLT is thus a consequence of precipitation of a new generation of austenite within these
Ni-rich fresh martensite pockets. However, as discussed previously, we did not find evidence for
unstable austenite in QL and we attributed the stability to size. Therefore, our view is that the
increase of stable austenite content from the second lower temperature T temper is due to
additional precipitation on top of existing L-generation austenite as well as nucleation and
growth. This two-step austenite evolution is evidenced by a combination of EDS and APT,
which show two compositionally distinct generations of thermally stable austenite; in the case of
co-precipitation, because the T temper is so short, the residual L-generation austenite does not
have time to equilibrate to the new austenite composition. Thus, we can separate the L- and T-
generation of austenite precipitation through Ni content; 15-17 wt.% L-generation austenite and
20-22 wt.% T-generation austenite. Schematically, this two-stage austenite core-and-shell
morphology evolution is shown in Figure 5.12. In addition, we can now provide a more complete
schematic overview of the QLT process compared to that presented in Chapter 2 (Figure 2.4),
complete with austenite composition and identification of phases for each stage of tempering
(Figure 5.13).
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Figure 5.12: QLT process produces thermally stable austenite around 15-17 wt.% Ni during the L treatment. During T, growth
resumes, resulting in additional growth in the range of 20-22 wt.% Ni.
Figure 5.13: QLT Process, understood in terms of austenite size and composition.
86
The thermal stability of both generations of QLT austenite appears to be due to their
small size: QLT austenite is still in the sub-micron regime, and the composition range falls
within the boundaries of the isothermal L and T’ treatments- both of which were found to be
unstable when sufficiently large. While the Ni-enriching T-step does provide additional thermal
stability to some of the austenite, it is perhaps not the primary factor for an increase in thermal
stability of the austenite as a whole. Perhaps a longer L-step could also produce the same volume
fraction of thermally stable austenite with similar morphology.
So why does QLT perform better in mechanical tests if similar volume fractions can be
produced isothermally? Zhang [25] is not the first to describe the mechanical superiority of a
QLT treatment in Ni steels; Kim et al. [13] in 1983 found that QLT treatments on a 5.5Ni steel
also produced superior Charpy toughness versus a range of isothermal treatments. It seems that
there must be microstructural reasons that a two-stage temper has superior mechanical properties
during dynamic deformation, be it Charpy impact energy or ballistic resistance. In both studies,
QLT-treated Ni-containing alloy is believed to undergo a TRIP mechanism wherein the austenite
dynamically transforms to martensite, and like thermal stability, the mechanical stability of
austenite is also a function of austenite composition and size. Recently, Yuan et al. [41]
suggested that a distribution of austenite in the microstructure with varied composition can have
a ‘spectral TRIP effect’, where the critical strains to transform individual austenite particles to
martensite are then spread over a large range, thereby increasing the work hardening capacity
and toughness of the material. In contrast, homogeneously composed austenite from that same
alloy was found to transform within a narrow strain window and was thus prone to failure with
lower toughness. A similar effect was also noted by Wang et al. [39,40] with respect to varied
austenite size- a distribution of austenite grain sizes with the same composition also helped
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postpone failure and improved overall toughness through staggered critical strains. The QLT
treatment that the 10Ni steel has been subjected to may combine both of these mechanisms
through its two stage temper: L- and T-generation austenite are chemically distinct, providing a
composition-based spectral TRIP; while the mixed sizes and morphologies arising from two
stages of growth provide a size-based spectral TRIP. Thus, the ballistic superiority of 10Ni QLT
can be explained through a combination of spectral TRIP effects, which are enabled through its
two stages of thermally stable austenite growth that can at least in part be rendered mechanically
unstable.
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Chapter 6: Microstructural Evolution in an Fe-10Ni-0.1C Steel
During Dynamic Deformation
6.1: Introduction
It is now established that Fe-10Ni-0.1C steel subjected to the QLT heat treatment results
in a microstructure containing two compositionally distinct, thermally stable generations of
austenite precipitates: the 15-17 wt.% Ni L-generation and the 20-22 wt.% T-generation.
However, the question remains- how does the microstructure evolve during high strain-rate
deformation? Specifically, what happens to each generation of austenite as deformation
commences, progresses, and then localizes, initiating adiabatic shear bands (ASB) that then
propagate through the material? This chapter examines these aspects using the compression
Kolsky-bar set-up to obtain the high strain rates needed to study the problem.
6.2: Kolsky-Bar Calibration using 4140 Steel
The supply of Fe-10Ni-0.1C steel with the QLT heat treatment steel was limited, and
therefore, Kolsky-bar experiments were first performed on 4140 tempered martensite for
purposes of calibration. A wide range of firing pressures was examined for each projectile bar
length, from low pressure to a regime where specimen fractured consistently or the maximum
permitted chamber pressure was reached, whichever occurred first. Strain/strain-rate pairing was
recorded in each instance and plotted for different projectile bar lengths (Figure 6.1) so that the
test parameters for specific predetermined strain/strain-rate regime could be selected a priori for
the 10Ni QLT experiments. While the 15.3cm bar achieved the greatest strain-rates, the total
height strain was limited by the upper allowable firing pressure; alternatively, the 30.5cm bar
could achieve much higher height strains while still reaching relatively high (2000s-1
) strain-
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rates, and so it was selected for the QLT experiments. Etched 4140 smooth cylindrical samples
were observed using SEM, and the presence of ASBs were confirmed; thus, this sample
geometry was determined sufficient for producing ASB in 10Ni QLT experiments.
Figure 6.1: Strain/strain-rate pairs for Kolsky experiments on 4140 tempered martensite. Each datum point correlates to a
specific firing pressure with the noted bar.
6.3: Kolsky-Bar Testing of 10Ni-QLT
As the focus of this study was the microstructural evolution during shear localization, low
strain and strain-rate Kolsky-bar experiments were not performed on QLT because they were
considered unlikely to localize and produce ASBs. Instead, gun pressure was increased up to and
just short of fracture in order to attempt to consistently induce shear localization. A variety of
samples in various stages of shear localization were produced, and the extent of shear
localization in each was determined using optical microscopy (examples are shown in Figure
6.2). The appearance of the white bands classified them as ‘transformed bands,’ suggesting they
had undergone dynamic recrystallization in the regions of most intense localization
[51,52,57,61,68]. No cracks or voids were observed in the samples containing these white bands.
0
500
1000
1500
2000
2500
3000
0 0.05 0.1 0.15 0.2 0.25
Stra
in-r
ate
(s-1
)
Height Strain
15.3cm
20.3cm
30.5cm
40.6cm
90
Figure 6.2: A diffuse and a well-developed adiabatic shear band
One particular sample was chosen for further investigation because it had two well-
developed shear bands in opposite corners which dissipated towards the center of the sample
(Figure 6.3). Thus from this single sample, several TEM specimens were FIB-lifted-out from
specific locations that corresponded to various stages of shear localization; a location far from
the shear band was included as a baseline and represented what might be considered as a
homogeneous plastic deformation region prior to the onset of localization. The nominal global
strain-rate for this particular sample was around 1900s-1
and the total height strain about 20% (It
should be noted however that the strain rates and strains experienced after localization
commences, and particularly within the ASB, will be significantly higher).
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Figure 6.3: Sample chosen for microstructural analysis. A well-developed shear band initiates at the bottom left, propagates
towards the center of the sample (top right) and eventually dissipates. Another band initiated in the opposite corner (not visible in
this image) that also propagated towards the center of the sample, though it was smaller and less developed than the one seen in
the montage above. The total height strain was about 20% and the strain-rate was about 1900s-1. Specific TEM specimen lift-out
locations are marked using numbers from 1 to 10. The relative displacements of the vertical etching bands (banding in the rolled
plate) provide a sense of the large shear experienced within the ASB.
6.4: Nanoindentation across the ASB
An array of nanoindentation hardness measurements was made across the well-developed
ASB (Figure 6.4). There is a measurable increase in hardness across the band- from 5.5GPa to
7GPa, in agreement with the increase from 6.5GPa to 8GPa previously measured by Wang and
Kumar across a shear band in this alloy that had been subjected to a different heat treatment [22].
In addition, the region with increased hardness extends well beyond the boundaries of the white
portion of the band as was also observed by Xue and Gray in 316L stainless steel [58,69]. They
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noted that the white band did not signify the boundary of the shear localization but that the
microstructure in the white band has dynamically recrystallized, whereas the regions flanking it
was in the process of forming subgrains.
Figure 6.4: Nanoindentation across the shear band (marked in purple). There is an increase in hardness (~5GPa to ~7.5GPa) that
extends beyond the boundary of the unetched portion of the shear band.
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6.5: Microstructural Analysis of the dynamically deformed Specimen
As noted above, the strain localized region extends in width beyond the unetched portion
of the band, the extent of which can be seen in SEM micrographs (Figure 6.5). The center of the
band itself appears featureless and corresponds to the ‘white band’ observed by optical
microscopy; adjacent to this region, the extent of shearing is evident through the highly
elongated microstructure parallel to the ASB. About 10-15µm away from the center of the band
and on either side of it the microstructure is significantly less aligned. The unetched portion of
the band is approximately 4µm wide. Inside this central portion of the band, BF TEM
micrographs show that significant microstructural changes have occurred: no trace of the former
lath martensite structure remains, and is instead replaced by a fine-grained, equiaxed structure
(Figure 6.6).
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Figure 6.5: SEM micrographs of QLT at increasing magnification (top to bottom), undeformed (left) and near the shear band
(right). Note the lack of features in the band, which gives it the unetched appearance. Outside of this region, the austenite
precipitates can be seen to be sheared parallel to the shear plane.
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Figure 6.6: SADP in undeformed QLT compared to in the band using the same aperture. On the left, the lath martensite structure
produces a highly textured SADP; on the right, the highly misoriented, equiaxed, mechanically recrystallized grains in the band
produce a ring-like SADP.
To characterize the microstructural evolution during adiabatic shear banding, we consider
four locations in particular: i) far from the band in a homogeneously deformed region; ii) ahead
of the white band, in the band plane, where localization had occurred but was not well-developed
(labelled ‘7’ in Figure 6.3); iii) adjacent to the developed band, and an orientation that was
normal to the band plane (labelled ‘5’ in Figure 6.3); and iv) in the band, normal to the band
plane (labelled ‘10’ in Figure 6.3). As this deformation mode is extremely rapid, any associated
phase transformation is expected to be diffusionless, and thus the Ni content can be used to link
deformed microstructure to undeformed microstructure. Therefore, we particularly look to
characterize regions with Ni content that ties them to the two generations of austenite in
undeformed material, the 15-17 wt.% L austenite and the 20-22 wt.% T austenite.
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6.5.1: Location far from the shear band:
Figure 6.7: TEM lift-out from a location that is far from the shear band (circled in blue dashed line).
The ‘striations’ seen in Figure 6.7 that etch light and dark result from rolling and
originate likely from macrosegregation but conveniently serve as markers for demarcating
localized deformation. Far from the ASB, these striations remain relatively parallel, suggesting
there is not significant local shearing in this region as opposed to the region flanking the ASB
where significant shearing is evident (Figure 6.7). Indeed, the microstructure seen in TEM BF in
a specimen obtained from the location circled in Figure 6.7 appears very similar to the
undeformed material, and the lath microstructure is still intact (Figure 6.8). Notably, precipitates
with Ni content ranging from 15-20 wt.% (indicative of both, L-generation and T-generation
austenite) were found to retain the FCC structure implying that the L- and T-generation austenite
are mechanically stable in this location (Figure 6.8).
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Figure 6.8: Far from the shear localized region, the lath martensite seen in the undeformed microstructure is still intact. Here, the
composition of two austenite precipitates are measured and found to be between 15-20 wt.% Ni. Thus these precipitates are a
mixture of L- and T-generation austenite, and MDP shows that they are mechanically stable.
6.5.2:Location ahead of shear band:
The differential etching response manifesting as bright and dark striations in the optical
micrograph shows that there is intense shear just ahead of the ‘white band’ (Figure 6.9a); in the
TEM, BF imaging confirms that the martensitic laths and/or austenite precipitate morphology
noted in the undeformed material is no longer visible and the microstructure looks more
‘recrystallized’ (Figure 6.9b). SADP over large areas shows a more or less continuous ring
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pattern implying the structure consists of multiple grains/subgrains that are significantly
misoriented with respect to each other, suggesting a good level of subgrain rotation has occurred,
likely due to the intense shearing that prevails in this location. However, since this process can
be considered diffusionless, the relationship between these newly-developed substructures and
their parent should still be discernable through Ni content measurements.
Figure 6.9: a) TEM specimen lift-out from a location ahead of the shear band. Notice the stark white band on the left, which
dissipates to the right. b) SADP from the location ahead of the band microstructure (aperture marked in orange). Ring pattern
suggests significant grain refinement and grain rotation and recognized from the relatively low magnification bright field image.
The fine austenite precipitates in the undeformed QLT condition is expected to have a
low-defect structure, but these subgrains show dislocation contrast within them implying
significant plastic deformation had occurred and enough to break these precipitate down into
subgrains. Figure 6.10 shows one such related event: beyond what is in contrast initially in
Figure 6.10b, there is a high-Ni region that extends into a wider area that is not in contrast- this
continuous, unbroken high-Ni region represents the trace of a single former austenite precipitate.
Through tilting (Figure 6.10a), its neighboring subgrain (Figure 6.10c) can be brought into
contrast and similarly indexed. The first subgrain is entirely from T-generation precipitation, as
demonstrated by its homogeneous Ni content of ~22 wt.%; MDP shows it is mechanically stable
austenite. Similarly, its neighboring subgrain is also from T-generation growth (20-22 wt.%) and
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is mechanically stable. Together, they make up a T-generation austenite precipitate that has
broken up through mechanical recrystallization into mechanically stable austenite subgrains.
This implies that the T generation austenite is stable despite the intense shear experienced in this
region locally and that if it is going to mechanically transform to martensite, it is going to require
even more intense deformation (i.e. much ‘later’ in the dynamic deformation process). The
microstructure also shows that the austenite precipitates from the QLT treatment are further
segmented during the local intense shearing process that occurs ahead of the propagating ASB,
so that a reduction in austenite subgrain size prior to TRIP can further discourage the TRIP
process and may even preclude it.
Figure 6.11 shows another austenite precipitate dynamically broken down into subgrains.
Similar to the previous example, the T-generation subgrain (Figure 6.11c) is mechanically stable.
However, its neighbor is comprised of two generations of austenite growth- the 15-17 wt.% L-
generation and 20-22 wt.% T-generation (Figure 6.11b). This mixture suggests that the parent
precipitate originally had a core/shell structure resulting from T-generation precipitation on top
of an existing L-generation precipitate. The L-generation austenite should be more mechanically
unstable due to its lower Ni content. However, this subgrain is mechanically stable, possibly
stabilized by the fragmentation of the original grain that reinforces the size effect.
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Figure 6.10: a) Two adjacent subgrains, which are brought into contrast by tilting. b) EDS on the first subgrain gives a
composition of about 22 wt.% Ni (T-generation), and MDP shows it is mechanically stable FCC. c) Its neighbor is also from T-
generation austenite (20 wt.% Ni), and is also mechanically stable austenite. Together, they show a T-generation austenite
precipitate has mechanically broken down into subgrains and these are mechanically stable.
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Figure 6.11: An L- and T-generation austenite precipitate has mechanically broken down into two subgrains (b and c) and can be
brought into contrast through tilting (a). The subgrain before tilting (b) is mechanically stable and has Ni content varying between
L- and T-generation, and was likely part of a core/shell structure before deformation. Its neighboring subgrain (c) is also
mechanically stable, but its composition shows that it is entirely growth from the T temper.
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6.5.3: Location adjacent to the shear band:
Next, we examine the deformed microstructure in the regions adjacent to the shear band
(Figure 6.12a). The region normal to the shear plane about 20µm from the edge of the white
band has a microstructure that appears lamellar with lamellae aligned parallel to the plane of the
shear band (Figure 6.12b). This morphology is seen throughout the TEM lift-out specimen,
which was dimensions that are approximately 20µm x 30µm and is larger than the original
martensite packet and block size. Thus, this lamellar microstructure is not the original
martensitic lath structure, but rather a consequence of microstructural evolution during shear
localization- the martensite structure has broken down and has been sheared parallel to the ASB
propagation direction. EDS analysis confirms thin strands of Ni-rich regions aligned parallel to
these lamellae, and an example is illustrated in Figure 6.13, where the strand is composed of
subgrains, some of which are in contrast while others are not. The subgrain that is in contrast has
Ni content ranging from 15-22 wt.%, suggesting it is a mixture of both L- and T-generation
austenite. MDP confirms that the subgrain is mechanically stable austenite (Figure 6.13).
Figure 6.12: a) Lift-out location with ASB demarcated by the two orange dashed lines and axes defined (shear plane normal is X
axis, and shear band is propagating in the Z direction). b) TEM BF image of area about 20µm from the band with rotated axes
marked. Microstructure is elongated in the shear plane (Y-Z).
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Figure 6.13: Both, the ferrite and the austenite precipitates adjacent to the band are aligned parallel to the shear plane (axes
marked) and have broken up into subgrains. The subgrain that is in contrast here has Ni content ranging from 15-22 wt.%,
suggesting it originates from a mixture of L- and T-generation austenite. MDP shows the subgrain to have the FCC structure.
6.5.4: Location inside shear band:
TEM specimens that were initially lifted out from within the band but aligned parallel to
the band plane provided EDS Ni distributions that narrowly peaked around the nominal
composition of the alloy. This is a direct consequence of an averaging effect due to the grain’s
extremely small dimension along the axis normal to the shear plane (schematically illustrated in
Figure 6.14). Therefore, an in-band sample was milled normal to the shear band to increase
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single-grain-depth through the foil thickness for TEM analysis, as the grain dimension along the
band propagation axis was comparable to the sample thickness (Figure 6.15). To verify the lift-
out sample thickness was in fact single-grain, EDS Ni distribution plots were compared to those
from undeformed material and were found to be closely matched (Figure 6.16- compare green
and red curves as opposed to the blue curve which was obtained from a TEM lift-out oriented
parallel to the plane of the shear band). To ensure the area thinned to electron transparency was
indeed within the band, the location of the white band relative to the sample was carefully
measured before milling to approximate its location in TEM.
Figure 6.14: Exaggerated schematic of TEM specimen lift-out geometries. The austenite precipitates are stretched parallel to the
shear band- when lift-outs are taken along this plane, there is an averaging effect in the EDS measurements. Lift-outs are taken
normal to the band to mitigate this effect.
Figure 6.15: Left: Location of in-band lift-out prior to FIB milling. The unetched portion of the band is approximately 4µm from
the lift-out edge, while the other boundary of the band is about 10µm into the lift-out. Right: The measurements taken on the
SEM are overlaid on a BF image of the lift-out specimen (which extends beyond what is shown, as marked by the black arrows).
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Figure 6.16: Normalized EDS Ni distribution of undeformed QLT (Green) in comparison to data obtained from a TEM specimen
that was lifted out parallel to the band (Blue) and another that was oriented normal to the shear band (Red). The parallel geometry
has a strong, narrow peak around the nominal Ni content due to an averaging effect rising from the ratio of grain depth to sample
thickness; the normal-to-band sample has a distribution that closely matches the undeformed material, suggesting it largely has
single-grain depth.
Figure 6.17: BF and corresponding DF image of the region of the deformed specimen that is estimated to include the white band
and the region immediately outside it. There is no observable difference in microstructure, but note that the demarcation of the
boundary is a best estimate based on markings in Figure 6.15.
0
0.05
0.1
0.15
0.2
0.25
0 5 10 15 20
Co
un
ts (
No
rmal
ize
d)
wt.% Ni
Undeformed QLT
In-Band Liftout (Parallel)
In-Band Liftout (Normal)
106
As seen in Figure 6.17, the microstructure estimated to be inside the white band has
mechanically recrystallized into a fine-grain, equiaxed microstructure with grains diameters
between 50-200 nm. Darkfield images confirm the grains have adequately rotated to be
differentiated from their neighbors. The grain size is much smaller than the subgrains observed
ahead of the band (Figure 6.10 and Figure 6.11), demonstrating further grain refinement has
occurred within the band as the shear band developed. The microstructure is also distinctly
different from the regions adjacent to the band (further away from the region shown Figure 6.17
and illustrated in the next sub-section), as no elongated lamellar grain morphology is visible.
There is no sharp, delineating microstructural features that separates the portion of the sample in
the unetched (white) portion of the shear band (labelled ‘inside white band’ in Figure 6.17) from
the region immediately outside of it (labelled ‘outside white band’), although the boundary itself
is only located based on best estimate from the images shown in Figure 6.15.
The low-Ni (previously ferrite lath) grains are between 50-300 nm in diameter, and some
have features that appear to be BCC twins (Figure 6.18). The high-Ni, former austenite
precipitates also appear to be extensively deformed and then mechanically recrystallized with a
morphology that resembles 100-200 nm-long strips within the ferrite and composed of small,
misoriented, equiaxed grains of diameter 20-50nm lined single-file along the strip. Figure 6.19
shows one-such high-Ni region located in the area of the sample believed to be just outside the
white band; the high-Ni strip is narrow and elongated and is surrounded by low-Ni ferrite. The
portion of the strip that is in contrast in Figure 6.19 has a Ni content of about 22 wt.% Ni,
suggesting it is likely T-generation austenite. Its MDP gives an FCC pattern with additional
twinning spots – thus we conclude it is heavily deformed, mechanically stable, T-generation
austenite.
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Figure 6.18: A small grain is in contrast. It is low in Ni (boundaries marked in orange on EDS gradient plot, 5-7 wt.%Ni). Inside,
a band-like substructure suggests twinning.
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Figure 6.19: The high-Ni region, formerly an austenite precipitate, is spread over a very narrow strip (about 20nm wide by
200nm in length). The portion that is in contrast is marked in the plot in orange and has a Ni content of 20 wt.%, suggesting it is
T-generation austenite; an MDP from this location confirms it to be FCC and additionally reveals twinning- it is mechanically
stable austenite.
In other regions within the shear band, there is evidence for the austenite being
mechanically unstable. Figure 6.20 shows a narrow strand along which the Ni content was high
but is surrounded by low-Ni regions in the portion of the sample along the border of the
approximated white band. The in-contrast grain in the strand is about 20nm in diameter and has
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composition of around 20 wt.% Ni, suggesting it had evolved from T-generation austenite.
However, it has a BCC MDP implying it had mechanically transformed to martensite. Since the
size of this particle is of the order of 20nm, a plausible scenario is that the parent more-or-less
equiaxed austenite particle that was deformed into the long strand-like configuration transformed
to the BCC martensitic structure (TRIP) and then fragmented by mechanical recrystallization
into multiple fine subgrains, one of which is in contrast in Figure 6.20. Similarly, an L-
generation grain (15-17 wt.% Ni) from the white-band section of the lift-out sample which is in
contrast is shown in Figure 6.21. Given its relatively lower Ni content, it is should be less
mechanically stable than a T-generation particle; indeed, it has also mechanically transformed to
martensite.
Due to the fine, highly-defective microstructure, it is rather difficult to determine
definitively whether the carbide particles are still present within the shear band. However,
occasionally, local measurements which indicate high Mo levels were detected by EDS. If the
carbides are plastically elongated and have indeed wholly or partially dissolved, the Mo is still
locally enriched.
110
Figure 6.20: The region presented in this image falls along the border of the white band. On the left, the narrow strand of high-Ni
(running left to right) is shown to be over 100nm long with only a small grain of about 20nm in diameter in contrast. On the right,
the composition of the grain is isolated and highlighted in orange: it is about 20 wt.% Ni, and therefore came from T-generation
austenite. The MDP is BCC [111]- this grain is very fine, high in Ni but was mechanically unstable.
111
Figure 6.21: Another region close to the border of the white band; here, a small grain that dynamically recrystallized from L-
generation austenite (15-17 wt.%, marked in orange) has mechanically transformed to martensite.
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6.6: Discussion
By comparing the microstructure at various locations relative to the shear band in
addition to using the baseline microstructure in the undeformed QLT condition, the
microstructural evolution during high strain-rate deformation and adiabatic shear banding can be
reconstructed. Far from the shear localized region, the lath microstructure is intact and there was
no evidence for deformation-induced phase transformation of austenite as precipitates with both
L- and T-generation compositions were found to be mechanically stable. Ahead of the shear
band, austenite precipitates are extensively sheared and begin to divide into subgrains (by
‘mechanical recrystallization’), and despite the large strains needed to form these subgrains, both
L- and T-generation austenite are still observed to be mechanically stable.
Adjacent to the most intense shear localization, SEM images show that the precipitates
are stretched along the axis of the plane’s propagation. TEM of this region from the perspective
normal to the shear plane shows a lath-like subgrain microstructure parallel to the shear plane,
similar to the microstructure observed by Xue and Gray in 316L stainless steel [58,69] and Wang
and Kumar in a steel with the same alloy composition but subjected to a different heat treatment
[22]. Together, these two microstructural perspectives show that the austenite precipitates are
pancaked parallel to the shear plane and are broken into finer subgrains. Furthermore, EDS and
MDP together show that subgrains from both L- and T-generation austenite are stable in this
region.
The unetched portion of the band appears featureless in SEM and is approximately 4µm
thick. TEM analysis shows there is significant grain refinement through mechanical
recrystallization, again corroborating Gupta and Kumar’s [29] and Wang and Kumar’s [22]
observations on steels with this alloy composition. The grain size is significantly smaller than
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those ahead of the band, which suggests that the grains continue to refine as the band develops.
The ferrite shows evidence for twinning, the L-generation austenite has transformed to
martensite, some T-generation austenite has transformed to martensite, and some T-generation
austenite appears to be stable and twinned.
The overall microstructural evolution matches the dynamic recrystallization theory put
forth by Xue and Gray [58,69] and Meyers et al. [51,71]. As the shear localizes, grains break into
finer and finer subgrains; with increased strain and high temperature in the most developed
portion of the band, the subgrains rotate to form fine, equiaxed, dynamically recrystallized grains
on the scale of 50-200nm. In particular, we have shown that if the original microstructure is
inhomogeneous (here a mixture of relatively small precipitates in a lath martensite matrix), the
grain refinement mechanisms for both phases appear to operate in parallel. Additionally, the
smaller austenite precipitates refine to smaller recrystallized grains than the ferrite. However,
unlike the austenite in the ballistic samples from Zhang’s experiments [25], the majority of the
austenite throughout this sample is mechanically stable. It is only within the band itself, which
represents a small percentage of the total volume, do we observe instability, and even then, some
T-generation austenite remains stable.
Zhang showed that quasi-static, tensile loading transformed half of the austenite in QLT
(9.5vol% austenite down from 18.98vol% austenite), while ballistic testing left no detectable
austenite in the impact region [25]. However, there are significant differences between Kolsky
experiments and quasi-static or even ballistic loading experiments. Talonen et al. [92] showed
that austenite strained in quasi-static tension tests was more mechanically unstable than high
strain-rate (200 s-1
) tensile loading where the austenite particle size is substantially reduced by
mechanical recrystallization. Therefore, it is reasonable to expect the austenite in QLT to more
114
readily transform in tensile tests than in high-strain-rate experiments. Zhang’s ballistic field tests
have similar strain-rates as Kolsky experiments, but the loading conditions are quite different-
ballistic impact tests affect an area much larger than the projectile size, whereas the Kolsky anvil
is over four times larger than the sample. The ballistic plate can thus dissipate the impact energy
over a much larger area than the Kolsky sample. Therefore, while the strain-rate experienced in
both types of testing may be in the dynamic regime, energy dissipation is likely very different in
the two cases and direct comparison is tenuous.
Nevertheless, the results from this effort provide microstructural justifications for the
superior high strain-rate deformation response of 10Ni QLT. Unstable shear localization occurs
when thermal softening overrides work hardening. Here, several strain-hardening mechanisms
enabled by the thermally stable austenite content in the steel are observed in this material: the
precipitated austenite in this alloy which begins with varied composition (Ni content), size
distribution, and a low defect morphology has been shown to multiply dislocations, form twins,
and eventually experience deformation-induced phase transformation. Each of these
mechanisms, observed in the Kolsky sample, increases the overall work-hardening capacity of
the material, thereby increasing the resistance to ASB formation and propagation.
115
Chapter 7: Conclusions
7.1: Microstructural Evolution during Heat Treatment of an Fe-10Ni-0.1C Steel
C Partitioning During Tempering
i) In the austenitized and quenched condition, roughly half of the C content in the alloy
is in the martensitic matrix; the rest remains tied to the refractory elements in the
form of undissolved Mo- and V-rich refractory MC and M2C carbides. C-segregation
to interfaces was noted in APT specimens of the as-quenched material.
ii) After 40 minutes of L (650°C) temper, the ferrite was entirely depleted in C. Some of
this C partitioned to austenite, while the rest partitioned to Mo-rich carbide particles.
iii) Negligible C was found in both austenite and ferrite in specimens subjected to heat
treatments such as QLT, Q25L, Q25T’ and Q125T’. The decrease in C content in
austenite is believed to be due to a combination of a dilution effect as the austenite
grows and the further precipitation and growth of carbide particles during extended
tempering.
Together these findings imply that the role of C in influencing precipitated austenite stability
in this alloy, if any, is restricted to the early stages of the L temper. In addition, it likely plays
an important role in the nucleation of the austenite particles when tempering commences.
116
Ni Partitioning During Tempering
i) Ni-rich austenite is found to precipitate during isothermal tempering. At 650°C, the
austenite is found to contain 15-17 wt.% Ni; at 540°C, the austenite is more Ni-rich
and incorporates 22-24 wt.% Ni.
ii) For both isothermal tempering treatments, L and T’, the austenite is found to be
thermally unstable after extended tempering times. The onset of thermal instability is
believed to be due to the increased precipitate size, and so precipitates from shorter
tempers are presumed to be primarily stabilized by their small size.
iii) The two-stage QLT treatment is found to have two chemically distinct generations of
austenite. The Ni content in thermally stable austenite from the L-generation does not
equilibrate during the T temper (within the allocated time at this temperature), and so
both the 15-17 wt.% Ni austenite from the L temper and the 20-22 wt.% Ni austenite
from the T temper are present in the final microstructure.
iv) The austenite in QLT is also believed to be primarily thermally stabilized by the fine
particle size.
7.2: Microstructural Evolution during High Strain-Rate Deformation of an Fe-10Ni-0.1C
Steel Subjected to a Two-Stage Heat Treatment
i) The Fe-10Ni-0.1C steel was subjected to a range of strains and strain rates using a
compression Kolsky-bar testing facility. Different extents of shear localization were
generated by varying the test parameters, including no localization, diffuse ASB, and
well-developed ASB.
ii) Well-developed shear bands were found to be of the ‘transformed’ type and had
increased hardness across their width.
117
iii) Microstructural analysis at various locations relative to a well-developed shear band
was performed, including: far from the shear band (homogenous plastic deformation
zone), ahead of the shear band, adjacent to the shear band, and in the shear band.
a. Far from the shear band, there was no evidence of deformation-induced
transformation of precipitated austenite to martensite.
b. Ahead of the band, the austenite precipitates were found to have broken into
subgrains. Both the L-generation and T-generation austenite were found to be
mechanically stable after subgrain formation.
c. Adjacent to the shear band, the ferrite had a lamellar morphology parallel to the
shear plane. Former austenite precipitates are also seen to be sheared parallel to
the band and had broken into fine subgrains.
d. In the shear band, the ferrite was found to have dynamically recrystallized as
equiaxed grains between 50-300nm in size. Twinning was observed in some
ferrite grains. The austenite precipitates were found drawn into strands between
100-200nm in length, and within individual strands, to have recrystallized into
equiaxed subgrains of about 20-50nm in diameter. L-generation austenite was
found to be mechanically unstable; twinning was observed in some stable T-
generation austenite grains; other T-generation grains were observed to have
mechanically transformed to martensite.
118
Chapter 8: Recommendations for Future Work
There are still some open questions, answers to which would further improve our
understanding of the response of this alloy to heat treatment and deformation. These are
identified below.
It was observed that the as-quenched Fe-10Ni-0.1C steel had C enrichment at martensite
interfaces. Similar observations have been previously made in other steels [34–36]; specifically,
if these interfaces are rich enough in C, they may remain as stable austenite after the quench [34–
36]. The presence of nanoscale residual austenite films was not examined in this study, but if
present, may play an important role in subsequent tempering, for example as a nucleation site for
Ni-rich austenite. This aspect requires further attention.
Furthermore, it was observed that for a short tempering time (40 mins at 650°C), the Ni-
rich precipitated austenite appeared somewhat richer (or equivalent) in C relative to the
untempered martensite, suggesting that C quickly partitions from the martensite to the
precipitated austenite in the early stages of tempering. However, the rate and extent to which C
partitions in the early stages of tempering was not addressed in this study in detail.
A study of microstructure evolution using short tempering times (e.g. 1-30 minutes)
could address the above points, namely: i) the potential role of C-rich retained austenite films on
subsequent Ni-rich austenite evolution, and ii) quantify the rate and extent of C partitioning in
early tempering as a function of tempering time and temperature.
It was found that the Ni content in the austenite phase was spatially uniform for the T’
tempers but was quite varied in the Q25L sample. The origin of this behavior is not currently
119
understood. Analysis of a wider range of L tempering times could conceivably provide further
insight.
The relationship between austenite size and thermal stability was qualitatively determined
in this study but has not been quantitatively assessed. It was determined that the austenite
examined in this study (i.e. with Ni content between 15-25 wt.%) is unstable if sufficiently large,
but the critical size as a function of Ni content remains unknown. This aspect is relevant if the
volume fraction of austenite is to be controlled through modified QLT-type heat treatments.
Additionally, the composition of austenite in the various tempers was measured using APT and
STEM EDS, but a size-independent, composition-based model for predicting Ms temperature
was not used (e.g. as done by Jain et al. [28]). Similar analysis on the compositions measured in
the isothermal samples in this study may provide further insight for future alloy design.
Several aspects of the microstructural evolution of 10Ni QLT and similar alloys during
deformation are still not well understood. Zhang showed the austenite volume fraction in 10Ni
QLT decreased by about 50% when deformed by quasi-static tensile tests [25]. However, in light
of the detailed microstructural analysis done by Wang et al. on Mn-containing TRIP steels
[39,40], similar experiments could be done on 10Ni QLT. Specifically, they determined that
different sizes of austenite particles with similar composition undergo different deformation
mechanisms (TWIP vs TRIP) during quasi-static deformation in a ‘spectral TRIP effect’ [39,40].
Additionally, Yuan et al. also proposed a composition-based spectral TRIP effect in C-alloyed
steel [41]. This study has shown that the austenite in 10Ni QLT varies both in size and
composition, and therefore may also undergo a spectral TRIP effect during quasi-static
deformation. However, detailed analysis of microstructural evolution during quasi-static
120
deformation has not been performed and is likely to yield more fundamental insights into the
deformation behavior of this family of steels.
This study only examined the microstructural evolution in 10Ni QLT during dynamic
deformation. A comparative Kolsky-bar experiment with similar in-depth microstructural
analysis of the alloy subjected to single-stage L, T, and T’ temper treatments could help refine
and isolate primary deformation mechanisms that affect performance improvements.
121
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