MULTI-SCALE MICROSTRUCTURE AND THERMO-MECHANICAL ...

The Pennsylvania State University

The Graduate School

Department of Engineering Science and Mechanics

MULTI-SCALE MICROSTRUCTURE AND THERMO-MECHANICAL

CHARACTERIZATION FOR SHAPE MEMORY ALLOY DESIGN

VIA ADDITIVE MANUFACTURING

A Dissertation in

Engineering Science and Mechanics

by

Beth A. Last

2019 Beth Last

Submitted in Partial Fulfillment

of the Requirements

for the Degree of

Doctor of Philosophy

May 2019

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The dissertation of Beth A. Last was reviewed and approved* by the following:

Reginald F. Hamilton

Associate Professor of Engineering Science and Mechanics

Dissertation Advisor

Chair of Committee

Todd A. Palmer

Professor of Engineering Science and Mechanics

Clifford J. Lissenden


Albert E. Segall


Allison M. Beese

Assistant Professor of Materials Science and Engineering

Judith A. Todd

P. B. Breneman Department Head and Professor of Engineering Science and Mechanics

Head of the Department of Engineering Science and Mechanics

*Signatures are on file in the Graduate School

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ABSTRACT

The layer-by-layer deposition process of additive manufacturing (AM) offers the capability

to design material microstructures on multiple length scales. For NiTi shape memory alloys,

designing material microstructures using AM would allow for unparalleled tailoring of the

multiscale martensitic transformation and shape memory response. However, the laser-based

directed energy deposition (LDED) AM technique produces localized microstructures which are

distinct from those found in conventionally processed alloys. This work characterizes the grain and

precipitate microstructures on multiple length scales for LDED fabricated NiTi alloys and assess

the capability for tailoring the martensitic transformation morphology shape memory response

through post-deposition heat treatments.

Build coupons were fabricated by LDED AM using elementally blended Ni and Ti powder

feedstock. The use of elemental powders allowed for a Ti-rich and a Ni-rich powder feedstock

composition to be blended; thus, both shape memory effect (Ti-rich) and superelastic (Ni-rich)

behaviors were investigated. Specimens were extracted from the fabricated build coupons to

investigate the localized microstructure and shape memory behaviors. A full-field deformation

analysis technique was employed to correlate the AM microstructure to the deformation

mechanisms.

The results of this work show that the NiTi LDED AM builds are inherently spatially

varying on multiple microstructure length scales. The grain structure resulting from the AM process

was similar for both feedstock compositions: fine grains within the interfacial regions formed by

overlapping passes/layers and larger columnar grains within bulk regions (i.e. away from these

interfaces). As a result of the spatially varying microstructure, as built LDED NiTi alloys exhibit a

hardening like response and localized strain concentrations. Post-deposition heat treatment of the

Ni-rich alloys reduced the spatial variation in the Ni4Ti3 precipitate microstructure and increased

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the localized superelastic strains compared to the as built condition, with the solutionizing and

precipitation aging treatment resulting in the most spatially uniform Ni4Ti3 precipitate morphology.

For the LDED alloys, shape memory effect recovery strains of -4.0 % (for Ti-rich alloys) and

superelastic recovery strains of -6.0 % (for solutionized and aged Ni-rich alloys) were achieved.

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TABLE OF CONTENTS

List of Figures .......................................................................................................................... viii

List of Tables ........................................................................................................................... xvii

List of Abbreviations ............................................................................................................... xviii

Acknowledgements .................................................................................................................. xix

Chapter 1 INTRODUCTION ................................................................................................... 1

1.1 Unique Shape Memory Alloy Functionality and Novel Applications ....................... 1 1.2 Hierarchy of Shape Memory Behavior ...................................................................... 6 1.3 Connection of Martensitic Transformation to Microstructure ................................... 11 1.4 State of the Art for Additive Manufactured NiTi Shape Memory Alloys .................. 14 1.5 Problem Statement ..................................................................................................... 17 1.6 Thesis Format and Outline ......................................................................................... 18

Chapter 2 FABRICATION OF NITI MATERIALS USING LASER-BASED DIRECTED

ENERGY DEPOSITION ................................................................................................. 20

2.1 Chapter Overview ...................................................................................................... 20 2.2 Laser-Based Directed Energy Deposition Parameter Development .......................... 20

2.2.1 Elemental Ni and Ti Powder feedstock characterization and blending ........... 20 2.2.2 Parameter Optimization ................................................................................... 24

2.3 Build Coupon and Specimen Preparation for Microstructure and Thermo-

Mechanical Characterization .................................................................................... 33 2.3.1 Removal of build coupons from substrates ..................................................... 33 2.3.2 Sectioning of builds into specimens ................................................................ 34 2.3.3 Post-deposition heat treatments ....................................................................... 35

2.4 Fabrication of Powder Bed Fusion builds .................................................................. 38 2.5 Chapter Summary ...................................................................................................... 39

Chapter 3 METHODOLOGY FOR CHARACTERIZATION OF MICROSTRUCTURE

AND SHAPE MEMORY BEHAVIOR ........................................................................... 40

3.1 Chapter Overview ...................................................................................................... 40 3.2 Microstructure Characterization ................................................................................. 40

3.2.1 Sample Preparation ......................................................................................... 40 3.2.2 Microstructure characterization methods ........................................................ 42

3.3 Thermo-mechanical Characterization ........................................................................ 44 3.3.1 Thermal-Induced Martensitic Transformation, characterized by

Differential Scanning Calorimetry ................................................................... 44 3.3.2 Pseudoelastic and Superelastic Behavior, characterized by an Isothermal

Mechanical Experiment .................................................................................... 46

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3.3.3 Shape Memory Effect Behavior, characterized by an Isothermal

Mechanical Experiment and Subsequent Heating Cycle .................................. 50 3.3.4 Strain Measurement Length Scales ................................................................. 51

3.4 Correlating Additive Manufactured Microstructure and Martensitic

Transformation ......................................................................................................... 54 3.5 Chapter Summary ...................................................................................................... 59

Chapter 4 SPATIAL CHARACTERIZATION OF THE THERMAL-INDUCED PHASE

TRANSFORMATION THROUGHOUT AS-DEPOSITED ADDITIVE

MANUFACTURED NiTi BULK BUILDS ..................................................................... 60

4.1 Chapter Overview ...................................................................................................... 60 4.2 Results and Discussion ............................................................................................... 60 4.3 Chapter Summary and Conclusions ........................................................................... 65

Chapter 5 ANISOTROPIC MICROSTRUCTURE AND SUPERELASTICITY OF

ADDITIVE MANUFACTURED NiTi ALLOY BULK BUILDS USING LASER-

BASED DIRECTED ENERGY DEPOSITION .............................................................. 66

5.1 Chapter Overview ...................................................................................................... 66 5.2 Results and Discussion ............................................................................................... 67

5.2.1 As built Microstructure and Pseudoelastic Response ...................................... 67 5.2.2 Aged Microstructure and Pseudoelastic Response .......................................... 75

5.3 Chapter Summary and Conclusions ........................................................................... 80

Chapter 6 MULTI-SCALE SHAPE MEMORY EFFECT RECOVERY IN NiTi ALLOYS

ADDITIVE MANUFACTURED BY SELECTIVE LASER MELTING AND

LASER DIRECTED ENERGY DEPOSITION ............................................................... 81


6.2.1 Compositional Analysis and Phase Transformation Temperatures ................. 81 6.2.2 Microstructural Analysis ................................................................................. 83 6.2.3 Martensite Deformation and Shape Memory Effect Recovery ....................... 87 6.2.4 Stress-Strain-Temperature Cycling ................................................................. 91


Chapter 7 CORRELATING MICROSTRUCTURE AND SUPERELASTICITY OF

DIRECTED ENERGY DEPOSITION ADDITIVE MANUFACTURED Ni-RICH

NiTi ALLOYS ................................................................................................................. 98


7.2.1 Microstructure analysis ................................................................................... 98 7.2.2 Thermal-induced and stress-inducted martensitic transformation behavior .... 105


Chapter 8 Ni-CONCENTRATION DEPENDENCE OF DIRECTED ENERGY

DEPOSITED NiTi ALLOY MICROSTRUCTURES ..................................................... 117

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8.1 Chapter Overview ...................................................................................................... 117 8.2 Results and Discussion ............................................................................................... 117 8.3 Chapter Summary and Conclusions ........................................................................... 122

Chapter 9 SUMMARY AND CONCLUSIONS ...................................................................... 123

Chapter 10 RECOMMENDATIONS FOR FUTURE WORK ................................................ 124

10.1 Designing microstructure via process parameter control ......................................... 124 10.2 Tension-compression asymmetric response of AM NiTi alloys .............................. 125 10.3 Precipitate morphology microstructure length scale dictating SIMT ...................... 126 10.4 Functionally Graded NiTi Structures ....................................................................... 127

Appendix A LDED NiTi summary table ................................................................................. 133

Appendix B PBF NiTi summary table ..................................................................................... 135

Appendix C NiTi build coupon inventory ............................................................................... 138

Appendix D Sample and Specimen Extraction Locations ....................................................... 140

Appendix E Nontechnical Abstract ......................................................................................... 144

Appendix F Publications and Presentations ............................................................................ 145

References ................................................................................................................................ 147

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LIST OF FIGURES

Figure 1.1 Stress-strain behavior for a wrought Ti-6Al-4V material (data from [4]), and a

NiTi shape memory alloy (deformation processed Ni50.8Ti49.2 at.% alloy, in as received

condition from SAES). The recoverable strains for the shape memory alloy are

significantly greater than the recoverable elastic strains for Ti-6Al-4V. ......................... 2

Figure 1.2 Crystal structures for the (a) austenite and (b) martensitic phases for NiTi. The

parent phase, austenite, has a B2 cubic structure; the product phase, martensite, has a

B19' monoclinic structure. ............................................................................................... 6

Figure 1.3 Schematic of temperature and stress relationship between the parent and product

phases of NiTi. ................................................................................................................. 8

Figure 1.4 Gibb’s free energy for the martensitic transformation between the austenite (A)

and martensitic (M) phases. Ms is the martensitic start temperature, or the starting

temperature for the forward (austenite to martensite) transformation; As is the austenite

start temperature (start temperature for reverse transformation)...................................... 8

Figure 1.5 Shape memory behaviors: (a) Thermal-induced martentisitc transformation; (b)

Superelastic behavior due to the stress-induced martensitic transformation; (c) Shape

memory effect behavior where the intial phase is twinned martensite; and (d) Shape

memory effect behavior under an applied stress, where the intial phase is detwinned

martensite. A is the austenite or parent phase; M is the twinned martensitic phase; M+

is the detwinned martensitic phase. .................................................................................. 9

Figure 1.6 Stress-temperature relationship for a material exhibiting shape memory effect

or superelastic behaviors. Ms is the martensitic start temperature; Mf is the martensitic

finish temperature; As is the austenite start temperature; Af is the austenite finish

temperature; Md is the temperature beyond which stress-induced martensite does not

form. ................................................................................................................................. 10

Figure 1.7 Ni-Ti equilibrium phase diagram. The equiatomic NiTi phase is highlighted, as

this is the phase which undergoes the martensitic phase transformation and exhibits

shape memory behavior. .................................................................................................. 12

Figure 1.8 Relationships between Ni content and martensite start temperature (Ms). Data is

taken from [17]. ................................................................................................................ 12

Figure 1.9 (a) Summary plot of additive manufactured NiTi references by publishing date,

separated based on AM technique (laser based directed energy deposition (LDED);

powder bed fusion (PBF)) and review articles. (b) Feedstock compositions utilized in

these studies, based on the AM technique and separated based on alloyed powders

versus elementally blended powders. ............................................................................... 14

Figure 2.1 Frequency and cumulative frequency plots for the particle sizes for the elemental

Ni and Ti powders measured (a) when received and (b) prior to deposition. .................. 22

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Figure 2.2 Particle morphologies for (a) and (b) elemental Ni powder, and (c) and (d)

elemental Ti powder. Both powders have a spherical shape. The Ti powder has some

finer satellites attached to the larger particles. ................................................................. 23

Figure 2.3 Schematic of the custom-built laser-based directed energy deposition system. ..... 25

Figure 2.4 Measurement of powder feedrate, or mass flow rate of the feedstock powder,

based on the voltage (dial) reading of the powder feeder. This mass flow rate was

measured for the Ni46.9Ti52.1 at.% powder. ....................................................................... 26

Figure 2.5 Schematic of the (a) dimensions for a single pass (b) the dimensions for the

hatch spacing and layer thickness deposition parameters. The region of overlap

between adjacent passes and the interlayer region due to the deposition of successive

layers are highlighted. ...................................................................................................... 27

Figure 2.6 (a) Single pass, single layer builds fabricated with increasing mass flow rates.

Dimensions for the (b) width and (c) height for these builds fabricated using varying

mass flow rates. Builds were fabricated with 1000 W laser power, a scan speed of 10.6

mm/s, and a substrate temperature of 25 °C. The build deposited using the 10.2 g/min.

mass flow rate had the smallest standard deviation in the width and height. The scale

bar in (a) is in mm. (Build coupon ID: B2 (5.0 g/min.), B3 (6.1 g/min.), B4 (8.1

g/min.), B5 (10.2 g/min.)) ................................................................................................ 29

Figure 2.7 Build coupons deposited with 1 pass and 6 layers, using 1000 W laser power,

10.2 g/min. mass flow rate, scan speed of 10.6 mm/s, and a substrate temperature of

25 °C. The builds delaminated from the substrate as pointed out with the red arrows.

The small ticks on the scale bar are 1 mm. (Build coupon ID: B6, B7) .......................... 30

Figure 2.8 Build coupons deposited with 1 pass and 6 layers, using (a) 1000 W laser power

and a dwell time of 3 s between layers; and (b) 500 W laser power and a dwell time

of 0 s between layers. Coupons were deposited using 10.2 g/min. mass flow rate, scan

speed of 10.6 mm/s, and a substrate temperature of 25 °C. The builds delaminated

from the substrate as pointed out with the red arrow. The small ticks on the scale bar

are 1 mm. (Build coupon ID: B8 and B12) ...................................................................... 31

Figure 2.9 NiTi build coupons fabricated using the laser-based directed energy deposition

technique with 14 layers and (a) 1 pass; (b) 3 passes; and (c) 6 passes. .......................... 32

Figure 2.10 (a) Build coupons after being removed from the substrate. (b) Cross-section of

substrate, dilution region, and build coupon after the build coupon has been removed

from the substrate. (c) Optical microscopy image of the cross-section of a 6 pass, 14

layer build coupon fabricated from the Ti-rich powder blend. The dilution region is

labeled. ............................................................................................................................. 33

Figure 2.11 (a) Dogbone and (b) compression block geometries used in this work.

Dimensions are from [80]. ............................................................................................... 34

Figure 2.12 (a) Compression specimens heat treated in an Ar environment. One side has

been polished to reveal the NiTi material. The dark gray boundaries on the edge are

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the surface oxide layer. (b) Shown are two of the polished compression specimens

next to the original 8 x 4 x 4 mm compression specimen geometry. ............................... 36

Figure 2.13 Change in thickness (Δt) between specimen after solutionizing heat treatment

and after grinding to remove the oxide scale layer. Specimens were solutionized in a

flowing Ar environment (flowing Ar), encapsulated in a quartz tube with a vacuum

environment (vacuum quartz tube), or encapsulated in a quartz tube with a back-filled

Ar environment (Ar quartz tube). ..................................................................................... 36

Figure 2.14 (a) Specimens which have been encapsulted in quartz tubes. (b) Vacuum sealed

quartz tubes after a 1000 °C, 24 h heat treatment. A film has formed on the inside of

the quartz tube. (c) Back-filled Ar sealed quartz tubes after a 1000 °C, 24 h heat

treatment........................................................................................................................... 37

Figure 3.1 DSC thermogram for an additive manufactured Ti-rich specimen, which shows

a single forward transformation peak and a single reverse transformation peak. The

martensitic start (Ms) and finish (Mf) transformation temperatures for the forward

martensitic transformation and the austenitic start (As) and finish (Af) transformation

temperatures for the reverse transformation, as well as the martensitic peak (Mp) and

austenitic peak temperatures (Ap), are identified. The integrated regions, identified in

blue and red, correspond to the enthalpies for the forward (ΔHF) and reverse (ΔHR)

transformations, respectively. .......................................................................................... 46

Figure 3.2 (a) Overview of the thermo-mechanical experimental set-up. Close-up figures

of (b) tension and (c) compression specimens within the load frame. ............................. 48

Figure 3.3 Pseudoelastic stress-strain response and determined shape memory and material

properties: critical transformation stress (σA→M), stress hysteresis (Δσ), maximum

applied strain (εmax), recovered or transformation strain (εrec), irrecoverable strain

remains (εirrec), Young’s modulus during loading (EA), Young’s modulus during

unloading (EM). ................................................................................................................ 49

Figure 3.4 Shape memory effect stress-strain-temperature response and determined shape

memory and material properties: critical stress of the twinned martensite (σM→M+),

residual strain (εres), recovered strain (𝛆recSME), Young’s modulus of the twinned

martensite (EM), Young’s modulus of the reoriented or detwinned martensite (EM+),

recovery start temperature (As*), and recovery finish temperature (Af

*). .......................... 51

Figure 3.5 Defined gage lengths for (a) engineering and true strain (b) strain using the

miniature extensometer on a tension and compression specimen, and (c) virtual gage

length for strain contours calculated using digital image correlation. (d) Compressive

mechanical responses from a single experiment for a conventionally processed

Ni49.9Ti50.6 at.% alloy. The engineering strain is plotted versus engineering stress

(black). The true strain is plotted versus the true stress (blue). The extensometer strain

is plotted versus the engineering stress (red) and the virtual extensometer strain is

plotted versus engineering stress (green). ........................................................................ 53

Figure 3.6 Machine vision systems with the camera and series of lenses ............................... 55

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Figure 3.7 (a) Low (0.75X magnification) and (b) high (6X) magnifications of a good

speckle pattern for digital image correlation. The speckle pattern has high contrast

between the white background and black speckles, there is a random and unique

speckle pattern, and the speckles cover approximately 3 by 3 pixels (in low

magnification image). ...................................................................................................... 56

Figure 3.8 (a) Gray-scale histogram for the “good” speckle pattern (shown in Figure 3.8

(a)). Comparison of the rigid body displacement (RBD) and the calculated RBD as

determined using the speckle pattern and digital image correlation, for RBDs of (b)

0.0 mm (stationary) and (c) incremental 0.2 mm. ............................................................ 57

Figure 3.9 Scale comparison of fabrication, microstructure, behavior, and martensitic

transformation aspect. The build coupon size and dimensions of the passes and layers

correspond to fabrication. The scale at which the microstructure was characterized is

identified by the composition, grains, and precipitates. The gage lengths for the

extensometer and virtual DIC extensometer are identified for the behavior. The size

of the martensite variants is significantly smaller. ........................................................... 59

Figure 4.1 Schematic illustrations of the build coupon showing (a) sections extracted along

the build direction [y0 – the beginning, ym – the middle, and ye – the end of the build

coupon] and (b) locations for differntial scanning calorimetry sample extraction along

the height/z-direction. ...................................................................................................... 61

Figure 4.2 (a) Cooling and (b) heating DSC thermograms for Ti-rich as built samples taken

at the same build height from sections y0, ym, and ye along the build direction. The

martensite and austenite start (Ms and As), finish (Mf and Af), and peak (Mp and Ap)

temperatures are identified. The forward (HF) and reverse (HR) enthalpy

measurements are derived from the peak area. ................................................................ 61

Figure 4.3 DSC thermograms with forward (Mp) and reverse peak (Ap) temperatures for

samples taken from section (a) y0 (b) ym, and (c) ye along the build/y-direction. ............. 62

Figure 4.4 TIMT, as measured using differential scanning calorimetry, for samples

extracted from the as built Ni-rich coupon. No transformation peaks are observed in

this temperature range. ..................................................................................................... 63

Figure 4.5 DSC thermograms and peak temperatures for Ni-rich samples that were directly

aged at (a, b) 450 °C, (c, d) 500 °C, and (e, f) 550 °C for various time durations.

Samples were all extracted from the same build height. .................................................. 64

Figure 5.1 XRD analysis taken from z-height locations of 2 mm and 8 mm above the

substrate of an as built Ni-rich compression specimen extracted from a large (6 pass,

14 layer) build coupon. .................................................................................................... 67

Figure 5.2 Back scatter electron images of a Ni-rich specimen micromachined from a 6.5

mm z-height (far from the substrate). Images were taken from two locations along the

build direction, a distances of 25 and 31 mm from the build origin. ............................... 68

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Figure 5.3 SEM images of the as built compression specimen, which are at (a) 8.5 (b) 6,

and (c) 2.5 mm z-heights. ................................................................................................. 69

Figure 5.4 (a) Tensile stress-strain response for as built specimens extracted 3.4 mm and

6.7 mm above the substrate. DIC axial strain contour images for specimens extracted

(b) 3.4 mm and (c) 6.7 mm above the substrate. The image numbers in (b) and (c)

correspond to the numbers in (a). The strains are axial measurements along the

loading direction and parallel to the build/y-direction. .................................................... 72

Figure 5.5 (a) Tensile transverse strain-axial strain response for as built specimens

extracted 3.4 mm and 6.7 mm above the substrate. Transverse and axial strains are

measured using virtual extensometer from DIC. Poisson’s ratio (νYX) is the ratio of the

transverse strain (εX) to the axial strain (εY). DIC transverse strain contour images for

specimens extracted (b) 3.4 mm and (c) 6.7 mm above the substrate. The image

numbers in (b) and (c) correspond to the numbers in (a). The strains are transverse

measurements perpendicular the loading direction and parallel to the pass-/x-

direction. .......................................................................................................................... 73

Figure 5.6 (a) Compressive stress-strain response for an as built specimen. The macroscale

strains are axial measurements along the loading direction and parallel to the build

height (z-dir). (b) DIC axial strain contour images numbered corresponding to

numbers along the stress-strain curve. Unloading images below loading images

correspond to the same macroscale strain. ....................................................................... 74

Figure 5.7 XRD analysis for a directly aged (550 °C, 3 h) Ni-rich sample. ............................ 76

Figure 5.8 Back scatter electron images of the aged (550 °C, 3 h) compression specimen,

which were taken at z-heights of (a) 8.5 mm; (b) 6 mm; and (c) 2.5 mm, measured

from the substrate. ............................................................................................................ 76

Figure 5.9 (a) Tensile stress-strain response and (b) DIC axial strain contour for aged (550

°C, 3 h) specimen, extracted at a height far from the substrate (5.3 mm). The strains

are axial measurements along the loading direction and parallel to the build/y-

direction. (c) Tensile transverse strain-axial strain response for aged measured using

virtual extensometer from DIC. Poisson’s ratio (νYX) is the ratio of the transverse

strain (εX) to the axial strain. (d) DIC transverse strain contour images for points

identified in (a). The strains are transverse measurements perpendicular the loading

direction and parallel to the pass-/x-direction. ................................................................. 78

Figure 5.10 (a) Compressive stress-strain response for an aged (550 °C, 3h) specimen. The

macroscale strains are axial measurements along the loading direction and parallel to

the build height (z-dir). (b) DIC axial strain contour images numbered corresponding

to numbers along the stress-strain curve. Unloading images below loading images

correspond to the same macroscale strain. ....................................................................... 79

Figure 6.1 Thermal-induced phase transformation temperatures (Ms, Mf, As, Af) and

enthalpies (HR, HF) measured from DSC analysis of samples sectioned along the build

height (z-direction) of (a) and (b) PBF and (c) and (d) LDED compression specimens.

.......................................................................................................................................... 83

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Figure 6.2 Back scatter electron images of the precipitate morphologies for (a) PBF and

(b) LDED alloys at z-heights of 7 mm, 4 mm, and 1 mm. ............................................... 84

Figure 6.3 XRD scans of the as built (a) PBF and (b) LDED specimens at various z-heights.

Specimens were mechanically cycled prior to analysis, which may have resulted in a

preferential orientation observed by the strong intensity for the 42° 2θ peak for the

LDED specimen. .............................................................................................................. 85

Figure 6.4 Optical microscopy images of the grain structure for the (a) PBF and (b) LDED

specimens at varying z-heights. ....................................................................................... 86

Figure 6.5 Compression stress-strain-temperature curves for (a) PBF and (b) LDED NiTi

alloys initially in the martensitic phases at T = 23 °C. Material parameters are defined

in Table 6.3. ..................................................................................................................... 88

Figure 6.6 Compressive stress-strain-temperature responses for a PBF NiTi alloy showing

multi-scale SME recovery using (a) macro-scale extensometer measurements and (b)

meso-scale in-situ full-field measurements. (b) In-situ DIC strain contour images with

numbers corresponding to those along the loading, unloading, heating, and cooling

segments in (a). ................................................................................................................ 90

Figure 6.7 Compressive stress-strain-temperature responses for an LDED alloy showing

multi-scale SME recovery using (a) macro-scale extensometer measurements and (b)

meso-scale in-situ full-field measurements. (b) In-situ DIC strain contour images with

numbers corresponding to those along the loading, unloading, heating, and cooling

segments in (a). ................................................................................................................ 91

Figure 6.8 Stress-strain-temperature cycling up to n = 15 cycles for (a) PBF and (b) LDED

alloys. The n=1 corresponds to Figure 6.6 (a) and Figure 6.7 (a). After unloading, the

specimen was heated for SME recovery. (c) Corresponding strain recovery ratios (η).

.......................................................................................................................................... 92

Figure 6.9 Full-field strain measurement contour images. (a) and (b) DIC analysis used

the undeformed image in n = 1 as the reference image for correlation in order to

determine the full-field strain measurement at the beginning of cycles n = 2, 3, 5, 10,

and 15. (c) and (d) In-situ DIC analyses for n=15 with the reference image at the start

of the cycle. ...................................................................................................................... 93

Figure 7.1 Back scatter electron images showing the microconstituent morphologies for (a)

as built alloys and alloys heat treated at 950 °C for (b) 10 h and (c) 24 h durations. In

(a) and (b), lenticular microconstituents are Ni4Ti3 precipitates and Ni3Ti secondary

phases appear globular. .................................................................................................... 99

Figure 7.2 XRD scans with increasing post-deposition heat treatment duration at 950 °C.

The specified locations are along the build height locations. Phases have been

identified as B2 ; B19' ; R-phase ; Ni4Ti3 ; and Ni3Ti . ........................................ 100

Figure 7.3 Evolution of composition with increasing post deposition heat treatment

duration at 950 °C. The dashed horizontal line is the input powder feedstock

xiv

composition. Circles represent average compositions. Squares represent maxima and

minima. ............................................................................................................................ 100

Figure 7.4 (a) Grain boundary structure, (b) grain orientation map, and (c) pole figure for

an as built Ni-rich compression specimen. The grain orientation is identified as normal

the specimen surface. ....................................................................................................... 102

Figure 7.5 (a) Grain boundary structure and (b) grain orientation map for a solution treated

(950 °C, 24 h) specimen. (c) Corresponding pole figure. ................................................ 103

Figure 7.6 Spatial resolution of the lenticular Ni4Ti3 precipitate morphology. Build height

(in the z-direction) dependence of the Ni4Ti3 precipitate morphology as a result of

solutionizing and aging (SL+Aged) versus directly aging as built material (Aged). (a)

Area fraction with selected SEM images inset and (b) Length along the major axis of

the lenticular precipitate. .................................................................................................. 104

Figure 7.7 (a) Forward and (b) reverse thermal-induced martensitic transformations for

directly aged (550 °C, 3h) and solutionized and aged (950 °C, 24 h followed by 550

°C, 3h) alloys. .................................................................................................................. 106

Figure 7.8 (a) Compressive stress-strain response for an as built alloy. The austenitic elastic

modulus EA is defined on the loading curve. The 0.2% offset critical transformation

stress (σA→M) and stress at the onset of elastic unloading are illustrated and specified

along the ordinate axis. The stress hysteresis (∆𝜎) is the differential between σA→M

and that onset stress. (b) Full-field axial strain contour images showing meso-scale

strain evolutions. Above each image are macro-scale strain (%) values, which are

identified by the points along the stress-strain curve in (a). ............................................. 107

Figure 7.9 (a) Compressive stress-strain response for a solutionized (950 °C, 24 h) Ni-rich

alloy. The austenitic elastic modulus EA, critical transformation stress (σA→M) and

stress at the onset of elastic unloading are illustrated and specified along the ordinate

axis. The stress hysteresis (∆𝜎) is the differential between σA→M and that onset stress.

(b) Full-field strain contour images showing meso-scale strain evolutions. Above each

image are macro-scale strain (%) values, which are identified by the points along the

stress-strain curve in (a). .................................................................................................. 108

Figure 7.10 (a) Compressive stress-strain response for an aged (550 °C, 3 h) alloy. The

austenitic elastic modulus EA is defined on the loading curve. The 0.2% offset critical

transformation stress (σA→M) and stress at the onset of elastic unloading are illustrated

and specified along the ordinate axis. The stress hysteresis (∆𝜎) is the differential

between σA→M and that onset stress. (b) Full-field strain contour images showing

meso-scale strain evolutions. Above each image are macro-scale strain (%) values,

which are identified by the points along the stress-strain curve in (a). ............................ 109

Figure 7.11 (a) Compressive stress-strain response for a solutionized and aged (950 °C, 24

h followed by 550 °C, 3 h) alloy. The austenitic elastic modulus EA, critical

transformation stress (σA→M) and stress at the onset of elastic unloading are illustrated

and specified along the ordinate axis. The stress hysteresis (∆𝜎) is the differential

between σA→M and that onset stress. (b) Full-field strain contour images showing

xv

meso-scale strain evolutions. Above each image are macro-scale strain (%) values,

which are identified by the points along the stress-strain curve in (a). ............................ 110

Figure 7.12 Full-field transverse strain contour images showing meso-scale strain

evolutions for the as built Ni-rich alloy. Above each image are macro-scale axial strain

(%) values, which are identified by the points along the stress-strain curve in Figure

7.8 (a). .............................................................................................................................. 111


evolutions for the solution treated alloy (950 °C, 24 h). Above each image are macro-

scale axial strain (%) values, which are identified by the points along the stress-strain

curve in Figure 7.9 (a). ..................................................................................................... 111


evolutions for the directly aged specimen (550 °C, 3 h). Above each image are macro-

scale axial strain (%) values, which are identified by the points along the stress-strain

curve in Figure 7.15 (a). ................................................................................................... 112


evolutions for the solution treated and aged specimen (950 °C, 24 h followed by 550

°C, 3 h). Above each image are macro-scale axial strain (%) values, which are

identified by the points along the stress-strain curve in Figure 7.17 (a). ......................... 112

Figure 8.1 Schematics of interfacial regions formed by overlapping passes and/or layers.

(a) Build coupon XY plane spanned by pass and build directions with periodic

interfacial regions formed by overlapping passes. (b) Build coupon YZ plane spanned

by build and height directions with periodic interfacial regions formed by overlapping

layers. (c) Build coupon XZ plane spanned by pass and height directions with

interfacial regions formed by overlapping passes and layers. .......................................... 118

Figure 8.2 XY, YZ, and XZ cross-sectional views of the Ni-rich NiTi alloy build. (XY)

Elongated/columnar grains are identified in the (a) bulk and equiaxed grains exist in

the (b) interpass/interfacial region. (YZ) Elongated grains exist in the region (a) and

away from the interlayer/interfacial regions. (XZ) Columnar grains exist within and

adjacent to interfacial (both interpass and interlayer) regions (b) .................................... 119

Figure 8.3 XY, YZ, and XZ cross-sectional views of the Ti-rich NiTi alloy build. (XY)

Elongated/columnar grains are identified in the (a) bulk and equiaxed grains exist in

the (b) interpass/interfacial region. (YZ) Elongated grains exist in the region (a) and

away from the interlayer/interfacial regions. (XZ) Columnar grains exist within and

adjacent to interfacial (both interpass and interlayer) regions (b) .................................... 120

Figure 8.4 SEM images showing precipitate morphologies for the Ni-rich alloy (a) within

an interfacial region and (b) in the bulk. (c) SEM images of the Ti-rich alloy ................ 121

Figure 10.1. Plan for extracting tension and compression specimens with the same loading

axis direction (a) Tension specimens with the width of the gage section parallel to the

layers. (b) Tension specimens with the width of the gage section perpendicular to the

layers (Build coupon IDs: B49, B50). .............................................................................. 126

xvi

Figure 10.2 (a) Image shows the High Power-High Deposition (HPHD) additive

manufacturing system. The HPDP system is within the glove box enclosure. (b)

Deposition head with four coaxially-fed powder nozzles. The hot plate and substrate

and the build coordinate axes are shown. (c) Close-up of the four powder nozzles. ....... 128

Figure 10.3 Measurement of powder mass flow rate of the feedstock powder, based on the

voltage (dial) reading of the powder feeder for the elemental Ti and elementally

blended NiTi powders. ..................................................................................................... 129

Figure 10.4 Deposition of compliant mechanism structures using the high power-high

deposition system. (a) During deposition, powder is directed into the melt pool formed

by the laser. (b) After deposition, the build is still hot at the top of the build as heat is

conducted through the previously deposited material. ..................................................... 131

Figure 10.5 Functionally graded and compliant mechanism build coupons fabricated using

elemental Ti and elementally blended NiTi powder using the high power-high

deposition system. ............................................................................................................ 132

xvii

LIST OF TABLES

Table 2.1. Calculated metrics from light diffraction analysis for elemental Ni and Ti

powders. The data from which these metrics are calculated are shown in Figure 2-1.

D10 and D90 refer to the particle size which encompasses 10% and 90% of the

cumulative particles. ........................................................................................................ 22

Table 2.2 Processing parameters for NiTi additive manufactured builds fabricated using

the laser-based directed energy deposition technique. ..................................................... 28

Table 3.1 Calculated Ni:Ti composition ratios for NiTi phases and secondary phases used

with EDS to identify phases. ............................................................................................ 43

Table 3.2 Crystal structures and parameters for the NiTi phases and secondary phases

identified using x-ray diffraction. The PDF# corresponds to the inorganic crystal

structure database (ICSD) number, or the reference crystal structure utilized to

identify the phases in the additive manufactured alloys. ................................................. 43

Table 6.1 Average composition measurements for compression specimens extracted from

a large build coupon. Specimen 1 was further sectioned along the build height. ............ 82

Table 6.2 AM fabricated Ti-rich NiTi material properties. ..................................................... 88

Table 7.1 Calculated material and shape memory properties for the as built, directly aged,

solutionizied, and solution treated and aged compressive responses. .............................. 115


the high power-high deposition (HPHD) laser-based directed energy deposition

system. ............................................................................................................................. 130

xviii

LIST OF ABBREVIATIONS AM additive manufacturing

AOI area of interest

ARL Applied Research Laboratory

BSE backscatter diffraction

CIMP-3D Center for Innovative Materials Processing through Direct Digital Deposition

CP commercially pure

CS conventional sintering

DIC digital image correlation

EBSD electron backscatter diffraction

EDM electrical discharge machining

EDS energy dispersive x-ray spectroscopy

FGM functionally graded material

HIP hot isostatic pressing

HPHD high power-high deposition

ICSD inorganic crystal structure database

LARS Laser Articulating Robotic System

LDED laser-based directed energy deposition

MIM metal injection molding

NSF National Science Foundation

PBF power bed fusion

PE pseudoelasticity

PM powder metallurgy

ROI region of interest

SE superelasticity

SEM scanning electron microscopy

SLM selective laser melting

SME shape memory effect

SMA shape memory alloy

XRD x-ray diffraction

xix

ACKNOWLEDGEMENTS

I would like to very sincerely thank and acknowledge Dr. Reginald Hamilton for all his

time, patience, motivation, help – I could not have done it without all your encouragement. I would

also like to thank Dr. Todd Palmer for all his assistance and expertise. I would like to thank my

thesis committee members: Dr. Albert Segall, Dr. Clifford Lissenden, and Dr. Allison Beese.

I would like to thank the other members in the Multifunctional and Adaptive Materials

Laboratory: (Dr.) Asheesh Lanba, (Dr.) Huilong Hou, Nicholas Costa, Jessica Spoll, Richard

LaSalle, Shashank Nagrale, and Emily Jenkins, for training me, assisting with experiments, and

discussing results which culminated in this dissertation. I would like to thank our collaborators:

Mohsen Taheri Andani and Dr. Mohammad Elahinia at the University of Toledo. I would also like

to thank these students from other research groups: (Dr.) Justin Kauffman, Andrew Iams, Scott

Meredith, and James Zuback.

I would like to acknowledge the assistance of Griffin Jones and Jay Tressler, for assisting

with deposition of the NiTi additive manufactured builds; Jay Keist, for assisting with

metallography and microscopy; Ed Good, for assistance with metallography; and The Center for

Innovative Materials Processing through Direct Digital Deposition for use of their laser deposition

systems. I would also like to thank the Materials Characterization Lab staff in the Millennium

Science Complex for their assistance with sample preparation and microstructure characterization.

Funding for my work was came from the National Science Foundation (NSF) Graduate

Fellowship under Grant No. DGE1255832 and the NSF Grant No. CMMI 1335283. Any opinions,

findings, and conclusions or recommendations expressed in this material are those of the author(s)

and do not necessarily reflect the views of the NSF.

Thank you to all my family and friends as well, who helped me to push through and stay motivated.

1

Chapter 1

INTRODUCTION

1.1 Unique Shape Memory Alloy Functionality and Novel Applications

Shape memory alloys (SMAs) are able to recover large deformations (strains) due to an

underlying martensitic phase transformation (MT) [1,2]. The MT is a solid state and diffusionless

transformation between a parent/austenitic and a product/martensitic phase. For SMAs, the MT is

reversible, meaning that the reverse MT (product to parent) follows the same pathway as the

forward transformation (parent to product) and the applied strain is recovered as the reverse

transformation occurs. The amount of recoverable strain is between 6-8% [3]. For non-SMA

materials, the recoverable strain corresponds to strain before the material has yielded. For non-

SMA metals, this strain is significantly less than the recoverable strain for NiTi SMAs. Figure 1-1

shows representative stress-strain responses for a wrought Ti-6Al-4V material and a deformation

processed Ni50.9Ti49.1 at.% alloy obtained from SAES Smart Materials Incorporated. Failure of the

Ti-6Al-4V alloy occurs at a strain around 17%; the response is truncated to the 0.2% offset yield

strength as permanent plastic deformation occurs beyond this point, and strain cannot be recovered.

By comparison, the NiTi SMA can recover more than six times the amount of strain that the Ti-

6Al-4V alloy can recover. This behavior is one of the unique responses of SMAs and is referred to

as the pseudoelastic (PE) or superelastic (SE) behavior.

2

Figure 1.1 Stress-strain behavior for a wrought Ti-6Al-4V material (data from [4]), and a

NiTi shape memory alloy (deformation processed Ni50.8Ti49.2 at.% alloy, in as received

condition from SAES). The recoverable strains for the shape memory alloy are

significantly greater than the recoverable elastic strains for Ti-6Al-4V.

Shape memory alloys are currently used in the automotive, aerospace, robotic, and

biomedical fields. Boeing developed a device which reduced noise and increased cruise efficiency

by controlling the deflection of SMA chevrons on the exhaust side of a jet engine [5]. NASA has

recently developed a wheel made from NiTi wires. These tires are able to deform while going over

rough terrain and, unlike pneumatic wheels, cannot be punctured [6]. NiTi wires are also used to

make stents; the NiTi alloy is deformed prior to insertion and the body’s temperature results in the

stent expanding to keep the blood vessel open [5]. NiTi alloys are also being used for bone implants

[5]. The NiTi response exhibits a stress hysteresis, shown in Figure 1-1, which closely mimics the

response of bone [7]. Other alloys which have been used for bone implants, like Ti-6Al-4V [8] or

stainless steel [7], do not exhibit a stress hysteresis. Additionally, the Young’s modulus of

pseudoelastic NiTi (25-50 GPa [3,9]) is close to that of cortical bone (10-20 GPa [10]). The moduli

of wrought Ti-6Al-4V and stainless steel are significantly higher (120 GPa for Ti-6Al-4V; 190 GPa

1000

800

600

400

200

0

Str

ess (

MP

a)

86420Strain (%)

Ni50.9Ti49.1 at.%

wrought Ti-6Al-4V

stress hysteresis

E = 120 GPa

E = 50 GPa

3

for stainless steel). Bone implants with responses more similar to natural bone are advantageous,

as they minimize stress shielding and bone fracturing [11–14].

Current applications of NiTi shape memory alloys utilize deformation processed NiTi

materials [15]. The fabrication process for commercially available deformation processed NiTi

consists of multiple steps: melting and casting of an ingot, hot- and/or cold-working to a final form,

and a shape memory (thermo-mechanical) treatment to improve shape memory behavior [1,2]. The

melting and casting process must be completed in a vacuum or an inert environment, to minimize

impurity pick-up. Moreover, ingots are typically re-melted multiple times to improve the

homogeneity of the composition [1,7]. Homogeneity is critical, as small local deviations in the

composition can have a significant impact on the shape memory behavior. For example, the

martensitic start temperature (Ms), the temperature at which the MT begins, can deviate by 10 °C

with a 0.1 at.% increase in Ni content [16]. Currently, casting processes have excellent control over

the Ms temperature, with the variation in temperature less than ±5 °C [1,7].

After casting, ingots are then processed into bar, sheet, rod, or wire forms using

deformation processing techniques like cold- or hot-working [7,17]. Cold working of NiTi is

difficult, especially at higher Ni contents, due to a strong work hardening response [15]. Hot

working of NiTi typically results in a surface oxide layer which needs to be removed. To achieve

more complex geometries, the bars, sheets, rods, or wires are then further machined. High tool

wear, undesirable chip formation, and the formation of burrs during turning are some of the

problems associated with machining of NiTi [7]. Machining also results in localized plastic

deformation, work hardening, and microstructure refinement, which locally alter the shape memory

response [18–24]. Thus, near-net shape fabrication techniques that minimize machining are of

interest.

Powder metallurgy (PM) has been proposed as a near net-shape technique for NiTi alloys.

Conventional sintering (CS), hot isostatic pressing (HIP), and metal injection molding (MIM)

4

methods are examples of PM techniques. For CS and MIM processes, powder is mixed with

lubricants (CS) or a binder (MIM) and then compacted or pressed into a green part. The green part

is then heated to form the final product. An additional HIP step can follow, to produce a fully dense

(greater than 99% dense) part, since as-sintered parts are not fully dense. The input material is either

alloyed NiTi or blended elemental Ni and Ti powders [7]. The part composition can be easily

altered by changing the initial feedstock composition. However, as PM processes typically only

sinter (not fully melt) the powder, local variations in composition may be present in the final part

due to compositional segregation during solidification. Alloyed NiTi powders typically result in

more homogeneous compositions and fewer secondary phases relative to PM parts fabricated from

elemental powders [25]. Undesired secondary phases are commonly observed in PM parts. Since

these precipitates do not undergo the reversible MT, they limit the achievable reversible strain

[7,26]. Impurity (i.e. O, C, N) pick-up is common due to the large surface area to volume ratios of

powder particles. Thus, handling of the powder, especially pure Ti powder, must be carefully

controlled to avoid impurity pick-up. The geometries for cast, PM, and MIM parts are also limited

to mold geometries (and ability to be removed from the mold).

Additive manufacturing (AM) has been proposed as an alternative net-shape and near-net

shape fabrication process for NiTi alloys. AM technologies build parts up layer-by-layer [27,28],

as opposed to subtractive manufacturing, where material is removed. By only solidifying material

in desired locations, there is either no (net) or minimal (near-net) post-processing and machining

needed to achieve the final part geometry. Beyond conventional processing techniques, AM offers

shape-, hierarchical-, material-, and functional- complexities [28]. The X, Y, and Z control over

where material is deposited means that highly complex shapes can be fabricated. More advanced

systems have 5-axis (XYZ, rotary axis, swivel axis) direction control for increased shape

complexity. For hierarchical complexity, multi-scale (micro-, meso-, and macro-scale) structures

can be designed and built; the multi-scale structures can be geometrical features integrated into the

5

part, like lattice structures [29], or microstructure features where specific features within the larger

part were designed to have location-specific microstructures. In terms of material complexities, the

feedstock material can be changed during processing, so that the composition of the part can be

altered as well. For functional complexity, full assemblies and structures can be fabricated at one

time [28].

Many current applications of AM benefit from the advantages AM offers. General

Electric’s LEAP fuel nozzle is an example of functional complexity. The cobalt-chrome alloy fuel

nozzle reduces the assembly from 20 individual parts to a single AM part [30]. Patient specific

dental implants are an example of shape complexity; the implants are design to fit perfectly for one

single patient. AM can advance current NiTi bone implants, by locally tailoring the material

response to match that of bone and by fabricating a surface which will form a better bond with bone

[11–14]. Human bone has a soft and porous interior (cancellous) bone and a harder exterior

(compact) bone. By controlling the AM deposition parameters, intentional porosity can be

introduced during fabrication to locally alter the NiTi bone implant response. Some traditional bone

implants have adhesion issues between the implant and the bone [8]. By controlling deposition

parameters, an optimized surface topology can be fabricated to minimize adhesion complications.

Additionally, the localized composition of the implant can be altered by changing or grading the

powder feedstock. This can be used to locally alter the mechanical response to more closely mimic

real bone.

Additive manufacturing also opens the potential for on demand structural and material

design. This is especially advantageous for NiTi shape memory alloys, as the ability to locally

control the composition, microstructure (i.e. porosity, grains, precipitates), and geometrical

structure will take advantage of the hierarchical nature of the shape memory response.

6

1.2 Hierarchy of Shape Memory Behavior

The martensitic phase transformation (MT) between the austenite and martensite phases

occurs in the atomic and micro-scales, resulting in an associated macro-scale shape change. The

macro-scale shape change due to the underlying MT gives rise to the superelastic/pseudoelastic and

shape memory effect (SME) behaviors. On the atomic scale, the parent austenite (A) phase

transforms to the product martensitic (M) phase. The crystal structures are shown in Figure 1-2.

Austenite is the high-temperature parent phase and has a cubic B2 crystal structure (Figure 1-2 (b)).

The martensite phases for NiTi SMAs typically exhibit a monoclinic (B19') crystal structure

[1,31,32] (Figure 1-2 (b)). The martensite phase, highlighted in red, is shown within the austenite

phase in Figure 1-2 (a).

(a) B2 Austenite (b) B19' Martensite

Figure 1.2 Crystal structures for the (a) austenite and (b) martensitic phases for NiTi. The

parent phase, austenite, has a B2 cubic structure; the product phase, martensite, has a B19'

monoclinic structure.

The martensitic transformation does not occur through diffusion. Rather it proceeds as a

cooperative movement of atoms. The cubic austenite structure is elongated to the monoclinic

martensitic structure. The crystal planes and directions for the austenite structure are related to the

martensitic structure through lattice correspondence, as identified in Figure 1.2. The transformation

requires a finite energy input (driving force) to proceed [1]. This energy can be an applied

mechanical force or a changing temperature.

7

The MT proceeds by nucleation and growth phenomena. As martensite is formed within

the parent austenite structure, a strain arises at the interface between martensite and austenite.

Twinning or slip lattice invariant shear mechanisms accommodate this strain. For SMAs, the

twinning mechanism usually occurs, rather than slip. The twinned martensitic structure,

schematically shown in Figure 1.3, naturally forms when cooling under zero applied stress from

the austenite. The twinned martensite forms as plates of corresponding martensite variant pairs

(CVPs). One variant of martensite has the same crystal structure and orientation. For NiTi, there

are 24 possible variants that can from during the MT from the austenite parent structure [33]. The

transformation of austenite to twinned martensite is accommodated by no volume change. The twin

boundaries are mobile; as stress is applied to the twinned martensitic structure, the twin boundaries

between variants move to accommodate the applied deformation. As the twin boundaries move,

more favorably oriented martensite variants grow at the expense of less favorably oriented variants

[1,2,31,32]. Heating from either the twinned or detwinned martensite phases causes a phase

transformation back to the parent austenite phase [1,2,31–33]. The temperature-stress coupling

between the austenite, twinned martensite, and detwinned martensite structures is illustrated in

Figure 1.3.

There is a driving force associated with the nucleation of martensite, which can be

illustrated using a Gibb’s free energy diagram, as shown in Figure 1.4. The thermodynamic

equilibrium temperature between the two phases is denoted as T0. The starting temperatures for the

forward and reverse transformations are denoted as Ms and As, respectively. The thermodynamic

equilibrium temperature does not equal the starting temperature for the forward and reverse

transformations. The Gibb’s free energy for the system upon the MT has a chemical component

(structural change from parent to product phase), a surface energy component (interfacial energy

between parent and product phases), an elastic energy component, and a non-chemical energy

8

Figure 1.3 Schematic of temperature and stress relationship between the parent and

product phases of NiTi.

component. For the MT to occur, a driving force (denoted ΔTs) is needed to overcome these energy

components. Additionally, for the MT to proceed beyond the starting temperature, further driving

force (cooling in this case) is required [1,34].

Figure 1.4 Gibb’s free energy for the martensitic transformation between the austenite (A)

and martensitic (M) phases. Ms is the martensitic start temperature, or the starting

temperature for the forward (austenite to martensite) transformation; As is the austenite

start temperature (start temperature for reverse transformation).

9

(a) Thermal-Induced MT (b) Superelasticity

(c) Shape Memory Effect (d) Shape Memory Effect – Applied Stress

Figure 1.5 Shape memory behaviors: (a) Thermal-induced martentisitc transformation;

(b) Superelastic behavior due to the stress-induced martensitic transformation; (c) Shape

memory effect behavior where the intial phase is twinned martensite; and (d) Shape

memory effect behavior under an applied stress, where the intial phase is detwinned

martensite. A is the austenite or parent phase; M is the twinned martensitic phase; M+ is

the detwinned martensitic phase.

The macro-scale shape memory behaviors arise from the underlying atomic and micro-

scale MT. The behaviors are schematically shown in Figure 1.5. For the thermal-induced

martensitic transformation (TIMT), the SMA is cooled under no applied load from the high

temperature austenite phase to the low temperature twinned martensite structure (Figure 1.5 (a)).

For the reverse transformation, the temperature is increased, and the material undergoes the reverse

transformation from twinned martensite structure to the austenite phase. The superelastic behavior

is schematically shown in Figure 1.5 (b). At temperatures where austenite is the stable phase, an

10

applied force will cause a stress-induced martensitic transformation (SIMT) and the alloy will

transform from austenite to a detwinned martensite structure. The shape memory effect behavior is

shown in Figure 1.5 (c) and Figure 1.5 (d). For SME, in general, a specimen is deformed at

temperature below the austenite start temperature (material is martensitic) and it regains its original

shape by increasing the temperature above the austenite finish temperature. For SME behavior in

Figure 1.5 (c), the application of force to the initial twinned martensite phase results in detwinning.

Subsequent heating results in strain recovery as the material undergoes the reverse MT. For the

SME behavior in Figure 1.5 (d), the detwinned martensite structure is the initial structure, as an

external load has been applied; strain recovery is accomplished by increasing the temperature above

the austenite finish temperature.

Figure 1.6 Stress-temperature relationship for a material exhibiting shape memory effect

or superelastic behaviors. Ms is the martensitic start temperature; Mf is the martensitic

finish temperature; As is the austenite start temperature; Af is the austenite finish

temperature; Md is the temperature beyond which stress-induced martensite does not form.

Superelastic and shape memory effect behaviors can occur in the same alloy, dependent

upon the temperatures, as long as the reversible MT occurs before permanent plastic

11

deformation/slip [1]. This relationship is shown in Figure 1.6. SME occurs at temperatures below

austenite start temperature (As) before being heated above the austenite finish temperature (Af); SE

behavior occurs above Af. Between As and Af., both SME and SE occur partially. The critical stress

to induce martensite follows the Clausius-Clapeyron relationship. However, above Md, stress

induced martensite can no longer form, and the austenite yields. The critical stress for slip is high

for the material described by Figure 1.6, thus at stresses below this critical slip stress, SME and SE

are realized. If the critical stress for slip is lower, SME or SE behaviors are not realized and

permanent plastic deformation occurs before the MT.

1.3 Connection of Martensitic Transformation to Microstructure

The material response is governed by its microstructure. For SMAs, the MT interacts with

the microstructure, and affects the critical stress for slip, the critical stress for the SIMT, the

transformation temperatures for starting and finishing the MT. By understanding the interaction

between the microstructure and the MT, the shape memory response can be controlled. Specifically,

the microstructure of NiTi alloys can be tailored, or designed, to meet the material requirements

for specific applications.

The transformation temperatures (start and finish temperatures for the forward and reverse

transformations) are related to the NiTi composition. The NiTi phase exists over a small

composition range, from approximately 49.5 to 56 at.% Ni at high temperatures, as shown in the

Ni-Ti phase diagram in Figure 1.7. As previously mentioned, the start temperature of the

martensitic transformation is related to the composition, where a 0.1 at.% increase in Ni content

corresponds to a 10 °C decrease in the transformation temperature, for Ni-rich compositions (Ni

content greater than 50.5 at.%) [35,36]. For Ti-rich compositions (Ni content less than 49.5 at.%),

the martensite start temperature is constant. This behavior is shown in Figure 1.8.

12

Figure 1.7 Ni-Ti equilibrium phase diagram. The equiatomic NiTi phase is highlighted, as

this is the phase which undergoes the martensitic phase transformation and exhibits shape

memory behavior.

Figure 1.8 Relationships between Ni content and martensite start temperature (Ms). Data

is taken from [16].

13

Microconstituent phases form when the alloy composition differs from the equiatomic

composition. Secondary phases, like Ti2Ni and Ni3Ti, and metastable phases which typically form

during aging treatments, like Ni4Ti3 and Ni3Ti2, decrease the volume fraction of NiTi material

which is able to undergo the MT, thus decreasing the amount of recoverable strain. The Ni-rich

secondary phases (Ni3Ti, Ni3Ti2, Ni4Ti3) typically form in Ni-rich alloys, and the Ti-rich secondary

phase (Ti2Ni) typically forms in Ti-rich NiTi compositions. However, as titanium has a high affinity

for oxygen, a Ti-rich oxide (Ti2NiO) can form. This oxide phase is stable at higher Ni

concentrations [1,37].

In superelastic NiTi alloys, however, Ni4Ti3 precipitates are considered desirable. The

Ni4Ti3 phase increases the yield strength of austenite (i.e. increases the critical stress for slip), which

ultimately increases the maximum strain and superelastic recoverable strain. When the precipitates

are small and coherent with the matrix, they create an internal stress field which may help or hinder

the MT (depending on the applied stress field). The lattice strains are generated due to a lattice

mismatch between the precipitate and the matrix crystal structures [37,38].

The grain microstructure also interacts with the MT and dictates the shape memory

behavior. A single crystal NiTi alloy can achieve higher recoverable strains relative to a

polycrystalline material. The polycrystalline material contains grain boundaries which impede the

MT. Additionally, the adjacent grains constrain the maximum deformation strain as compatibility

between grains/across grain boundaries must be maintained [39]. The amount of recoverable strain

varies for the different crystal orientations [40]. Thus, alloys which have grains preferentially

oriented, or textured, will have the highest recoverable strain. The preferred orientation is also stress

state dependent, with the same orientation or texture resulting in different theoretical recoverable

strains for tension relative to compression. This results in an asymmetric tension-compression

response [40–43].

14

The microstructure-MT-property relationships are known and understood from

conventionally processed NiTi alloys. By leveraging this knowledge and by developing

understanding of the microstructure which results from the AM process, AM can potentially be

used for designing the microstructure.

1.4 State of the Art for Additive Manufactured NiTi Shape Memory Alloys

Interest in additive manufactured NiTi alloys has increased in recent years, as shown in

Figure 1.9 (a). Most AM NiTi studies have utilized powder bed fusion (PBF) techniques rather than

laser-based directed energy deposition techniques. However, an advantage of DED is the ability to

blend small quantities of elemental Ni and Ti powders to a specific composition, expanding the

composition range for elementally blended feedstock significantly larger relative to alloyed

feedstocks. The wider range of compositions by using elemental powders is shown in Figure 1.9

(b). In this figure, the utilized feedstock composition is noted for each reference on AM NiTi.

Summary tables of these works are provided in Appendix A and Appendix B.

(a) (b)

Figure 1.9 (a) Summary plot of additive manufactured NiTi references by publishing date,

separated based on AM technique (laser based directed energy deposition (LDED); powder

bed fusion (PBF)) and review articles. (b) Feedstock compositions utilized in these studies,

based on the AM technique and separated based on alloyed powders versus elementally

blended powders.

15

Initial works on AM NiTi can be classified as feasibility studies. These studies focused on

depositing structures using powder bed fusion and laser-based directed energy deposition

techniques, using either alloyed NiTi powders or elementally blended Ni and Ti powders. These

studies confirmed that the fabricated materials exhibited shape memory behavior, primarily by

confirming the thermal-induced martensitic transformation [44–52].

Studies have also focused on relating the AM processing parameters to porosity,

composition, transformation temperatures, impurities, and grain size. The AM processing

parameters which are typically varied include the laser power, scanning speed, hatch spacing, or

layer thickness. These parameters can be summarized into a unifying term of laser energy density

(Equation 1.1) or linear heat input (Equation 1.2).

𝐿𝑎𝑠𝑒𝑟 𝑒𝑛𝑒𝑟𝑔𝑦 𝑑𝑒𝑛𝑠𝑖𝑡𝑦

= 𝐿𝑎𝑠𝑒𝑟 𝑃𝑜𝑤𝑒𝑟

𝑆𝑐𝑎𝑛 𝑠𝑝𝑒𝑒𝑑 ∗ 𝐻𝑎𝑡𝑐ℎ 𝑠𝑝𝑎𝑐𝑖𝑛𝑔 ∗ 𝐿𝑎𝑦𝑒𝑟 𝑇ℎ𝑖𝑐𝑘𝑛𝑒𝑠𝑠

Equation 1.1

𝐿𝑖𝑛𝑒𝑎𝑟 𝐻𝑒𝑎𝑡 𝐼𝑛𝑝𝑢𝑡 = 𝐿𝑎𝑠𝑒𝑟 𝑃𝑜𝑤𝑒𝑟

𝑆𝑐𝑎𝑛 𝑠𝑝𝑒𝑒𝑑

Equation 1.2

The hatch spacing is the distance between adjacent passes (i.e. hatches, tracks) of deposited material

and the layer thickness is the distance between adjacent layers. There is a threshold energy density,

below which parts are no longer fully dense, and porosity increases. For PBF parts, this value is

approximately 55 J/mm3 for fabricating a 99% dense build [50,53,54]. For LDED, the energy

density needs to be at least 50 J/mm3 for fabricating a 90% dense build [45,55]. Increasing the

energy density significantly beyond this value does not guarantee a denser build, rather the

opposite. A decrease in the part density is observed with increasing energy densities beyond the

threshold value [53]. This has been attributed to the balling effect resulting in the formation of

voids or an increase in the melt pool volume resulting in increased gas entrapment [56]. Increasing

energy densities are also correlated to increased Ni evaporation [17,57,58], as Ni has a lower

16

melting temperature than Ti [54,59]. The decreasing Ni content also means decreased

transformation temperatures [16]. The higher energy densities are also correlated with higher peak

temperatures and slower cooling rates [55,57,60–63]. At these higher energies, as the melt pool is

larger, the material stays molten for a longer period of time, and results in increased impurity pick

up [17] and coarser grains [45,55,60,62–64].

Primarily, the focus of AM NiTi studies has been on developing correlations between the

processing parameters and the material response. Few works have explored the fabrication-

microstructure relationships for AM NiTi. Parts built using AM processes are fabricated one layer

at a time to achieve the final part geometry. Specifically, for the LDED AM technique, powder

feedstock is melted and solidified to form a pass, adjacent passes are deposited to form a layer, and

successive layers are deposited to build the part. Multiple adjacent passes and successive layers are

deposited due to the melt pool size being significantly smaller than the final part dimensions. The

adjacent passes and successive layers are also overlapped in order to form an effective bond

between passes and/or layers and to minimize lack of fusion defects [28]. The overlapping of passes

or layers creates these interfacial regions, where previously deposited material is re-melted or re-

heated. The interfacial regions, though localized, represent a significant fraction of the final build.

For example, 25% overlap between passes is typical [28]. Within the interfacial regions, the thermal

histories, and thus microstructures, will be spatially varying when compared to material outside of

these regions. All AM builds have some percentage overlap between passes or layers.

Understanding the influence that the localized re-melting and re-heating at these interfacial regions

has on the microstructure is crucial.

The spatial variation in the grain structure [60,61,65,66] and precipitate morphology

[65,67,68] have been noted in past studies. In terms of the grain structure, the studies attributed the

smaller grain sizes to overlapping tracks resulting in partial re-melting of previously deposited

material or changing heat transfer conditions from the center to the edge of a part (i.e. thermal

17

history). Oliveira et al. [68] only observed Ni4Ti3 precipitates within the heat affected zone (HAZ)

of a Ni-rich alloy, with no precipitates at the center of the pass. The variation in the secondary

phase formation was attributed to the different temperatures of these locations. They postulated that

the HAZ experienced a temperature which was conducive to precipitate the Ni4Ti3 phase (below

600 °C [37]), whereas the temperature and rapid solidification rate at the center of the pass

precluded precipitation.

1.5 Problem Statement

Additive manufacturing has the potential for tailoring shape memory behavior by designing

the microstructure over fine and coarse length scales. By being able to control multiple

microstructure length scales and dimensions, the martensitic transformation and the shape memory

thermo-mechanical response can be tailored. The goal of this research is to develop understandings

of the relationships between the laser-based directed energy deposition AM technique and the grain,

precipitate, and composition microstructure length scales, as well as understand how the additive

manufactured microstructure impacts the shape memory response. To accomplish this goal, NiTi

builds were fabricated using LDED AM. Elemental Ni and Ti powders were blended to control the

feedstock and build compositions; the deposited builds were confirmed to exhibit the fundamental

thermal-induced martensitic response and thus are shape memory alloys. The grains, precipitate,

and composition microstructures were spatially resolved and correlated to the interfacial regions

between passes/layers where localized re-melting and re-heating occurs. To understand the

influence of the additive manufactured microstructure on the MT, the macro-scale shape memory

response and the meso-scale full-field deformation response were characterized for as built and

post-deposition heat treated alloys.

18

The multi-scale characterization of the Ti-rich and Ni-rich microstructures provides insight

into the effect of the localized re-melting and re-heating within interfacial regions for additive

manufactured builds. The composition, precipitate morphology, grain structure, and grain

orientation were investigated, as these microstructure features are known to impact the MT and

SMA behavior. The two different compositions allowed for both SME (for Ti-rich build) and SE

(for Ni-rich build) to be investigated. The influence of the microstructure on the micro-scale MT

and behavior was interpreted from the multi-scale strain analysis approach. Post-deposition heat

treatments were employed to tailor the precipitate morphology to improve the superelastic

recoverable strain.

1.6 Thesis Format and Outline

A general introduction to shape memory alloys and the motivation for using additive

manufacturing to fabricate NiTi SMAs is provided in Chapter 1. Chapter 2 covers the fabrication

of the additive manufactured builds. The initial powder feedstock is characterized and the

methodology for blending the powder feedstocks is discussed. The additive manufacturing

processing parameters and LDED system are detailed. Finally, the post-deposition heat treatments

and sectioning of the build coupons is discussed. Chapter 3 provides a summary of the experimental

approach and the methodology utilized to characterize the microstructure and shape memory

responses. The characterization of the thermal-induced martensitic transformation is discussed in

Chapter 4. The microstructure of the Ni-rich builds and the superelastic response is characterized

in Chapter 5. Chapter 6 focuses on the microstructure and shape memory effect behavior of the

LDED Ti-rich builds and the response for builds fabricated through a powder bed fusion technique.

This work was completed in collaboration with the University of Toledo. Chapter 7 further

characterizes the microstructure and superelastic behavior of the Ni-rich alloys, and employs post-

19

deposition heat treatments for improving the response. The microstructure for the Ti-rich and Ni-

rich builds is further characterized and spatially resolved in Chapter 8. This work is summarized in

the conclusions section in Chapter 9. Chapter 10 briefly discusses the future research and work for

AM NiTi.

20

Chapter 2

FABRICATION OF NITI MATERIALS USING LASER-BASED

DIRECTED ENERGY DEPOSITION

2.1 Chapter Overview

Build coupons were fabricated by additive manufacturing using laser-based directed

energy deposition. Elemental Ni and Ti powders were blended to create the input material (powder

feedstock). Prior to fabricating the final build coupons, powder mass flow rate, laser power, and

scan speed were varied to produce a series of parameter development coupons to minimize lack of

fusion defect. Additionally, the substrate temperature was optimized to prevent delamination of the

builds from the substrate. The post-processing of the final build coupons included removing the

builds from the substrates, sectioning the build coupons for microstructure and thermo-mechanical

characterization, and conducting post-deposition heat treatments to alter the microstructure.

2.2 Laser-Based Directed Energy Deposition Parameter Development

2.2.1 Elemental Ni and Ti Powder feedstock characterization and blending

The composition of NiTi alloys which exhibit shape memory behavior can be classified as

Ti-rich (Ni content less than 49.5 at.%), near-equiatomic (Ni content between 49.5 and 50.5 at.%),

or Ni-rich (Ni content greater than 50.5 at.%). Microstructure and shape memory behavior are very

dependent on composition. Using elemental Ni and Ti powders allows for the blending of powder

feedstock with varying overall compositions. This research utilized Ni and Ti powders to blend two

21

different feedstock compositions. Additionally, elemental Ni and Ti powders are inexpensive

relative to pre-alloyed powders [25,69].

Elemental powders are typically specified, at a minimum, by their particle size,

morphology, and chemical composition. For traditional powder metallurgy processes,

characterization of the powder also includes packing density (or interparticle friction),

compressibility, and the internal structure of the powder (e.g. porosity) [70]. In 2014, ASTM issued

a standard (ASTM F3049 - Standard Guide for Characterizing Properties of Metal Powders Used

for Additive Manufacturing Processes [71]) which provided a starting point for characterizing

powders for AM. The characteristics deemed important for metal AM powders included the particle

size, morphology, chemical composition, flow characteristics, and density. The powder

characterization for this work follows the suggested practices outlined in that standard.

The elemental Ni and Ti powders were purchased from ATI Powder Metals (Pittsburgh,

PA). The manufacturer reports the powders were fabricated by gas atomization from commercially

pure (CP) Ni or Ti ingots. The elemental powders were sieved using a -100/+325 mesh, which

limits the particle sizes to between 45 to 150 μm [70]. The spherical powder (Figure 2.2), with this

range of powder size, has good flowability. Powders need to have good flowability characteristics

to ensure continuous flow of powder during deposition. Small particle sizes, especially below 100

μm in size, tend to agglomerate [70], which may clog the powder nozzle(s) during deposition. The

particle sizes of the Ni and Ti powders were chosen to be similar, in order to minimize particle size

segregation [70], and thus compositional non-uniformity, during handling.

The average particle diameter of the elemental Ni and Ti powders were measured using a

laser diffraction technique. Measurements were taken once when the powders were received (Sept.

25, 2012, termed “when received”) and prior to their use (analysis completed on Aug. 25, 2013

with depositions occurring between Sept. 24 through 27, 2013; termed “prior to deposition”). The

elemental powders were tumbled/re-blended; the prior to use measurements were taken to ensure

22

there was no change. The particle sizes are shown in Figure 2.1, and the calculated metrics are

shown in Table 2.1. The slight variations are likely a result of powder sampling differences

(sampling technique unknown for as received; multiple small samples were taken and combined

for the prior to deposition sample). For the Ni powder, the average particle size prior to deposition

was 101.7 ± 40.7 μm, which was slightly larger than the Ti powder size of 91.4 ± 45.2 μm.

(a) (b)

Figure 2.1 Frequency and cumulative frequency plots for the particle sizes for the

elemental Ni and Ti powders measured (a) when received and (b) prior to deposition.

Table 2.1. Calculated metrics from light diffraction analysis for elemental Ni and Ti

powders. The data from which these metrics are calculated are shown in Figure 2-1. D10

and D90 refer to the particle size which encompasses 10% and 90% of the cumulative

particles.

When Received Prior to Deposition

Ni Ti Ni Ti

Mean (μm) 108.6 91.8 101.7 91.4

Median (μm) 117.5 104.0 110.0 102.9

Standard Deviation (μm) 45.2 46.9 40.7 45.2

D10 (μm) 70.6 59.0 67.8 59.5

D90 (μm) 174.4 166.1 163.5 162.6

Scanning electron microscopy shows that both elemental powders have a spherical morphology.

Figure 2.2 shows the particle geometries of the Ni and Ti powders. The Ni powder appears free of

satellites (Figure 2.2 (a) and (b)), whereas the Ti powder shows some satellites on the particle

surfaces (Figure 2.2 (c) and (d)).

23

(a) (b)

(c) (d)

Figure 2.2 Particle morphologies for (a) and (b) elemental Ni powder, and (c) and (d)

elemental Ti powder. Both powders have a spherical shape. The Ti powder has some finer

satellites attached to the larger particles.

Two feedstock compositions were blended: a Ti-rich feedstock (Ti53.0Ni47.0 wt.%) and a Ni-

rich feedstock (Ni58.0Ti42.0 wt.%). These compositions were chosen because they represent

compositional extremes for NiTi alloys which exhibit shape memory behavior. Additionally, both

shape memory effect (for the Ti-rich) and superelastic (for the Ni-rich) behaviors are anticipated at

room temperature for these compositions. The elemental Ni and Ti powders were dry mixed to

these selected compositions, without additives or lubricants, within a sealed plastic container. The

container was backfilled with Ar gas and sealed prior to convective mixing, to prevent oxidation.

Mixing was carried out in a Type T2C Turbula® Mixer (Willy A. Bachofen AG Maschinenfabrik,

Switzerland). The sealed container was set in three-dimensional movement for 1 hour in order to

achieve random mixing and ensure the elemental powders were evenly dispersed [70]. The blended

powder was then poured into the powder feeder/hopper system.

24

2.2.2 Parameter Optimization

Similar to other fabrication processes, the laser-based directed energy deposition method

has variable process parameters which affect final product properties. The processing parameters

for a powder fed AM technique can be classified as energy source parameters, material parameters,

environmental parameters, manipulation and toolpath parameters, and “other” parameters [28], as

classified by ASTM F3187 (Standard Guide for Directed Energy Deposition of Metals) [72].

Description of Deposition Parameters

The energy source parameters for laser-based directed energy deposition include laser

wavelength, laser power or laser pulse conditions (if the laser is not continuous), spot size at the

part, beam profile, and position of the focal point [72]. The LDED system utilized in this work was

CIMP-3D’s Laser Articulating Robotic System (LARS). This system was custom-built by CIMP-

3D and had an IPG Photonics® YLR-12000 Yb-fiber laser. The laser was delivered through a 200

mm collimator and a 200 mm focal length lens. The laser wavelength was between 1070-1080 nm.

The laser power was initially varied in the parameter development coupons, but the final build

coupons were fabricated using a laser power of 1000 W. The spot size at the part was approximately

4 mm in diameter, and the focus head standoff was 212 mm [48]. Figure 2.3 shows a schematic of

the custom-built LDED system used in this work. The optics for aligning the laser are contained

within the deposition head.

The material parameters include the base material alloy (i.e. substrate), filler material alloy

(i.e. feedstock), mass flow rate, powder capture efficiency, and powder/wire characteristics [72].

The substrate for this work was a commercially pure Ti substrate (McMaster-Carr). Ti was selected

as the substrate material, as to not introduce any unwanted elements into the AM build. Prior to

25

Figure 2.3 Schematic of the custom-built laser-based directed energy deposition system.

depositions, the substrate surface was finished to surface roughness of 0.8 μm (32 μin) arithmetic

average (Ra). Figure 2.4 shows the measured feedrate (or mass flow rate) for the Ti-rich powder

feedstock, which was determined experimentally by flowing the powder for 60 s and measuring

the weight of the powder captured in a container. The feedrate therefore defines the amount of

material being provided through the nozzle, and not necessarily the amount of material entering the

melt pool. The independent variable for these experiments was voltage supplied to the outflow

mechanism of the powder hopper. Increasing the voltage increased the feedrate. The feedrate was

measured for the Ti-rich feedstock and assumed similar for the Ni-rich feedstock. A voltage reading

of 40 V was selected, which corresponded to a powder feedrate of 12 g/min. [48]. The feedrate of

the powder and the parameters which determine the energy directed to the melt pool are related:

higher energy corresponds to a larger amount of material which can be melted (higher feedrate).

The environmental parameters include chamber gas composition, supplemental gas

composition and flow rate, nozzle orifice geometry and diameter, and vacuum level [72]. The

chamber gas is the atmosphere within the sealed enclosure surrounding the LDED system. In this

work, the chamber gas was Ar. The inert Ar environment was necessary to minimize oxygen pick-

26

Figure 2.4 Measurement of powder feedrate, or mass flow rate of the feedstock powder,

based on the voltage (dial) reading of the powder feeder. This mass flow rate was measured

for the Ni46.9Ti52.1 at.% powder.

up, especially with the elemental Ti powder [7,73]. The supplemental gas, which assists powder

flowability and protects the optics within the deposition head, was also Ar. The Ar supplemental

gas flow rate was 9.4 L/min. [48]. The nozzle orifice, where the powder flowed out toward the melt

pool, was a ring. However, the small opening of the ring nozzle did occasionally result in powder

clogging the orifice. This prevented powder from flowing into the melt pool during one parameter

development build, ultimately resulting in a failed build. To prevent this from reoccurring, the

nozzle orifice was periodically cleaned using a wire brush. This step mitigated powder clogging.

The manipulation and toolpath parameters include the travel speed (or scan speed), layer

height, hatch spacing, and mechanical arrangement [72]. The travel speed is measured as the rate

at which the deposition head (containing the laser and powder nozzles) moves relative to the

substrate. The travel speed was varied during the parameter development stage and was set at a

constant 10.6 mm/s when depositing the final build coupons. Figure 2.5 (a) shows the dimensions

of the pass width and the layer thickness for a single deposited pass, determined from the process

parameters and the resulting bead geometry. Figure 2.5 (b) shows two adjacent passes, the centers

of which are separated by the hatch spacing. Since the hatch spacing is less than the pass width, the

two adjacent passes are overlapped. Figure 2.5 (b) also shows two successive layers, the tops of

27

which are separated by the layer height. Since the layer height is less than the layer thickness, the

layers are overlapped. Passes and layers are typically overlapped in order to minimize lack of fusion

defects [28]. For this work, the pass width was 3.6 mm and the hatch spacing was 1.9 mm. The

layer thickness was 1.0 mm and the layer height was 0.3 mm.

(a) (b)

Figure 2.5 Schematic of the (a) dimensions for a single pass (b) the dimensions for the hatch

spacing and layer thickness deposition parameters. The region of overlap between adjacent

passes and the interlayer region due to the deposition of successive layers are highlighted.

Other parameters include the substrate temperature and the dwell time between passes and

layers [72]. For this work, the substrate was bolted to a ThermoScientific HP131225Q hot plate.

For depositing the build coupons, the hot plate temperature was set to 250 °C. The actual substrate

temperature was measured and recorded prior to deposition, by measuring the temperature of the

substrate near the starting point for the deposition. The temperature was measured using a Fluke

meter temperature probe. The dwell time is the time between individual passes and layers during

which the deposition head is stationary. The final build coupons were fabricated with no dwell time

between passes and layers. The processing parameters used for depositing the build coupons are

listed in Table 2.2.

The goal of the parameter development builds was to minimize process-related defects.

Depositing dense structures required that (i) single passes had uniform widths and thicknesses along

the entire continuous track length (ii) adjacent passes and successive layers were overlapped with

no gaps or porosity; and (iii) that there was no cracking or delamination between passes, layers, or

in the overlap regions. Large build coupons were deposited, to take advantage of LDED’s fast

28


the laser-based directed energy deposition technique.

Parameter Value

(Laser) Energy

Source

Parameters

Laser wavelength 1070 nm

Laser power 1000 W

Spot size (at part) 4 mm

Position of focal point 212 mm

Material

Parameters

Base Material Alloy CP Ti

Filler Material Alloy elementally blended Ni and Ti powders

Feedrate 12 g/min.

Powder characteristics described in 2.2.1

Environmental

Parameters

Chamber Gas Ar

Supplemental Gas/Flow rate Ar; 9.4 L/min.

Vacuum Level N/A

Manipulation

and Toolpath

Travel speed 10.6 mm/s

Layer height 1.0 mm

Hatch spacing 1.9 mm

Other Temperature 250 °C (set on hot plate beneath substrate)

Dwell Times 0 s between passes, 0 s between layers

deposition rates [74] for fabricating large structures, as might be a typical need in an industry

application.

The first parameter which was varied for this parameter development exercise was the mass

flow rate of the powder. A series of single pass, single layer tracks were deposited and evaluated

to determine the mass flow rate that produced the most uniform dimensions in the deposited

structure. Any fluctuations in width or height of a pass may result in gaps or porosity when an

adjacent pass or a successive layer is deposited. A series of four single pass builds were deposited

with mass flow rates ranging from 5.0 g/min. to 10.2 g/min, as shown in Figure 2.6 (a). The width

and height of the builds are reported in Figure 2.6 (b) and (c), respectively. Similar trends are

observed for both the width and height dimensions: the build dimensions (width and height) were

the smallest for the lowest mass flow rate, and increased to maximum with increasing mass flow

rate until 8.1 g/min. Further increasing the flow rate to 10.2 g/min. resulted in a smaller width and

height. However, the width and height dimensions for the 10.2 g/min. flow rate build had the

smallest standard deviation. That is, the 10.2 g/min. flow rate build had the most uniform

29

dimensions along the entire deposited pass. Visual inspection of the deposited passes revealed that

there was sufficient energy necessary to fully melt the powder for all tested powder flow rates.

(a)

(b) (c)

Figure 2.6 (a) Single pass, single layer builds fabricated with increasing mass flow rates.

Dimensions for the (b) width and (c) height for these builds fabricated using varying mass

flow rates. Builds were fabricated with 1000 W laser power, a scan speed of 10.6 mm/s, and

a substrate temperature of 25 °C. The build deposited using the 10.2 g/min. mass flow rate

had the smallest standard deviation in the width and height. The scale bar in (a) is in mm.

(Build coupon ID: B2 (5.0 g/min.), B3 (6.1 g/min.), B4 (8.1 g/min.), B5 (10.2 g/min.))

The number of passes and layers were increased, and passes and layers overlapped, to

fabricate larger coupons. The overlap between the passes was set at 40% (calculated as ratio

between hatch spacing and actual pass width), and the overlap between layers was set at 60%

(calculated as ratio between layer spacing and actual layer height). Figure 2.7 shows the two

fabricated parameter development coupons with large overall dimensions (1 pass, 6 layers, with

build dimensions of 4 mm width and 2.5 mm height). The coupons delaminated from the substrate.

30

This delamination ruins the build, as the distance between the build and the deposition head (laser

and powder nozzle) is no longer constant throughout the fabrication process. Additionally, in

normal operation the excess heat generated by the laser is conducted through the previously

deposited material into the substrate. A coupon that has delaminated from the substrate will have

less area through which heat can be conducted to the substrate, forcing higher heat flux through the

rest of the part that remains attached to the substrate. The next section of parameter development

focused on reducing thermal stresses in order to deposit a build coupon which remains attached to

the substrate throughout fabrication.

Figure 2.7 Build coupons deposited with 1 pass and 6 layers, using 1000 W laser power,

10.2 g/min. mass flow rate, scan speed of 10.6 mm/s, and a substrate temperature of 25 °C.

The builds delaminated from the substrate as pointed out with the red arrows. The small

ticks on the scale bar are 1 mm. (Build coupon ID: B6, B7)

It was hypothesized that during deposition high thermal stresses were causing the build

coupons to delaminate from the substrate. These stresses may have resulted from the large thermal

gradients building up within the deposit [74–77]. The high temperature at of the melt pool resulted

in the deposited structure experiencing more thermal expansion than the substrate, giving rise to

the stress [74–77]. It was also hypothesized that a brittle Ti-rich (i.e. Ti2Ni) phase was formed at

the interface between the Ti substrate and the NiTi build, which failed due to the presence of

residual stress. The processing parameters of dwell time, laser power, and substrate temperature

were varied in order to fabricate larger build coupons which remained adhered to the substrate. The

dwell time between layers was increased to 3 s (from 0 s / no dwell time). A parameter development

build coupon was again fabricated with 1 pass and 6 layers. The build coupon is shown in Figure

31

2.8 (a). The height of the build coupon is not uniform, with a difference of 3 mm between the lowest

and highest points. This build coupon also delaminated from the substrate. In Figure 2.8 (b), the

build coupon was deposited with a lowered laser power (500 W compared to 1000 W), to decrease

the heat input (amount of energy directed to the melt pool). We postulated that the lower power

would decrease the maximum preheat temperature observed in the build and decrease the cooling

rate, thus decreasing the thermal stress. The build height for this coupon was uniform. However,

it still delaminated from the substrate, as the cooling rates were still high.

(a)

(b)

Figure 2.8 Build coupons deposited with 1 pass and 6 layers, using (a) 1000 W laser power

and a dwell time of 3 s between layers; and (b) 500 W laser power and a dwell time of 0 s

between layers. Coupons were deposited using 10.2 g/min. mass flow rate, scan speed of

10.6 mm/s, and a substrate temperature of 25 °C. The builds delaminated from the

substrate as pointed out with the red arrow. The small ticks on the scale bar are 1 mm.

(Build coupon ID: B8 and B12)

For the next parameter development coupon, the substrate was pre-heated. This was in

order to further decrease the cooling rate. The coupon was again deposited with 1 pass and 6 layers,

as shown in Figure 2.9. The temperature of the hot plate (at the base of the substrate) was set to

250 °C. The surface of the substrate, however, was above 250 °C due to conduction from the melt

pool. The measured temperature at the substrate was approximately 350 °C for all builds. This pre-

heating of the build plate proved to eliminate all delamination issues.

32

Deposition of Builds

Build coupons with varying numbers of passes and layers were deposited using the

optimized deposition parameters. Identical build coupons were fabricated using each of the two

powder feedstock compositions discussed previously. Figure 2.9 shows example build coupons

fabricated using 14 layers and (a) 1 pass, (b) 3 passes, or (c) 6 passes. All the build coupons

remained attached to the substrate during deposition. A table listing the build coupons, deposition

parameters, and build plans (i.e. number of layers and passes) is in Appendix A.

(a) (b)

(c)

Figure 2.9 NiTi build coupons fabricated using the laser-based directed energy deposition

technique with 14 layers and (a) 1 pass; (b) 3 passes; and (c) 6 passes.

33

2.3 Build Coupon and Specimen Preparation for Microstructure and Thermo-Mechanical

Characterization

2.3.1 Removal of build coupons from substrates

Build coupons were removed from the substrate using wire electrical discharge machining

(EDM) (Acura-Cut, Pleasant Gap, PA). The material which remained on the substrate had heights

between 0.2 and 0.4 mm. Figure 2.10 (a) shows the substrates after the build coupons have been

removed using EDM. The removed coupons are shown in Figure 2.10 (b). A cross section of the

substrate (and the portion of the build coupon below the EDM sectioning line) is shown in Figure

2.10 (c). The dilution region is adjacent to the substrate. In the dilution region, additional Ti from

the substrate is mixed with the NiTi alloy in the melt pool. The dilution region was further

investigated for a 6 pass, 14 layer build fabricated from the Ti-rich powder, shown in Figure 2.10

(d); the thickness of the dilution region was approximately 0.15 mm. Thus, the removed coupons

are entirely above the dilution region.

(a) (b)

(c)

Figure 2.10 (a) Build coupons after being removed from the substrate. (b) Cross-section of

substrate, dilution region, and build coupon after the build coupon has been removed from

the substrate. (c) Optical microscopy image of the cross-section of a 6 pass, 14 layer build

coupon fabricated from the Ti-rich powder blend. The dilution region is labeled.

34

2.3.2 Sectioning of builds into specimens

The large build coupons were sectioned into smaller microstructure samples, differential

scanning calorimetry (DSC) samples, and thermo-mechanical test specimens. The recorded

sample/specimen details include build coupon ID, specimen ID, location from which the

sample/specimen was extracted, material condition (as built, heat treated, etc.), and the type of

characterization performed. The locations of the extracted specimens are identified based on the

XYZ additive manufacturing coordinate system, with the origin point of the build coupon

representing (0,0,0). The locations are measured to the geometric center of the specimen.

Microstructure and DSC samples were sectioned using a slow speed saw with a diamond

tip blade, to provide a clean cut with minimal deformation. Thermo-mechanical test specimens

were sectioned from the larger build coupons using wire EDM to the specified dogbone or

compression block geometries. The dogbone and compression block geometries and dimensions

are shown in Figure 2.11 and are based on the dimensions for small-scale specimens [78]. The

EDM surface layer was removed by light abrasion prior to post-deposition heat treatment or

thermo-mechanical testing.

(a) (b)

Figure 2.11 (a) Dogbone and (b) compression block geometries used in this work.

Dimensions are from [78].

35

2.3.3 Post-deposition heat treatments

Post-deposition heat treatments were utilized to modify the microstructure of a subset of

extracted samples, specifically the builds fabricated from the Ni-rich composition. Thermo-

mechanical treatments are common for conventionally processed Ni-rich alloys, to bring about the

optimal microstructure and the desired properties [2,37]. Heat treatments included direct aging

treatment (precipitation aging heat treatment on as built material), solutionizing treatment, and

solutionizing treatment followed by an aging heat treatment.

The purpose of the solutionizing treatment was to homogenize the composition and

dissolve unwanted secondary phases. For NiTi alloys, solutionizing occurs above 850 °C, as shown

on the Ni-Ti phase diagram [1,37]. The solutionizing heat treatments require an inert environment,

due to the high temperatures and long times of these treatments, to minimize oxidation. The inert

environments within the furnace were variously created by (i) flowing inert Ar gas over the

specimen for the duration of the heat treatment; (ii) encapsulating the specimen in a quartz tube

which has been vacuumed sealed; or (iii) encapsulating the specimen in a quartz tube which has

been backfilled with Ar gas. The inert Ar environment was created by flowing ultra-high purity Ar

gas, at a rate of 10 cm3/min., through either a RapidTemp tube furnace (CM Furnaces, Bloomfield,

NJ) or a 1700 BLF (bottom loading furnace) (Carbolite, Hope Valley, UK). Heat treatment of the

specimens in quartz tubes was completed in a Lindberg Blue M furnace (Thermo Fisher Scientific,

Waltham, MA).

The effectiveness of the inert environments was assessed based on the degree of oxide or

scale formation on the surface of the specimen. A more perfectly inert environment would result in

less scale or oxide formation. Compression specimens were heat treated in the three different inert

environments at 1000 °C for 36 hours. For the specimens in the flowing Ar gas atmosphere, visual

inspection of the surface revealed a noticeable thickness of a dark gray oxide layer on the sides, as

36

shown in Figure 2.12 (a). As the scale does not undergo the reversible MT, it had to be removed.

Subsequent grinding on all size sides using SiC sandpaper successfully removed this oxide layer.

Figure 2.12 (b) shows the size difference between a specimen which has not been heat treated at

all and specimens which have been solutionized in Ar with the unwanted oxide layer removed. The

oxide layer was thick, as removal of this oxide layer significantly reduced the specimen dimensions.

(a) (b)

Figure 2.12 (a) Compression specimens heat treated in an Ar environment. One side

has been polished to reveal the NiTi material. The dark gray boundaries on the edge

are the surface oxide layer. (b) Shown are two of the polished compression specimens

next to the original 8 x 4 x 4 mm compression specimen geometry.

Figure 2.13 Change in thickness (Δt) between specimen after solutionizing heat treatment and

after grinding to remove the oxide scale layer. Specimens were solutionized in a flowing Ar

environment (flowing Ar), encapsulated in a quartz tube with a vacuum environment

(vacuum quartz tube), or encapsulated in a quartz tube with a back-filled Ar environment

(Ar quartz tube).

A surface oxide layer was also present on both the specimens that were heated treated in vacuum

and back-filled Ar environments in the sealed quartz tubes This surface layer was also removed by

37

grinding. The change in thickness (Δt) between the heat treated specimen dimensions before and

after removing the surface oxide layer for the three heat treatment environments are shown in

Figure 2.13. The thickness of the oxide layer in a backfilled Ar quartz tube was the smallest, and

thus represents the most effective inert environment for minimizing oxidation.

For the quartz tubes, a film on the inside of the tube was observed only for the vacuum

environment. The quartz tubes (before (a) and after ((b) and (c)) heat treatment) are shown in Figure

2.14. The film was deposited on the tube surface near where the specimen was located (Figure 2.14

(b)); no film was observed for the Ar back-filled sealed quartz tubes (Figure 2.14(c)). The back-

filled Ar sealed quartz tubes were therefore utilized for all subsequent high temperature

solutionizing heat treatments.

(a)

(b) (c)

Figure 2.14 (a) Specimens which have been encapsulted in quartz tubes. (b) Vacuum sealed

quartz tubes after a 1000 °C, 24 h heat treatment. A film has formed on the inside of the

quartz tube. (c) Back-filled Ar sealed quartz tubes after a 1000 °C, 24 h heat treatment.

The purpose of precipitation aging treatments was to precipitate (or grow) the desirable

Ni4Ti3 phase. The lower temperature (450 to 550 °C) precipitation aging treatments were conducted

in Lindberg Blue M furnace. At this lower temperature, minimal surface oxidation was observed.

The thin surface layer was removed through grinding with SiC sandpaper prior to further

38

characterization. Depending on the exact aging temperature and time, desirable Ni4Ti3 precipitates

or undesirable Ni3Ti2 or Ni3Ti phases form and grow [1,2,37].

2.4 Fabrication of Powder Bed Fusion builds

Additional compression specimens were fabricated using the additive manufacturing

powder bed fusion technique at the Nitinol Commercialization Accelerator group within the

Dynamic and Smart Systems Lab at the University of Toledo. These specimens were used for a

limited comparison to the LDED samples, as discussed in subsequent sections. Specimens were

fabricated on a Type PXM selective laser melting (SLM) commercial workstation (3DSystems,

Rock Hill, SC). Specimens were fabricated in an inert Ar gas environment. An alloyed powder with

a Ni50.09Ti49.92 at.% composition and a powder size of 25-75 μm was employed. The small powder

size used in this technique was appropriate for PBF techniques, where the powder is spread using

a re-coater blade and the layer thickness is on the order of the powder diameter [74]. For these

depositions, laser power was 250 W, laser beam diameter was 0.08 mm, pass width was 0.20 mm,

scan velocity was 1250 mm/s, hatch spacing was 0.12 mm, and layer thickness was 0.03 mm

[53,79]. The calculated laser energy density was 56 J/mm3 (Equation 2.1), which is similar to the

energy density used for depositing the LDED build coupon (60 J/mm3). However, the linear heat

input for PBF alloys was significantly smaller to that for the LDED alloys (0.2 J/mm for PDF

relative to 94 J/mm for LDED) Individual compression specimens were fabricated to the same

dimensions shown in Figure 2.11 (b). The compression loading direction aligned with the z-

direction (build height direction). Specimens were deposited on a porous support structure

fabricated using the same NiTi feedstock. The support structures were fragile and easily broke apart

from the specimens using minimal force [79].

39

2.5 Chapter Summary

Build coupons were fabricated with the LDED AM technique using two feedstock

compositions. Processing parameters were determined by fabricating a series of parameter

development builds to manufacture builds free of defects. However, subsequent build coupons

consisting of multiple passes and layers delaminated from the substrate. A substrate pre-heat was

applied to decrease the impact of large thermal excursions experienced by the deposited material

[48]. Build coupons were removed from the substrate, above a dilution region, and further sectioned

into microstructure and thermo-mechanical specimens. For the Ni-rich alloys, post-deposition heat

treatments were employed for tailoring the as built microstructure.

40

Chapter 3

METHODOLOGY FOR CHARACTERIZATION OF

MICROSTRUCTURE AND

SHAPE MEMORY BEHAVIOR


The microstructure and thermo-mechanical responses of the NiTi additive manufactured

builds were characterized. The additive manufactured microstructure was characterized based on

the composition (using energy dispersive spectroscopy), phase (using x-ray diffraction), precipitate

morphology (using scanning electron microscopy and X-ray dispersive spectroscopy), and grain

structure (using electron-back scatter diffraction technique within SEM or optical microscopy). The

shape memory behavior was investigated based on the thermal-induced martensitic transformation

(using differential scanning calorimetry), the stress-induced martensitic transformation and the

shape memory effect (both using thermo-mechanical experimentation). A full-field deformation

analysis technique of digital image correlation (DIC) was employed to correlate the additively

manufactured microstructure to the martensitic transformation behavior.

3.2 Microstructure Characterization

3.2.1 Sample Preparation

The specific microstructure characterization technique dictates the necessary sample

preparation. For example, detected backscattered electrons for the EBSD technique come from the

first 50-100 nm depth of a specimen [80]. Any deformation or distortion on the specimen surface

will complicate data interpretation. Therefore, polishing processes must impart little to no

41

deformation on samples that are to be analyzed with EBSD. Polishing processes that meet this

criteria include vibratory polishing and electropolishing [80,81].

Microstructure samples were sectioned and mounted in epoxy. A clear epoxy was used for

samples which had to be removed from the epoxy, as heating the epoxy made it malleable. The

epoxy resin and hardener (EpoThinTM 2 Epoxy Resin and Hardener, Buehler, Lake Bluff, IL) were

mixed in a 100:45 weight ratio and had a 10 hour curing time at room temperature. A black

DuroFast carbon compound (Struers, Cleveland, OH) was also used with a ProntoPress-20

mounting press (Struers). The recipe was 30 MPa pressure and 180 °C temperature for 4 minutes,

with a 3 minute cooling period.

Grinding and polishing were accomplished by using successively finer grits of SiC

sandpaper. For NiTi alloys, an undesired martensitic phase transformation can occur if a large force

is applied during grinding. To minimize the grinding force, 600 grit (P1200, 15 μm particle size)

sandpaper was the coarsest grit utilized. Successively finer grits of sandpaper were utilized to polish

the surface to a final 1200 grit (P4000, 5 μm particle size). Between each grit, samples were cleaned

by ultrasonicating in a water bath and rinsing with acetone to minimize contamination. Mounted

specimens were polished on a MetPrep 3TM (Allied High Tech Products Inc., Rancho Dominguez,

CA) or a RotoForce 4 (Struers) automatic polisher, with minimal applied force (3 lbs on individual

mounts for the MetPrep; 30 lbs distributed over six mounts for the RotoForce). Water was used as

a lubricant. The samples were polished using 3 μm and 1μm diamond slurry (martensitic materials)

or colloidal silica (austenitic materials).

A PACE Technologies GIGA-0900 vibratory polisher was used with a NAPPAD polishing

pad as the final polishing step. Three polishing media were trialed, to determine which created the

best sample surface for EBSD: (i) colloidal silica, (ii) diamond, and (iii) chemical-mechanical

polish (CMP) slurry, a media composed of silica and alumina. Samples from the same build coupon

were polished using the previously described methodology prior to final vibratory polishing. One

42

sample each was vibratory polished for 2 hours using one of the vibratory polishing media. Then,

the zero solutions (pixel areas where the EBSD software is unable to determine the phase and

orientation of the material due to a low signal or inability to determine the Kikuchi lines) were

analyzed. After the vibratory polish time (2 h), the colloidal silica had the lowest zero solution

percentage (20 %), compared to the diamond (73%) or CMP slurry (51%). Thus, 0.04 μm colloidal

silica was chosen as the vibratory polishing media for all subsequent samples.

An additional step of chemical etching was necessary to reveal the grain boundaries, grain

facets, and additive manufacturing features for optical microscopy. Samples were etched using

Kroll’s reagent (10 vol.% HNO3, 5 vol.% HF, remainder water) for between 15 and 150 seconds.

This etchant is preferred for NiTi alloys [60,82–85] and suggested for Ti-alloys in ASTM E407

(Standard Practice for Microetching Metals and Alloys) [86]. For Ti alloys, the HF acid attacks the

surface of Ti alloys while HNO3 brightens the surface of Ni alloys [87]. Kroll’s reagent was

swabbed on the sample surface, rather than the sample being immersed, to reduce staining [87] and

to have greater control over the etching process. The quality of the polished sample surface can

influence the etching response [81,87] and thus a polished surface free of visible scratches was

confirmed using optical microscopy prior to etching. As NiTi readily forms a passivation layer on

the surface of the sample [9], etching was completed immediately after fine polishing, so that the

passivation layer did not have time to form.

3.2.2 Microstructure characterization methods

The composition and secondary phase morphology characterized using scanning electron

microscopy (SEM), specifically the techniques of energy dispersive x-ray spectroscopy (EDS) and

backscatter electron (BSE), respectively. In addition to the visualization of secondary phases and

precipitates, the area fraction, size (length, width/major and minor axis dimensions), and spacing

43

of the secondary phases were measured. The compositions of the phases were identified by

comparing the measured composition to the calculated Ni:Ti ratio for different phases (Table 3.1).

The SEMs utilized in this work are: a Philips XL30 environmental SEM (ESEM)

(FEI/ThermoFisher Scientific), a Quant 200 ESEM (FEI/ThermoFisher Scientific), a Quant 250

ESEM (FEI/ThermoFisher Scientific), and a Helios NanoLab 660 focused ion beam (FIB)/SEM

(FEI/ThermoFisher Scientific). X-ray diffraction (XRD) augmented the phase identification,

especially for identifying the austenite, martensite, and R- NiTi phases as the composition of these

phases is identical. Table 3.2 lists the crystal structure details for the austenitic and martensitic NiTi

phases and the secondary phases identified using XRD in this work.

Table 3.1 Calculated Ni:Ti composition ratios for NiTi phases and secondary phases used

with EDS to identify phases.

Phase Ni:Ti content

(at.%) (wt.%)

B2 NiTi, B19' NiTi, R-phase NiTi 50.0:50.0 55.1:44.9

Ti2Ni 33.3:66.7 38.0:62.0

Ni4Ti3 57.1:42.9 62.0:38.0

Ni3Ti 75.0:25.0 78.6:21.4

Table 3.2 Crystal structures and parameters for the NiTi phases and secondary phases

identified using x-ray diffraction. The PDF# corresponds to the inorganic crystal structure

database (ICSD) number, or the reference crystal structure utilized to identify the phases

in the additive manufactured alloys.

Phase Crystal System a; b; c (Å) α; β; γ (°) PDF#/ICSD#

B2 NiTi cubic 3.015; 3.015; 3.015 90; 90; 90 04-004-9090

B19' NiTi monoclinic 2.92; 4.725; 4.031 90; 90; 98 97-016-6012

R-phase NiTi hexagonal 7.257; 7.257; 5.383 90; 90; 120 97-015-7605

Ti2Ni cubic 11.333; 11.333; 11.333 90; 90; 90 04-003-6277

Ni4Ti3 hexagonal 11.2632; 11.2632; 5.0969 90; 90; 120 01-078-4623

Ni3Ti hexagonal 5.0924; 5.0924; 8.2975 90; 90; 120 00-051-1169

The grain structure was visualized using optical microscopy and electron backscatter

direction (EBSD). An Olympus MX50 optical microscope (Shinjuku, Toyko, Japan) and a Stemi

508 stereoscope (Carl Zeiss AG, Germany) were employed. The grain size was measured using

ASTM E112 (Standard Test Methods for Determining Average Grain Size) [88]. The morphology

44

(shape of the grains) was characterized by the aspect ratio (ratio of length to width). EBSD was

utilized to visualize the grain structure and determine the grain orientation. A Helios NanoLab 660

FESEM (FEI, Waltham, MA) equipped with a NordlysMax2 EBSD detector (FEI), and AZtecHKL

software (Oxford Instruments, Abingson, Great Britain) was used. The austenitic phase and the

Ni4Ti3, Ni3Ti, Ti2Ni secondary phases were able to be identified using EBSD. The martensite

variants could not be identified, as a program-specific crystal structure file could not be created for

the B19' or R phases. Post-processing of the EBSD data was completed in MapStitcher, Tango, and

Mambo (Oxford Instruments) software programs, allowing for very large areas (18 mm2) to be

analyzed as a single area.

3.3 Thermo-mechanical Characterization

3.3.1 Thermal-Induced Martensitic Transformation, characterized by Differential

Scanning Calorimetry

The MT results in measurable changes in the hardness, yield strength, Young’s modulus,

heat capacity, latent heat of transformation, lattice spacing, and thermal conductivity between the

austenitic and martensitic phases [2]. Differential scanning calorimetry measures the heat flow

necessary to maintain a constant temperature change rate, relative to a reference sample. When an

endothermic phase transformation occurs, additional heat is necessary to maintain the constant

temperature change; for an exothermic phase transformation, additional heat removed. The reverse

MT (martensite to austenite) is an endothermic reaction and the forward MT (austenite to

martensite) is an exothermic reaction. From this experiment, the temperatures at which the forward

and reverse MT begin and end, as well as the latent heat of transformations, can be measured.

45

The calorimeter is a DSC 8500 (Perkin-Elmer, Shelton, CT). The measured sample of

known weight is placed in an aluminum pan within the sample chamber. The reference chamber

contains an empty aluminum pan. In DSC, the heat flow to maintain a constant heating/cooling rate

in the sample chamber is always compared to the reference chamber. The experiment conformed

with ASTM F2004 (Standard Test Method for Transformation Temperature of Nickel-Titanium

Alloys by Thermal Analysis) [89]. The experiment details are as follows: (i) Heat sample from 23

°C to 100 °C, at a rate of 10 °C/min. (ii) Hold at 100 °C for 2 minutes to allow the sample to

equilibrate. (iii) Cool sample from 100 °C to -120 °C (minimum temperature) at a rate of 10 °C/min.

(iv) Hold at -120 °C for 2 minutes to allow sample to equilibrate. (v) Heat sample from -120 °C to

100 °C at a rate of 10 °C/min. (vi) Return sample to 23 °C. The maximum temperature for the DSC

experiment should exceed the finishing temperature for the reverse MT (martensite to austenite) by

at least 30 °C (Af + 30 °C) and the minimum temperature should exceed the finishing temperature

for the forward MT (austenite to martensite) by at least 30 °C (Mf – 30 °C) [89].

The transformation and peak temperatures of the MT, as well as the enthalpies of

transformation, are calculated, following ASTM F2004 [89]. An example DSC scan is shown in

Figure 3.1 and shows the described temperatures and enthalpies. The positive endothermic peak

corresponds to the reverse MT. The negative exothermic peak corresponds the forward MT. The

start and finish temperatures for the forward and reverse transformations are determined using the

tangent method [89] and are identified as the martensitic start and finish temperature (Ms, Mf) for

the forward transformation and the austenitic start and finish temperatures (As, Af) for the reverse

transformation. The peak temperatures are the maximum or minimum temperatures of the peaks

(Mp, Ap). The forward and reverse enthalpies (ΔHF, ΔHR) are calculated by integrating under the

curve. A temperature hysteresis exists between the forward and reverse MTs and is calculated as

the difference between the austenite finish and martensite start temperatures (𝐴𝑓 − 𝑀𝑠) or the

difference between the peak temperatures (𝐴𝑝 − 𝑀𝑝) [90]. There is some ambiguity in the start

46

and finish temperatures determined using the tangent method, therefore the peak temperatures are

utilized where appropriate, and the thermal hysteresis is calculated from the peak temperatures.

Figure 3.1 DSC thermogram for an additive manufactured Ti-rich specimen, which shows

a single forward transformation peak and a single reverse transformation peak. The

martensitic start (Ms) and finish (Mf) transformation temperatures for the forward

martensitic transformation and the austenitic start (As) and finish (Af) transformation

temperatures for the reverse transformation, as well as the martensitic peak (Mp) and

austenitic peak temperatures (Ap), are identified. The integrated regions, identified in blue

and red, correspond to the enthalpies for the forward (ΔHF) and reverse (ΔHR)

transformations, respectively.

3.3.2 Pseudoelastic and Superelastic Behavior, characterized by an Isothermal

Mechanical Experiment

The pseudoelastic and superelastic behaviors in the additive manufactured alloys were

characterized with an isothermal mechanical experiment. For pseudoelastic behavior, the stable

phase is the austenite phase, which undergoes a stress-induced martensitic transformation when

stress is applied. The reverse transformation occurs when the stress is removed. The criteria

necessary for a pseudoelastic response is a testing temperature above the austenite finish

temperature (Af) and a high yield strength such that the alloy is able to undergo the SIMT prior to

inducing permanent strain [91]. Additionally, there is a temperature criterion for a complete

pseudoelastic response, which was calculated using Equation 3.1 .

47

𝑇 > 𝐴𝑓𝑑 + 𝜎0

𝑆𝐼𝑀𝑑𝑇

𝑑𝜎 Equation 3.1 [91]

where 𝑻 is the temperature criterion for complete pseudoelasticity; 𝑨𝒇𝒅 is the finish temperature

of the first reverse transformation after deformation; 𝝈𝟎𝑺𝑰𝑴 is the minimum stress for stress-

induced MT at Ms; and 𝒅𝑻

𝒅𝝈 is the inverse slope of linear dependence of the critical stress for

stress-induced martensite on the testing temperature.

Thermo-mechanical experiments were conducted on the MTS 810 servo-hydraulic load

frame. Figure 3.2 (a) shows the entire thermo-mechanical set-up. The load cell, which measured

the force, had a 20 kN capacity. The temperature was controlled using a custom-built induction

coil, which required a control system (labeled in Figure 3.2 (a)). The temperature of the specimen

was measured using a T-type thermocouple, which was spot welded to the side of a specimen using

a HotSpot TC thermocouple welder (DCC Corporation, Pennsauken, NJ). The specimen was heated

to the test temperature and allowed to thermally equilibrate (approximately 10 minutes) prior to

starting the experiment. To accommodate the small dogbone and compression block geometries,

additional adapters and tension fixtures or compression platens were added to the load train,

between the load cell and the displacement linear variable differential transformer. The load train

is identified by the red box in Figure 3.2 (a). Close-up views of tension and compression specimens

within the fixtures/platens are shown in Figure 3.2 (b) and Figure 3.2 (c), respectively.

The thermo-mechanical experiments for characterizing the superelastic behavior followed

ASTM standards E8 (Standard Test Methods for Tension Testing of Metallic Materials) [92], E9

(Standard Test Methods of Compression Testing of Metallic Materials at Room Temperature) [93]

and F2516 (Standard Test Method for Tension Testing of Nickel-Titanium Superelastic Materials)

[94]. The specimen was gripped within the fixture (tension specimen) or fixed within the platen

(compression specimen) and a small pre-load stress (less than 10 MPa) is applied. The tension

specimens were loaded in displacement control at a rate of 0.15 mm/min. [92], which corresponds

to a quasi-static strain rate of 1x10-3 s-1. To conform to the compression testing standard, a lubricant

48

(a)

(b) (c)

Figure 3.2 (a) Overview of the thermo-mechanical experimental set-up. Close-up figures of

(b) tension and (c) compression specimens within the load frame.

(Enerpac Hydraulic Oil, Milwaukee, WI) was applied to the platens to prevent barreling of the

specimen during loading. The rate of loading for the compression experiments was 0.04 mm/min.

[93], which also corresponds to a quasi-static strain rate of 1x10-3 s-1. Specimens were loaded to a

49

pre-defined strain or stress level, then unloaded in displacement control, at the same displacement

rate as was used during loading.

The parameters of interest for the stress-strain data from a pseudoelastic experiment

include the critical transformation stress (σA→M), the stress hysteresis (Δσ), the maximum applied

strain (εmax), and the recovered strain (εrec). These parameters are defined in Figure 3.3, which shows

a pseudoelastic response for a conventionally processed Ni50.8Ti49.2 at.% alloy. The critical

transformation stress is determined using a 0.2% offset method. The stress hysteresis is defined as

the stress differential between critical transformation stress at loading and the onset stress during

unloading. The recovered strain or transformation strain is the portion of maximum strain which is

recovered or regained during unloading. If an alloy does not recover all of the applied strain, a

portion of irrecoverable strain remains (εirrec). Two additional material properties, Young’s modulus

during loading (EA) and Young’s modulus during unloading (EM), were also calculated.

Figure 3.3 Pseudoelastic stress-strain response and determined shape memory and

material properties: critical transformation stress (σA→M), stress hysteresis (Δσ), maximum

applied strain (εmax), recovered or transformation strain (εrec), irrecoverable strain remains

(εirrec), Young’s modulus during loading (EA), Young’s modulus during unloading (EM).

50

3.3.3 Shape Memory Effect Behavior, characterized by an Isothermal Mechanical

Experiment and Subsequent Heating Cycle

The shape memory effect behavior in the additive manufactured alloys were characterized

with an isothermal mechanical experiment followed by a heating cycle. The material is initially

composed of correspondent variant pairs of martensite, in a twinned structure. The twins move and

the CVPs reorient to form a detwinned or reoriented martensite structure when stress is applied.

Subsequent heating of the material results in the reverse MT and the material transforms into the

austenite phase. The MTS load frame (shown in Figure 3.2 (a)), with temperature control and

temperature measuring capabilities was utilized for running SME experiments.

The thermo-mechanical experiments for characterizing the SME behavior followed ASTM

standards E8 [92] and E9 [93], following the description in 3.3.2 . Specimens were loaded to a pre-

defined strain or stress level and unloaded in displacement control at the same displacement rate as

was used during loading. During subsequent heating, this small stress was maintained (i.e. machine

control was switched to load control). The heating and cooling rates were approximately 10 °C/min.

The properties of interest are determined from the stress-strain-temperature plot for the

SME behavior, as shown in Figure 3.4. These properties include the critical stress of the twinned

martensite (σM→M+), the residual strain (εres), and the recovered strain (𝜀𝑟𝑒𝑐𝑆𝑀𝐸). The critical stress is

determined using a 0.2% offset, or as the deviation from linear elastic behavior. The residual strain

is the strain which remains after unloading. The recovered strain is the strain recovered during the

subsequent heating cycle. The recovery ratio (𝜂) is calculated as the ratio of recovered strain relative

to the residual strain (𝜂 = 𝜀𝑟𝑒𝑐

𝑆𝑀𝐸

𝜀𝑟𝑒𝑠⁄ ). The Young’s modulus of the twinned martensite (EM) and

Young’s modulus of the reoriented or detwinned martensite (EM+) are determined from the linear

portions of the loading response. During heating, the recovery start temperature (𝐴𝑠∗) and recovery

finish temperature (𝐴𝑓∗) are determined using a method of tangents. The recovery temperature

51

differential (Δ𝑇∗) is calculated as the difference between the recovery finish and recovery start

temperatures (Δ𝑇∗ = 𝐴𝑓∗ − 𝐴𝑠

∗).

Figure 3.4 Shape memory effect stress-strain-temperature response and determined shape

memory and material properties: critical stress of the twinned martensite (σM→M+), residual

strain (εres), recovered strain (𝛆recSME), Young’s modulus of the twinned martensite (EM),

Young’s modulus of the reoriented or detwinned martensite (EM+), recovery start

temperature (As*), and recovery finish temperature (Af

*).

3.3.4 Strain Measurement Length Scales

Throughout thermo-mechanical testing, strain was measured using different techniques

which spanned a range of gage lengths. Engineering strain (𝜀𝑒𝑛𝑔 = ∆𝒍

𝒍𝟎) is determined as the

movement of the actuator in the load frame (i.e. the change in length; ∆𝑙) divided by the original

specimen length (𝑙0). The calculation of engineering strain is based on the original specimen

dimensions. As the specimen is deformed, the specimen dimensions may vary significantly from

the original dimensions. At high strains (typically beyond yielding), the engineering strain value is

significantly altered from the “true” strain value. True strain (𝜀𝑡𝑟𝑢𝑒 = ln (1 + 𝜀𝑒𝑛𝑔)) considers the

52

changing specimen dimensions at large strains. However, the displacement of the actuator is

measured based on the response of the specimen as well as the fixtures, adapters, collets, and other

equipment pieces between the actuator (which measures the displacement) and the fixed crosshead.

Thus, the engineering strain, and the calculated true strain, are not truly indicative of the specimen

response.

Strain measurement techniques including extensometer and virtual digital image

correlation (DIC) extensometers establish the specimen response. To measure strain using an

extensometer, the device is attached to the specimen and the deformation and strain are measured

across a selected gage length (for example, 5 mm). DIC is a full-field deformation analysis

technique, where the deformation is tracked based on a speckle pattern applied to the specimen

surface, the process for which is described in greater detail in the next section. Figure 3.5 shows

gage lengths for engineering and true strain (a), which is orders of magnitude larger than the gage

length for the extensometer (b), which is orders of magnitude larger than the gage length for the

strain measurements using digital image correlation (c). Figure 3.5 (d) shows the differences in the

responses of the various strain measurement techniques. The engineering stress-engineering strain

and the true stress-true strain measurements have a smaller Young’s modulus compared to the

engineering stress-extensometer strain and engineering stress-virtual DIC extensometer

measurements. Additionally, the recoverable strain for the engineering and true strain responses

are higher than for the extensometer and virtual extensometer responses and may be misleading.

As the responses using the extensometer and virtual extensometer are more representative of the

specimen response, these strain measurements will be used when possible throughout this work.

53

(a) (b)

(c) (d)

Figure 3.5 Defined gage lengths for (a) engineering and true strain (b) strain using the

miniature extensometer on a tension and compression specimen, and (c) virtual gage length

for strain contours calculated using digital image correlation. (d) Compressive mechanical

responses from a single experiment for a conventionally processed Ni49.9Ti50.6 at.% alloy.

The engineering strain is plotted versus engineering stress (black). The true strain is plotted

versus the true stress (blue). The extensometer strain is plotted versus the engineering

stress (red) and the virtual extensometer strain is plotted versus engineering stress (green).

54

3.4 Correlating Additive Manufactured Microstructure and Martensitic Transformation

Digital image correlation is a non-contact deformation analysis technique which is able to

track the displacement over the entire surface of a specimen [95]. DIC is likened to many small

strain gages over the region of interest (ROI) on the specimen surface [96] and the deformation

measurements are considered meso-scale. These finer scale deformation measurements enable

localized strain contour evolutions to be visualized. For this technique, a speckle pattern is applied

to the surface of a specimen, images of the specimen are captured in real-time and software is used

to track the speckle movement and calculate corresponding displacements and strains. For shape

memory alloys, the martensitic transformation gives rise to localized strains. These localized strains

can be observed and quantified using digital image correlation, as the gage length size for this

technique is small (this work achieved gage lengths on the order of hundreds of microns). From the

strain contours, the microstructures can be correlated with the underlying phase transformation

morphology. Additionally, the correlation software can calculate the strains in the axial, transverse,

and shear directions. Therefore, Poisson’s ratio can then be determined from the elastic portion of

the material response.

The machine vision system consists of a camera and software which correlates the speckle

pattern and calculates the displacements and strain values. Images were captured using a

Grasshopper GRAS-20S4M/C CCD camera (Point Grey Research Inc., Canada), with a 1600 x

1200 pixel array. Shown in Figure 3.6, a series of lenses (2X adapter, 12X variable zoom lens,

0.25X lens attachment) were used to magnify the specimen surface and maximize the specimen

size relative to the pixel array. A rough and fine stage assisted with focusing the camera on the

specimen surface. A gooseneck light was utilized for illuminating the speckled specimen surface.

VIC2D software (Correlated Solutions, Irmo, SC) was used to correlate the images and create the

strain contours.

55

Figure 3.6 Machine vision systems with the camera and series of lenses

The quality of the applied speckle pattern affects the quality of the DIC results. The

requirements for a good speckle pattern include (i) high contrast, (ii) randomness and uniqueness,

and (ii) an appropriate speckle size (covers 3 by 3 pixels) [95,97]. As the camera captures an image,

a gray-scale value is assigned to the individual pixel, with black representing a gray-scale value of

0 and white representing a value of 255. The camera can easily identify the speckles from the

background when there is high contrast between the background and the speckles. The speckle

pattern needs to be random and unique such that the intensity matching method is able to identify

the intensity and geometry unique to each speckle. The speckle size needs to be appropriate for the

camera magnification. Displacements will not be tracked accurately if the speckles are too large or

small [95,97].

The speckle size was established based on the desired gage length for DIC. In an ideal case,

the DIC gage length would be on the order of the precipitate size (few microns) or the grain size

(few hundred of microns), such that the strain between precipitates or within individual grains could

be determined. Additionally, even at this fine scale, an ideal experiment would allow us to capture

the deformation over the entire specimen dimensions. For this work, the deformation over the entire

specimen dimensions was obtained at the expense of increasing the DIC gage length. The speckle

size was also dictated by the magnification of the camera, as well as the parameters used in the

correlation (e.g. step size and filter size).

56

The desired gage length necessitated very fine speckles and a good speckle pattern.

Speckles were applied to the specimen surface using a micron-CM B airbrush (IWATA, Portland,

OR). Specifically black speckles were applied to a white background [98]. The background layer

of Golden Airbrush titanium white paint (Golden Artist Colors, New Berlin, NY) was thin and

uniform. This thin paint layer was assumed to be perfectly bonded to the specimen surface such

that the strains in the surface were perfectly transferred to the background layer; the white layer

was thin so that the strains are not distortion or influenced by the paint layer. The speckles were

painted with Golden Airbrush carbon black paint (Golden Artist Colors). The compressed air

pressure, specimen to airbrush tip distance, and painting speed were selected to produce a good

speckle pattern. A good speckle pattern is shown in Figure 3.7, at (a) low and (b) high

magnifications for the gage section of a tension specimen, which was used as a standard for all

other speckle patterns in this work.

(a) (b)

Figure 3.7 (a) Low (0.75X magnification) and (b) high (6X) magnifications of a good speckle

pattern for digital image correlation. The speckle pattern has high contrast between the

white background and black speckles, there is a random and unique speckle pattern, and

the speckles cover approximately 3 by 3 pixels (in low magnification image).

The speckle pattern was assessed individually for each specimen by determining the gray-scale

histogram and comparing known rigid body displacement to the calculated rigid body

displacement. A good speckle pattern has a Gaussian distribution of gray-scale values, with a large

57

standard deviation/wide bell-shaped curve that is evenly distributed between black (0 gray-scale

value) and white (gray-scale value of 255). Figure 3.8 (a) shows the gray-scale histogram for the

“good” speckle pattern shown in Figure 3.7 (a). For the rigid body displacement (specimen was not

deformed) comparison, the specimen was incrementally moved a known distance and images were

captured after each movement. These images were then correlated and the calculated displacement

was compared to the set displacement [99]. The good speckle pattern was moved a displacement

of 0.1 mm. This set displacement and the calculated displacement from DIC generally agree, as

shown in Figure 3.8 (b) and Figure 3.8 (c).

(a)

(b) (c)

Figure 3.8 (a) Gray-scale histogram for the “good” speckle pattern (shown in Figure 3.7

(a)). Comparison of the rigid body displacement (RBD) and the calculated RBD as

determined using the speckle pattern and digital image correlation, for RBDs of (b) 0.0 mm

(stationary) and (c) incremental 0.2 mm.

The parameters utilized in the VIC2D correlation software were set to produce the best

correlation based on the specimen speckle pattern. The subset size, step size, and filter size were

58

selected based on the speckle size. The subset size is the area or region of the image which is used

to track the displacement. Thus, the subset must be large enough such that the correlation software

is able to identify unique subsets. The subset size was determined as the region encompassing 3 by

3 speckles, with a speckle covering a 3 by 3-pixel array. The step size is the spacing of the points

analyzed as the images are correlated. A step size of 1 means that every data point within the subset

is correlated. However this step size requires significant computation time. A step size equal to

one-quarter that of the subset size acts as a good initial guess which would result in a good

correlation while not becoming too time-intensive. The filter was set at 5 (smallest value) if the

step size was large, and set higher (around 15) if the step size was small [100]. The filter is the size

of the window over which the data was smoothed. The subset weights were set to Gaussian weights,

which should provide the “best combination of spatial resolution and displacement resolution”

[100]. The interpolation scheme was set to optimized 8-tap, which provided the most accurate

displacement. The criterion was set to normalized squared differences, as this selection is not

affected by changes in lighting (like squared differences) and will always converge/produce a result

(which may not be the case for zero-normalized squared differences).

The gage length of the virtual DIC extensometer is calculated based on the step and filter

size, which is dependent on the subset size, speckle sizes, and machine vision magnification. The

gage length is calculated using Equation 3.2. The virtual gage length is calculated in units of pixels

from Equation 3.2 and converted to units of millimeters based on the camera set-up and

magnification.

𝑽𝒊𝒓𝒕𝒖𝒂𝒍 𝒈𝒂𝒈𝒆 𝒍𝒆𝒏𝒈𝒕𝒉 [𝒑𝒊𝒙𝒆𝒍𝒔] = 𝑺𝒕𝒆𝒑 𝒔𝒊𝒛𝒆 ∗ 𝒇𝒊𝒍𝒕𝒆𝒓 𝒔𝒊𝒛𝒆 Equation 3.2

The virtual gage length ranged from 170 to 300 μm for this work [65,67].

A limitation of this work is that the deformation analysis scale (extensometer and virtual

DIC extensometer) is orders of magnitude larger than the scale at which the martensitic

59

transformation occurs. Individual martensite variants have dimensions on the nanometer scale,

whereas the gage length for the virtual DIC extensometer is on the micron scale. Thus, the scope

of interpretations of microstructure interaction on the MT and shape memory response is limited to

comparisons between the strain contour morphologies and strain values.

Figure 3.9 Scale comparison of fabrication, microstructure, behavior, and martensitic

transformation aspect. The build coupon size and dimensions of the passes and layers

correspond to fabrication. The scale at which the microstructure was characterized is

identified by the composition, grains, and precipitates. The gage lengths for the

extensometer and virtual DIC extensometer are identified for the behavior. The size of the

martensite variants is significantly smaller.

3.5 Chapter Summary

The microstructure and mechanical response of additive manufactured alloys was

extensively characterized. Specifically, the composition, phase, precipitate morphology, and grain

structure microstructure aspects were characterized, as these features are known to impact the

martensitic transformation and the shape memory response. The shape memory response of these

alloys was characterized in accordance with common shape memory characterization

methodologies. Additionally, the full-field deformation analysis technique of digital image

correlation was used to interpret the effect of the microstructure on the martensitic transformation,

and thus the shape memory response.

60

Chapter 4

SPATIAL CHARACTERIZATION OF THE THERMAL-INDUCED

PHASE TRANSFORMATION THROUGHOUT AS-DEPOSITED

ADDITIVE MANUFACTURED NiTi BULK BUILDS


The goal of this chapter was to confirm that the build coupons, fabricated from elementally

blended Ni and Ti powder using laser-based directed energy deposition AM technique, exhibited

characteristic shape memory behavior. Specifically, the thermal-induced martensitic

transformation behavior was investigated using differential scanning calorimetry for selected

locations within the build coupons. The previous works on additive manufactured NiTi shape

memory alloys investigated the TIMT in small volume depositions and did not address the role that

deposition size can have on the shape memory response. The homogeneity of the transformation

response is discussed with respect to the fabrication process. The work shown in this chapter was

published in [48].

4.2 Results and Discussion

Small differential scanning calorimetry samples were extracted from build coupons at

selected locations along the build direction and through the build height. Sections were extracted

along the build (laser travel) direction, designated by y0, ym, and ye locations, as shown in Figure

4.1 (a). From the sections, smaller differential scanning calorimetry samples were extracted along

the build height, as shown in Figure 4.1 (b).

61

(a) (b)

Figure 4.1 Schematic illustrations of the build coupon showing (a) sections extracted along

the build direction [y0 – the beginning, ym – the middle, and ye – the end of the build coupon]

and (b) locations for differntial scanning calorimetry sample extraction along the height/z-

direction.

The DSC thermograms for extracted Ti-rich samples are shown for the forward austenite

to martensite and reverse martensite to austenite transformations, in Figure 4.2 (a) and Figure 4.2

(b), respectively. The thermograms correspond to samples extracted from the same height (4.4 mm)

including similar peak temperatures (Ap = 97 °C; Mp = 59 °C) and forward (HF = 20 J/g) and reverse

enthalpies (HR = 20 J/g), is observed at the same heights along the build direction. The measured

Ms temperatures (approximately 50 °C) are comparable to Ms temperatures for conventional

polycrystalline NiTi SMAs (approximately 70 °C) with similar Ni concentrations [16].

(a) (b)

Figure 4.2 (a) Cooling and (b) heating DSC thermograms for Ti-rich as built samples taken

at the same build height from sections y0, ym, and ye along the build direction. The

martensite and austenite start (Ms and As), finish (Mf and Af), and peak (Mp and Ap)

temperatures are identified. The forward (HF) and reverse (HR) enthalpy measurements

are derived from the peak area.

62

The TIMT occurs at every location within the build. The DSC thermograms for specimens

extracted at different heights along the y0, ym, and ye locations are shown in Figure 4.3 (a), Figure

4.3 (b), and Figure 4.3 (c), respectively. At the 4.4 mm height and higher, the forward MT peak

temperatures (Mp) in each location are 57 ± 4 °C and those for the reverse MT (Ap) are 98 ± 3 °C.

The enthalpy measurements at 4.4 mm and above are 19 ± 2 J/g for each location, which suggest

equivalent transformations. The 1.8- and 3.3-mm locations exhibit a contrasting thermal response,

which may reflect the influence of the substrate, residual stresses, or variable microstructures.

(a) (b)

(c)

Figure 4.3 DSC thermograms with forward (Mp) and reverse peak (Ap) temperatures for

samples taken from section (a) y0 (b) ym, and (c) ye along the build/y-direction.

The Ni-rich builds required post-deposition heat treatment to bring about the TIMT and

observe peaks in the DSC thermograms. The as built Ni-rich build coupon, shown in Figure 4.4,

did not exhibit the TIMT; the expected Ms temperature is most likely below the -120 °C limit of

the equipment, based on the high Ni-content of this alloy.

63

Figure 4.4 TIMT, as measured using differential scanning calorimetry, for samples

extracted from the as built Ni-rich coupon. No transformation peaks are observed in this

temperature range.

Direct aging heat treatments were required to bring about the TIMT. Samples, extracted

from the same build height, were heat treated at three different temperatures (450 °C, 500 °C, 550

°C) and for four different time durations (1.5 h, 3 h, 5 h, 10 h). As shown in the thermo-grams in

Figure 4.5, the treatments resulted in a multi-stage MT, where multiple forward and reverse peaks

are observed. Each heat treatment resulted in two forward MT peaks, MP1 and MP2, in the cooling

DSC thermo-grams. The MP2 peak appears as a broad shoulder. For the 450 °C (Figure 4.5 (a)) and

500 °C (Figure 4.5 (c)) aging temperatures, one reverse MT peak, AP, arises during heating. Two

endothermic events exist for 550 °C. The peak temperature, AP1, corresponds to a broad shoulder

adjacent to the AP2 peak. Multistep MTs in NiTi alloys are typically attributed to an initial MT from

the austenitic B2 phase to an intermediate R-phase that is followed by a transition to the martensitic

B19' phase [101,102]. It is well known that a lower thermal driving force is required to bring about

the R-phase transition [1,37,101,102], which causes the MT to take place at higher temperatures in

the aged Ni-rich AM NiTi alloys.

The evolutions of the characteristic peak temperatures with increasing aging temperature

and increasing hold time are plotted in Figure 4.5 (b), Figure 4.5 (d), and Figure 4.5 (f). The

characteristic temperatures are typical of traditionally fabricated NiTi alloys with similar Ni

concentration [16]. Increasing the hold time from 1.5 h to 10 h facilitates increased peak

temperatures for each aging treatment. The trend of increasing times resulting in increasing peak

64

temperatures is consistent with observation for other NiTi alloys [46,52,101]. Comparing

characteristic transformation temperatures for the same hold time at each temperature, the

characteristic temperatures decrease with increasing heat treatment temperature. Lower forward

MT temperatures suggest an increased driving force is necessary to facilitate the transformation as

the aging temperature increases.

(a) (b)

(c) (d)

(e) (f)

Figure 4.5 DSC thermograms and peak temperatures for Ni-rich samples that were

directly aged at (a, b) 450 °C, (c, d) 500 °C, and (e, f) 550 °C for various time durations.

Samples were all extracted from the same build height.

65

4.3 Chapter Summary and Conclusions

This work reports on the spatial variation of the TIMT temperatures in NiTi builds

fabricated via laser-based directed energy deposition process using Ti-rich (Ti52.1Ni47.9 at.%) and

Ni-rich (Ni53.0Ti47.0 at.%) elemental powder blends. Specimens were micromachined from selected

locations along the build path and height to spatially characterize the MT temperatures. Only the

transformation temperatures very near the substrate are affected, which may be a result of residual

stress or microstructure variation. DSC analyses on specimens taken from higher build heights

showed equivalent transformation temperatures and enthalpy measurements, suggesting similar

microstructures. The phase transformation takes place for the Ti-rich alloy in the as built state;

whereas, the Ni-rich builds required aging treatments typical of conventional NiTi alloys. The

disparity between characteristic temperatures for Ti-rich and Ni-rich builds is consistent with

conventional thermomechanical processed alloys. This work confirms that the NiTi alloys

fabricated using LDED AM can exhibit the TIMT throughout the build and that the transformation

temperatures of as built Ni-rich alloys can be systematically controlled via heat treatment, with

trends similar to conventionally processed NiTi materials.

66

Chapter 5

ANISOTROPIC MICROSTRUCTURE AND SUPERELASTICITY OF

ADDITIVE MANUFACTURED NiTi ALLOY BULK BUILDS

USING LASER-BASED DIRECTED ENERGY DEPOSITION


The goal of this work was to spatially resolve the microstructure and the underlying stress-

induced martensitic phase transformation (SIMT) morphology of additively manufactured Ni-rich

alloys. Previous work explored the spatial variation of the thermal-induced MT using DSC ([48]).

This work focused on the Ni-rich alloys and utilized the same post-deposition heat treatments to

tailor the microstructure and behavior. The microstructure and SIMT response were spatially

characterized by extracting specimens from selected locations. Ultimately, a spatially varying

Ni4Ti3 precipitate structure was characterized. In the as built material, Ni4Ti3 precipitates are

observed with morphologies typical of wrought Ni-rich SMAs after aging treatments. The spatially

varying microstructure is characterized by finer precipitate morphology farthest from the substrate

and coarse morphologies nearest to the substrate. The strain analysis revealed the SIMT

predominately occurs in the finer precipitate morphology. An overaging heat treatment decreased

the degree of anisotropy and facilitated larger recoverable transformation strains. The work

presented in this chapter was published in [67].

67


5.2.1 As built Microstructure and Pseudoelastic Response

The phases and crystallographic structures present in the as built material were first

characterized using a combination of XRD and electron microscopy. Two locations were

characterized, at heights close to the substrate (2 mm) and near the top of the build (8 mm), as

shown in Figure 5.1. The B2 austenite phase is present throughout the build, which is expected of

the Ni-rich alloy at room temperature based on previous analysis [48]. At the 2 mm location, the

XRD scan is dominated by the B2 austenite parent phase, while smaller peaks at 2θ = 43.3° and

62.1° indicate the presence of the Ni4Ti3 phase. At a z-height of 8 mm, the XRD scan is again

dominated by the B2 austenite phase with other smaller peaks corresponding to the Ni4Ti3 phase at

2θ values of 37.7°, 43.3°, 62.1°, and 84.4°. The peak at the 2θ value of 43.3°, however, can be

indexed as either Ni4Ti3 or Ni3Ti. SEM and EDS analyses were utilized to investigate the presence

of the Ni3Ti phase at both locations. This analysis showed that the Ni3Ti phase was not present at

the 8 mm height, but the Ni3Ti phase was present at the 2 mm height.

Figure 5.1 XRD analysis taken from z-height locations of 2 mm and 8 mm above the

substrate of an as built Ni-rich compression specimen extracted from a large (6 pass, 14

layer) build coupon.

68

The morphology and volume fractions of these Ni3Ti and Ni4Ti3 precipitates was further

evaluated using SEM at a range of locations along the longitudinal and transverse directions of the

build. The as built microstructure at two locations along the y-direction, at a constant z-height of

6.5 mm, are shown in Figure 5.2. Ni4Ti3 precipitates exist as fine platelets approximately 0.6 ± 0.1

μm in length and comprising an area fraction of approximately 7 %. Similar Ni4Ti3 precipitate sizes

(0.6 ± 0.1 μm) and area fractions (7 ± 2%) were observed at different locations at this same height,

indicating that along the longitudinal direction, the observed microstructures are similar. In the

SEM images, defects appear as black features. The defects may be micropores [44,46]. The defects

may also be microvoids remaining from unwanted secondary nickel-titanium-oxide phases

[16,103], which were pulled out during the metallographic preparation process.

Figure 5.2 Back scatter electron images of a Ni-rich specimen micromachined from a 6.5

mm z-height (far from the substrate). Images were taken from two locations along the build

direction, a distances of 25 and 31 mm from the build origin.

The Ni4Ti3 precipitate size varies significantly in the height direction in comparison to

along the build direction. The microstructures at z-heights of 2.5 mm, 6.0 mm, and 8.5 mm above

69

the substrate are presented in Figure 5.3 (a), Figure 5.3 (b), and Figure 5.3 (c), respectively. As the

height increases in the as-deposited build, the microstructure, particularly the morphology and the

area fraction of the Ni4Ti3 precipitates, changes. For example, at the 2.5 mm location (Figure 5.3

(c)), the Ni3Ti phase appears globular and a cross-hatch pattern of needle-like Ni4Ti3 precipitate

plates is observed. These precipitates are large (1.6 ± 0.7 μm) and constitute a large portion of the

material (area fraction of 33%). Ni3Ti phase is not observed at the other height locations. With

increasing build height, shown in Figure 5.3 (a) and Figure 5.3 (b), the area fraction and size of

precipitates decreases as the z-height increases. At the intermediate z-height of 6.0 mm, the area

fraction of Ni4Ti3 is 32% and the length is 1.7 ± 0.7 μm. Near the top of the build, at a z-height of

8.5 mm, the Ni4Ti3 precipitates have a length of 0.6 ± 0.1 μm and an area fraction of 15%.

Figure 5.3 SEM images of the as built compression specimen, which are at (a) 8.5 (b) 6, and

(c) 2.5 mm z-heights.

70

The anisotropic microstructure observed in the transverse direction arises from different

thermal histories experienced at each location through the build volume due to the layer-by-layer

AM process [57,104–106]. The as built AM NiTi alloy build coupons contain Ni4Ti3 precipitate

with fine platelet geometries that are typical of aged Ni-rich NiTi SMAs [15,107–109]. Along the

longitudinal direction (y-direction) precipitate morphologies were comparable. In stark contrast,

the precipitate morphologies, sizes, and area fractions vary through the build height. For the Ni-

rich material (Ni52.4Ti47.6 at.%), the Ni4Ti3 phase will precipitate when aged at temperatures below

700 °C, after quenching (i.e. a fast cooling rate) from a high temperatures [37]. AM achieves fast

cooling rates [28,110], which can lead to supersaturated Ni-rich solutions [28] which would have

large driving forces for precipitation of Ni-rich secondary phases [1,37,103,111]. As additional

layers are deposited, the previously deposited layers are reheated [28,105,112,113], the Ni-rich

secondary phases are precipitated. The spatially dependent microstructure can be attributed to the

localized layer-by-layer AM process resulting in spatially varying thermal histories, as well as the

substrate pre-heat leading to lower cooling rates. Near the substrate, the precipitate morphology is

coarse and exhibits the largest area fraction. Ni4Ti3 precipitates are metastable and observed in the

early stages of aging at low temperatures [15,31,37]. As a result, the Ni4Ti3 precipitates are finer

and the area fraction is smallest farthest from the substrate. The Ni3Ti secondary phase is known to

be the final product of decomposition in Ni-rich (> 50.5 at.% Ni) NiTi alloys [31,37], and only

existed closest to the substrate.

The tensile superelastic response for the as built Ni-rich alloy is shown in Figure 5.4 for

two specimens extracted at heights of 3.4 and 6.7 mm above the substrate. At a height of 3.4 mm,

the tensile mechanical response is initially elastic and the Young’s modulus of austenite was found

to be 60 GPa. The stress-induced martensitic transformation from austenite to martensite occurs at

a critical transformation stress of 400 MPa. Further loading, however, results in fracture at a macro-

scale failure strain of 1.6 % and a failure stress of 480 MPa. In order to avoid premature failure of

71

the specimen during testing and observe strain recovery via superelasticity, the specimen extracted

from a height of 6.7 mm was loaded to a stress below this level. The austenitic Young’s modulus

was found to be 50 GPa and the critical transformation stress was 390 MPa. Superelastic strain

recovery begins immediately upon unloading culminating in a macroscale recovery strain of 2.0 %.

The meso-scale full-field tensile strain contours are shown in Figure 5.4 (b) and Figure 5.4

(b) for the tension specimens extracted at 3.4 and 6.7 mm heights, respectively. The numbered

images correspond to the numbered points along the stress-strain curves. Images 1-3 correspond to

the elastic region and show a single predominant color in the ROI. The strain contours display little

to no change during the elastic deformation of austenite, indicating that no SIMT is taking place.

The subsequent images 4 through 6 are beyond the macro-scale critical stress and provide insight

into the underlying SIMT. The maximum meso-scale strains achieved by both tension specimens

is 2.5 %, which exceeds the macro-scale strains of 1.5 % for the 3.4 mm z-height specimen and 2.0

% for the 6.7 mm specimen. When comparing the two heights, areas of maximum meso-scale (2.5

%) strain appear concentrated for the 3.4 mm z-height. In image 6 for that specimen, shown in

Figure 5.4 (b), dashed lines encircle the regions with the maximum meso-scale strain of 2.5 %

within the ROI. Additionally, an area pointed out by an arrow in images 4 through 6 shows a strain

value of 0%, which suggests that the region has not undergone the SIMT. For the specimen at a

height of 6.7 mm, shown in image 6 of Figure 5.4 (c), large ellipses encircle the areas of maximum

meso-scale (2.5 %) strain. There is a larger spatial distribution of localized strain such that the

strain contours are diffuse compared to the concentrations observed for the specimen at 3.4 mm.

Thus, the results reveal that the SIMT is predominant for the tensile specimen extracted farthest

from the substrate.

72

(a)

(b) (c)

Figure 5.4 (a) Tensile stress-strain response for as built specimens extracted 3.4 mm and

6.7 mm above the substrate. DIC axial strain contour images for specimens extracted (b)

3.4 mm and (c) 6.7 mm above the substrate. The image numbers in (b) and (c) correspond

to the numbers in (a). The strains are axial measurements along the loading direction and

parallel to the build/y-direction.

The tensile axial and transverse strains were also calculated using a virtual extensometer

with a 5 mm gage length. The transverse strain-axial strain plot is shown in Figure 5.5 (a). The

strain values are similar for the specimens extracted at heights of 3.4 and 6.7 mm and the Poisson’s

ratio (νYX) was calculated as 0.44. The meso-scale full-field transverse strain contours are shown in

Figure 5.5 (b) and Figure 5.5 (c) for the tension specimens extracted at 3.4 and 6.7 mm heights,

respectively. The numbered images correspond to the same points identified in Figure 5.4 (a). The

regions which accrue the largest axial strains also accrue the largest transverse strains, as identified

73

by ellipses. Similar to the axial contours, diffuse contours are observed for the specimen at a height

of 6.7 mm compared to the concentrations observed for the specimen at 3.5 mm.

(a)

(b) (c)

Figure 5.5 (a) Tensile transverse strain-axial strain response for as built specimens

extracted 3.4 mm and 6.7 mm above the substrate. Transverse and axial strains are

measured using virtual extensometer from DIC. Poisson’s ratio (νYX) is the ratio of the

transverse strain (εX) to the axial strain (εY). DIC transverse strain contour images for

specimens extracted (b) 3.4 mm and (c) 6.7 mm above the substrate. The image numbers

in (b) and (c) correspond to the numbers in (a). The strains are transverse measurements

perpendicular the loading direction and parallel to the pass-/x-direction.

The as built compressive superelastic response is shown in Figure 5.6. The mechanical

response is elastic to the compressive critical stress of 670 MPa. Beyond this stress, the response

exhibits a nearly linear slope and a strain hardening-like response. Strain recovery ensues upon

unloading with a residual strain of -0.2 % remaining after complete unloading. Hence the

transformation strain, which is the difference between the maximum strain and the strain value

upon complete unloading, was nearly -3.1 %. The meso-scale strain contours are shown in Figure

74

5.6 (b). Contour images showing strain values below -0.5 % correspond to the elastic response.

The image showing -0.7% strain corresponds to the critical transformation stress. The maximum

meso-scale strain (-4.7 %) exceeds the macroscale strain (-3.3 %). Dashed lines in the images

encircle diffuse areas of higher strain magnitude compared to the surrounding ROI.

(a)

(b)

Figure 5.6 (a) Compressive stress-strain response for an as built specimen. The macroscale

strains are axial measurements along the loading direction and parallel to the build height

(z-dir). (b) DIC axial strain contour images numbered corresponding to numbers along the

stress-strain curve. Unloading images below loading images correspond to the same

macroscale strain.

The multiscale mechanical deformation analysis, including DIC analysis for microscale

strain measurements, confirms the underlying reversible SIMT for superelasticity and reveals that

the presence or absence of the SIMT can be correlated to the anisotropic microstructure. For the

SE, the SIMT is crystallographically reversible, in which the martensite (product) phase reverts

back to the austenite (parent) phase in the original orientation and follows the reverse path of the

75

forward MT, due to the lattice correspondence between the parent and product phases [1,114].

Hence corresponding loading and unloading images exhibit similar strain contours with the regions

that underwent the forward (i.e. during loading) SIMT first having the highest microscale strains at

the onset of the reverse SIMT during unloading. Moreover, regions that transformed first are the

last to undergo the reverse SIMT. Because coarse secondary phases and larger precipitate area

fractions are found close to the substrate, less B2 NiTi phase is available to undergo the SIMT.

Consequently, close to the substrate strain values were the lowest for tensile and compressive stress

states. The coarse precipitate morphology facilitated localized concentration of higher strains

nearest the substrate. Farthest from the substrate, in the finer precipitate microstructure, a larger

volume fraction of B2 is available to undergo the SIMT. Thus, the strain values were highest and

the contours were diffuse. The results confirm that the AM NiTi materials undergo the reversible

SIMT and exhibit SE at the finer deformation measurement scale. Typically, Ni-rich NiTi alloys

undergo a thermo-mechanical process to bring about the desired Ni4Ti3 microstructure and

superelastic behavior [1,37]; the SE behavior is observed in the as built LDED AM alloys.

5.2.2 Aged Microstructure and Pseudoelastic Response

As built alloys were aged in order to alter the microstructure by precipitating and growing

the desired Ni4Ti3 phase. Figure 5.7 and Figure 5.8 show the phases and precipitate morphologies

for a directly aged (550 °C, 3 h) specimen. For the XRD scan in Figure 5.7, one location was

characterized. The B2 austenitic NiTi phase is still the primary phase after the direct aging

treatment. Additionally, the only other phase identified is the desirable Ni4Ti3 phase. The precipitate

morphologies were investigated at z-heights of 8.5 mm, 6.0 mm, and 2.5 mm. Ni4Ti3 precipitates

are observed at all heights. The measured Ni4Ti3 lengths and area fractions are similar for the three

analyzed locations (30% area fraction, 1.3 ±0.5 μm length). Overall, it appears the direct aging heat

76

Figure 5.7 XRD analysis for a directly aged (550 °C, 3 h) Ni-rich sample.

Figure 5.8 Back scatter electron images of the aged (550 °C, 3 h) compression specimen,

which were taken at z-heights of (a) 8.5 mm; (b) 6 mm; and (c) 2.5 mm, measured from the

substrate.

77

treatment decreased the differential precipitate microstructure through the height, resulting in more

consistent area fractions and precipitate lengths. However, there remains some degree of variation,

which may be inherited from the starting as built microstructure.

The tensile stress-strain response and the corresponding strain contours for an aged (550

°C, 3 h) specimen are shown in Figure 5.9. The aged specimen achieved a higher macro-scale strain

at fracture (3.1 %) compared to the as built specimen extracted far (6.5 mm) from the substrate (2.0

%; Figure 5.4 (b)). The meso-scale tensile strain value of 4.5% for the directly aged material is also

higher compared to the as built material (2.5 %; Figure 5.4 (d)). Strain contour images 4-7 (Figure

5.9 (b)) illustrate the SIMT. The regions of maximum meso-scale strain (4.5 %) are circled by

dashed lines in image 6. Within the contour, the maximum strain regions are concentrated. Aging

facilitated concentrated high strain regions throughout the ROI, in contrast to the diffuse contours

observed for the as built material. The transverse strain-axial strain and the transverse strain

contours for the aged specimen are shown in Figure 5.9 (c) and Figure 5.9 (d), respectively.

Concentrated regions of localized strain are encircled within the transverse strain contours. The

maximum transverse strain values (-3.0 %) are also larger compared to the as built transverse strains

(-1.8 %).

The compressive stress-strain response and strain contours are shown in Figure 5.10.

Complete compressive strain recovery via SE was observed for the aged specimen in contrast to

the residual strain present in the as built alloy after unloading. Even though the as built and aged

specimens were loaded to equivalent compressive stresses, the macroscale strain for the aged

material (-4.0%) was higher than the as built material (-3.3%). The micro-scale strain values

achieved by the aged specimen (-6.25%) were also greater than the micro-scale strains achieved by

the as built specimen (-4.7%). Beyond the critical transformation stress, as shown in image 4,

localized areas of lower strain appear towards the bottom of the ROI. Higher strain appears in the

78

(a) (b)

(c) (d)

Figure 5.9 (a) Tensile stress-strain response and (b) DIC axial strain contour for aged (550

°C, 3 h) specimen, extracted at a height far from the substrate (5.3 mm). The strains are

axial measurements along the loading direction and parallel to the build/y-direction. (c)

Tensile transverse strain-axial strain response for aged measured using virtual

extensometer from DIC. Poisson’s ratio (νYX) is the ratio of the transverse strain (εX) to the

axial strain. (d) DIC transverse strain contour images for points identified in (a). The

strains are transverse measurements perpendicular the loading direction and parallel to

the pass-/x-direction.

upper regions at the top of the ROI ranging from -4.7% to -6.25%. Unloading and loading images

correspond to similar macroscale strain values and thus the results confirm the underlying

reversible SIMT. Despite the decreased anisotropy of the precipitate morphology, the spatial

distribution of the compressive strain contours within the ROI appears similar through the build

height for aged and as-deposited materials. In both cases, the highest microscale strains are

observed farthest from the substrate.

A heat treatment at 550 °C for 3 hours decreased the variation in the Ni4Ti3 precipitate

microstructure and higher micro-scale DIC strain magnitudes and macroscale extensometer strain

were observed after aging. Equivalent Ni4Ti3 precipitate morphologies were observed through the

build height for the aged material and the Ni3Ti phase was not present. The heat treatment time and

79

(a) (b)

Figure 5.10 (a) Compressive stress-strain response for an aged (550 °C, 3h) specimen. The

macroscale strains are axial measurements along the loading direction and parallel to the

build height (z-dir). (b) DIC axial strain contour images numbered corresponding to numbers

along the stress-strain curve. Unloading images below loading images correspond to the same

macroscale strain.

temperature has been identified as an overaging heat treatment [40,43]. Coarsening the precipitate

morphology via heat treatment removed the disparity of precipitate volume fraction through the

build height and DIC analysis revealed localized SIMT strains increased compared to the as built

alloy. Higher transformation strain has been related to incoherent Ni4Ti3 precipitates compared to

semi-coherent precipitates and based on the postulation that incoherent precipitates do not curtail

the detwinning contribution to the transformation strain [40,43], it is reasonable to assume that an

overaging heat treatment compromised the coherency of the pre-existing Ni4Ti3 precipitates, which

can result in the larger transformation strain observed after aging. Despite the equivalent precipitate

morphologies and decreased anisotropy, both aged and as built microscale compressive strains

show localized areas of the lowest values closest to the substrate and the highest strains farthest

from the substrate. The result suggests that factors known to control the transformation strain

morphology including a gradient in composition, the B2 crystallography, and/or the grain

morphology may persist throughout the build height. Thus, systematic fine scale probing

80

microstructure analyses are paramount for correlating fabrication-microstructure-property

relationships for AM NiTi SMAs.


The relationship between the microstructure and superelastic behavior in Ni-rich NiTi

builds fabricated using laser-based directed energy deposition additive manufacturing was

investigated. By using elementally blended Ni and Ti powders, the feedstock composition was

tailored to obtain Ni-rich build coupons for which superelasticity is expected. Ni4Ti3 precipitates

were observed in the as built alloys, even though they are typically observed only after an aging

heat treatment of Ni-rich NiTi SMAs. This work is the first to report on the consequences of the

spatially varying AM microstructures on the superelastic behavior, using full-field deformation

measurements to correlate the underlying SIMT to the AM microstructure. For the as built material,

the variation of the size and area fraction of the precipitates was most significant through the build

height, whereas the precipitate morphology was, by comparison, uniform in the build direction.

The SIMT was predominant in material farthest from the substrate and micro-scale full-field strain

measurements, via digital image correlation analysis, reveal the largest transformation strain levels

and diffuse strain contours. Aging the as built alloy (i) coarsened the Ni4Ti3 precipitates, which

reduced the degree of spatial variation through the build height, and (ii) caused larger macro- and

micro-scale transformation strains. The evolution of the DIC strain contours during loading and

unloading underscore that the underlying SIMT proceeds in a reversible manner and thus AM Ni-

rich NiTi SMAs exhibit superelasticity.

81

Chapter 6

MULTI-SCALE SHAPE MEMORY EFFECT RECOVERY IN NiTi

ALLOYS ADDITIVE MANUFACTURED BY SELECTIVE LASER

MELTING AND LASER DIRECTED ENERGY DEPOSITION


The goal of this chapter was to characterize NiTi alloys fabricated using both powder bed

fusion and laser-based directed energy deposition techniques, in order to provide new insights into

the different techniques. Near equiatomic NiTi alloy compositions were investigated. The

composition and thermal-induced phase transformation temperatures are spatially resolved to

assess the different AM techniques. The martensitic microstructure exists at room temperature, and

thus deformation is recovered by the shape memory effect during heating. A multi-scale

deformation analysis elucidates differential transformation morphologies for the PBF and LDED

techniques. This work was published in [79], in collaboration with Dr. Mohammad Elahinia and

Mohsen Taheri Andani from the University of Toledo.


6.2.1 Compositional Analysis and Phase Transformation Temperatures

Samples for DSC analysis were sectioned from compression specimens along the build

height. Nine EDS measurements were taken on each sample and the average compositions of four

samples are in Table 6.1. Locations are measured with respect to the surface of the compression

specimens nearest the support/substrate. The PBF compositions are near-equiatomic with standard

deviations between ±0.1 and ±0.4 at.%. In stark contrast, the LDED samples exhibit variable

82

compositions spanning near-equiatomic to Ni-rich and deviations between ±0.6 at.% and ±1.7 at.%.

The average of the 36 measurements for PBF and LDED are, respectively, Ni50.1Ti49.9 ±0.3 at.%

and Ni51.0Ti49.0 at.% ±1.2 at.%.

Table 6.1 Average composition measurements for compression specimens extracted from

a large build coupon. Specimen 1 was further sectioned along the build height.

AM

technique Specimen

Sample: approximate

sectioning height (mm)

Average

composition (at.%) Fig. number

PBF 1

6.0 Ni50.0Ti50.0 (± 0.1)

Figure 6.1 (a),

(b)

4.0 Ni50.2Ti49.8 (± 0.3)

3.0 Ni50.1Ti49.9 (± 0.2)

1.0 Ni50.1Ti49.9 (± 0.4)

2 - Ni49.8Ti50.2 (± 0.2)

LDED

1

6.0 Ni50.2Ti49.8 (± 0.9)

Figure 6.1 (c),

(d)

4.0 Ni51.0Ti49.0 (± 0.6)

3.0 Ni51.8Ti48.2 (± 0.8)

1.0 Ni51.1Ti48.9 (± 1.7)

2 - Ni49.4Ti50.6 (± 1.6)

3 - Ni51.3Ti49.7 (± 0.3)

The location dependence of transformation temperatures is shown in Figure 6.1. In Figure

6.1 (a), PBF transformation temperatures range from 15 °C up to 80 °C with Ms = 56 ±3 °C and Af

= 84 ±2 °C. Despite the differential compositions, the Ms temperatures correspond to those expected

for Ti-rich or near-equiatomic NiTi alloys [16]. In Figure 6.1 (c), LDED transformation

temperatures range from 55 °C up to 110 °C with Ms = 77 ±2 °C and Af = 106 °C ±2 °C. The

variations for each of the transformation temperatures is less than 5 °C. Selected PBF and LDED

samples were subjected to five thermal cycles and temperatures varied by less than ±5 °C; hence,

the temperatures in Figure 6.1 (a) and Figure 6.1 (c) are representative of stable values. The

location dependencies of transformation enthalpies are shown in Figure 6.1 (b) and Figure 6.1 (d).

For PBF samples, a narrow variation is apparent in Figure 6.1 (b). The measurements for LDED

samples in Figure 6.1 (d) exhibit a marked increase. Close to the Ti substrate, the melt pool may

pick up additional Ti material from the substrate and Ti-rich secondary phases may form results in

an overall increase in Ni content, and a decreased enthalpy [16] for the LDED specimen.

83

Additionally, residual stresses may impede the MT, resulting in the observed enthalpy variation as

a function of height. Small variations for PBF contrasts the differential enthalpy measurements for

LDED and underscore microstructural inhomogeneity for LDED that is consistent with trends in

compositional analysis.

(a)

(b)

(c) (d)

Figure 6.1 Thermal-induced phase transformation temperatures (Ms, Mf, As, Af) and

enthalpies (HR, HF) measured from DSC analysis of samples sectioned along the build height

(z-direction) of (a) and (b) PBF and (c) and (d) LDED compression specimens.

6.2.2 Microstructural Analysis

The compression specimen surface (XZ) utilized for DIC analysis was investigated. BSE

images of the precipitate morphologies are shown in Figure 6.2; XRD analysis results of the phases

are shown in Figure 6.3; and optical microscopy images of the grain structure are shown in Figure

6.4. Microconstituent phases are not evident in Figure 6.2. XRD analysis confirms the martensitic

B19' phase is predominate at each location for PBF and LDED alloys. The Ti2Ni and austenitic B2

NiTi phase are evident in XRD scans (Figure 6.3 (a)) of PBF alloys, albeit trace amounts relative

84

to martensite. The SEM images in Figure 6.2 (b) show black features for the LDED specimen,

which may result from entrapped gasses or metallographic surface preparation [67]. Though LDED

specimen compositions vary with height (Table 6.1), the B19' phase is present at every location.

Note that the single peak in the spectra reflect texture. The XRD scans were taken after thermo-

mechanical cycling and the texture may be attributed to residual stabilized martensite. For the PBF

specimen (Figure 6.4 (a)), grains are columnar with the long axis oriented in the build height/z-

direction. The average grain length is 0.15 ±0.05 mm, which is larger than the layer thickness (30

μm), thus the grains traverse multiple layers. The grains in LDED alloys (Figure 6.4 (b)) are

equiaxed by comparison with larger sizes (0.7 ±0.4 mm) equivalent to the layer thickness. Note

that the globular dark features in Figure 6.4 (b) are attributed etchant burn. The lines within the

grains for the LDED specimen at a z-height of 7 mm may be twins.

(a) (b)

Figure 6.2 Back scatter electron images of the precipitate morphologies for (a) PBF and (b)

LDED alloys at z-heights of 7 mm, 4 mm, and 1 mm.

85

(a) (b)

Figure 6.3 XRD scans of the as built (a) PBF and (b) LDED specimens at various z-heights.

Specimens were mechanically cycled prior to analysis, which may have resulted in a

preferential orientation observed by the strong intensity for the 42° 2θ peak for the LDED

specimen.

A comparison of the microstructure between the Ti-rich LDED and PBF builds provides

insight to the microstructure differences between the two AM techniques. The laser energy

densities are similar (around 60 J/mm3) and martensite is the primary phase. Therefore, these builds

are expected to have comparable transformation temperatures and exhibit shape memory effect

behavior. Generally, the composition and the microstructure are more uniform for the PBF builds.

This is in stark contrast to the LDED builds. The uniform microstructure of the PBF builds may be

due to the smaller melt pool relative to LDED. The melt pool size depends on the laser power, the

scan speed, and the spot size [28,115], which influences the grain morphology [74,115,116]. A

comparison of these processing parameters shows that the LDED deposition had a larger melt pool

size, which results in slower cooling rates and larger grain sizes [74,104]. The grains of the LDED

builds are significantly larger than the grains of the PBF builds (0.7 mm for LDED compared to

86

(a) (b)

Figure 6.4 Optical microscopy images of the grain structure for the (a) PBF and (b) LDED

specimens at varying z-heights.

0.15 mm for PBF). Additionally, the direction of the largest thermal gradients within the large melt

pool for LDED can deviate from the build height/z-direction [74], which may result in variable

grain sizes and orientations. By comparison, the columnar grains are all oriented in the z-direction

for the PBF build. The more uniform composition in the PBF build may be a result of using alloyed

powder for the PBF builds as opposed to the blended elemental powders used for the LDED builds

[117].

87

6.2.3 Martensite Deformation and Shape Memory Effect Recovery

The compressive SME behavior was characterized for specimens with the loading direction

along the height/z-direction, at a test temperature corresponding to the material being in the

martensitic phase. Characteristic material properties are summarized in Table 6.2. The initial strain

response is characterized by the elastic modulus 𝐸𝑀. The critical stress 𝜎𝑐𝑟𝑀→𝑀+

defines the deviation

from linear elasticity attributed to the onset of the reorientation and detwinning. M denotes stress-

free twinned martensite and M+ denotes oriented+detwinned martensite, which represents multiple

contributions at increasing strain levels: M → preferentially oriented correspondent-variant pair,

oriented CVPs → detwinned martensite, and elastic deformation of detwinned martensite [96].

Beyond the yield stress 𝜎𝑦𝑀+

, the detwinned martensite plastically deforms. Upon unloading, a

residual strain 𝜀𝑟𝑒𝑠 remains. Recovery of the residual strain via heating begins at the 𝐴𝑠∗ temperature

and ends at 𝐴𝑓∗ . For the LDED material, the 𝐴𝑠

∗ = 90 °C, which exceeds As = 80 °C. The larger 𝐴𝑠∗

with respect to As reflects that residual martensite is stable after deformation [96]. Upon cooling,

the strain increases once the temperature is below the Ms temperature.

Recovery ratios for PBF and LDED materials are in Table 6.2. The amount of strain

recovered via SME, 𝜀𝑟𝑒𝑐𝑆𝑀𝐸, is the difference between the strain at a temperature of 𝐴𝑠

∗ and the strain

value after heating and being cooled to room temperature (final strain of a cycle). The recovery

ratio, η, is the ratio between the recovered strain (𝜀𝑟𝑒𝑐𝑆𝑀𝐸) and the residual strain upon unloading

(𝜀𝑟𝑒𝑠). The stress-strain-temperature responses for the PBF and LDED specimens shown in Figure

6.5 (a) and Figure 6.5 (b) correspond to stress levels well beyond 𝜎𝑦𝑀+

. SME recovery is incomplete

as the martensite is permanently deformed at the high stress levels.

88

Table 6.2 AM fabricated Ti-rich NiTi material properties.

PBF

Figure 6.5

(a)

LDED

Figure 6.5

(b)

Forward MT start temperature Ms (°C) 55 75

Reverse MT start temperature As (°C) 50 80

Reverse MT finish temperature Af (°C) 80 105

Elastic modulus (twinned martensite) EM (GPa) 60 65

Poisson’s ratio (twinned martensite) νZX 0.44 0.48

Critical stress (twinned martensite) 𝝈𝒄𝒓𝑴 →𝑴+ (𝑴𝑷𝒂) 150 150

Elastic modulus

(reoriented/detwinned martensite) EM+ (GPa) 30 25

Yield stress (reoriented/detwinned

martensite) 𝝈𝒚

𝑴+ (𝑴𝑷𝒂) 1050 1070

Residual strain 𝜺𝒓𝒆𝒔 (%) 6.7 5.6

Recovered strain 𝜺𝒓𝒆𝒄𝑺𝑴𝑬 (%) 3.2 2.3

Recovery ratio 𝜼 = 𝜺𝒓𝒆𝒄𝑺𝑴𝑬 𝜺𝒓𝒆𝒔⁄ (%) 48 41

Recovery start temperature 𝑨𝒔∗ (°𝑪) 95 90

Recovery finish temperature 𝑨𝒇∗ (°𝑪) 110 120

Recovery temperature differential ∆𝑻∗ = 𝑨𝒇∗ − 𝑨𝒔

∗ (°𝑪) 15 30

.

(a) (b)

Figure 6.5 Compression stress-strain-temperature curves for (a) PBF and (b) LDED NiTi

alloys initially in the martensitic phases at T = 23 °C. Material parameters are defined in

Table 6.2.

The SME responses in Figure 6.6 (a) and Figure 6.7 (a) were further characterized to a

stress level below the martensite yield stress, to prevent plastic deformation. Once the critical stress

89

𝜎𝑐𝑟𝑀→𝑀+

is exceeded, the stress-strain response for the LDED specimen exhibits a hardening-like

slope (Figure 6.7 (a)) compared to the response for the PBF specimen (Figure 6.6 (a)). The

difference in behavior beyond the critical stress, along with the contrasting 𝐴𝑠∗ temperatures, and

recovery temperature differentials underscore that the two techniques exhibit dissimilar martensitic

transformation morphological evolutions. Microscale in-situ full-field DIC strain contour images

in Figure 6.6 (b) and Figure 6.7 (b) show the local strain fields during mechanical deformation and

subsequent heating. The residual strain is completely recovered via the SME. Recovery begins at

𝐴𝑠∗ temperatures below the stress-free As temperature in Figure 6.6 and Figure 6.7, and thus the

martensite becomes unstable compared to the responses in Figure 6.5.

For the PBF specimen, the images 1-3 in Figure 6.6 (a) show uniform contours throughout

the ROI corresponding to elastic strain. Images 4 and 5 are associated with the onset of detwinning

near the critical stress and microscale localized strains (-2.5%), are concentrated in the regions in

image 4 that exceed the macro-scale strain (-0.5%). In images 6-8, the strains increase causing the

contour color to change in a diffuse manner throughout the ROI. The ROI in image 8 exhibits a

relatively uniform contour contrast. Localized strains (-5.5%) are concentrated in the encircled

regions in images 8-10. As the strains exceed the macro-scale strain (-2.7%), they correspond to

the elastic deformation of detwinned martensite. During unloading, strain recovery takes place

throughout the ROI resulting in image 13 resembling image 8. Hence, the full-field measurement

underscores that elastic deformation of detwinned martensite ensues when macro-scale strains

exceed the value (-2%) corresponding to image 8. During heating martensitic regions revert to

austenite via SME. Images 14-21 illustrate that the austenite/martensite boundary traverses the

ROI. The microscale strains within image 22 primarily show regions of 0% strain contours, which

show the undeformed austenite structure has been recovered via SME recovery. Upon cooling from

above Af and once reaching the Ms temperature, the macro-scale strain increases to an approximate

value of -0.5%, and corresponding images 23-38 also depict a slight increase in strain.

90

(a) (b)

Figure 6.6 Compressive stress-strain-temperature responses for a PBF NiTi alloy showing

multi-scale SME recovery using (a) macro-scale extensometer measurements and (b) meso-

scale in-situ full-field measurements. (b) In-situ DIC strain contour images with numbers

corresponding to those along the loading, unloading, heating, and cooling segments in (a).

The macro-scale stress-strain-temperature response and the corresponding microscale in-

situ full-field DIC strain contour images for the LDED NiTi alloy are shown in Figure 6.7. For the

elastic response, images 1-3 illustrate a uniform strain field consistent with elastic deformation.

Deformation near the critical stress corresponds to images 4-6. Boundaries are delineated that

illustrate strain levels increase primarily within the band in the ROI. During the hardening-like

response in the stress-strain curve, rising strains are predominant within the banded region in

images 5-9. Localize strains reach about -3.5%, while macro-scale measurements are just above -

2%. Localized strains (-5.5%) are concentrated in the encircled regions in images 9 and 11 that

exceed the macro-scale strain (-3.2%). Notice that contour strain levels increase in the upper right

corner of the ROI. Those levels decrease during unloading. Moreover, the banded region can be

delineated after complete unloading in image 14. From images 14 to 15, strain recovers from the

onset of heating and encircled localized regions of 0 % strain in the upper right of image 15 are

91

larger in size compared to image 14. During SME recovery in images 15-19, the contour colors to

change in a diffuse manner. Regions of 0% strain are prevalent in image 19. Upon cooling and near

the Ms temperature, the strain contours and macro-scale strain values reveal a small strain increase.

(a) (b)

Figure 6.7 Compressive stress-strain-temperature responses for an LDED alloy showing

multi-scale SME recovery using (a) macro-scale extensometer measurements and (b) meso-

scale in-situ full-field measurements. (b) In-situ DIC strain contour images with numbers

corresponding to those along the loading, unloading, heating, and cooling segments in (a).

6.2.4 Stress-Strain-Temperature Cycling

The repeatability of martensite deformation was investigated for 15 stress-strain-

temperature cycles for the PBF and LDED NiTi alloys in Figure 6.8 (a) and Figure 6.8 (b), with

the corresponding SME recovery ratios shown in Figure 6.8 (c). For the cyclic stress-strain

responses shown in Figure 6.8Error! Reference source not found., the end condition was a stress l

evel close 𝜎𝑐𝑟𝑀→𝑀+

, therefore the pre-strain values may differ between PBF and LDED specimens

and between cycles. The 𝐴𝑠∗ does not vary significantly for PBF (T = 62 ±5 C) or LDED (T = 29

±1 C). After 5 cycles, the critical stress 𝜎𝑐𝑟𝑀 →𝑀+ decreases such that transitions from linear-elastic

92

to non-linear martensite reorientation/detwinning responses are indistinguishable in stress-strain

responses. The EM modulus reduces from 60 to about 10 GPa for the PBF alloy and from 65 to 20

GPa for the LDED NiTi alloy. The moduli are nearest the values for the EM+ of detwinned

martensite in Table 6.2, which suggests residual preferentially oriented detwinned martensite may

exist at the onset of deformation. For the PBF and LDED alloys, the recovery ratio stabilizes by

the eighth cycle in Figure 6.8 (c). For both the PBF (Figure 6.8 (a)) and LDED (Figure 6.8 (b))

NiTi alloys, cycle 10 and cycle 15 are similar and show that the responses are stable. For the

stabilized 100% recovery ratio, macro-scale residual strains recovered via SME are around -2 %.

(a) (b)

(c)

Figure 6.8 Stress-strain-temperature cycling up to n = 15 cycles for (a) PBF and (b) LDED

alloys. The n=1 corresponds to Figure 6.6 (a) and Figure 6.7 (a). After unloading, the

specimen was heated for SME recovery. (c) Corresponding strain recovery ratios (η).

93

PBF LDED

correlated

to ref.

image

n=1 n=2 n=3 n=5 n=10 n=15 n=2 n=3 n=5 n=10 n=15

(a) (b)

in-situ

DIC

n=15

(c) (d)

Figure 6.9 Full-field strain measurement contour images. (a) and (b) DIC analysis used the

undeformed image in n = 1 as the reference image for correlation in order to determine the

full-field strain measurement at the beginning of cycles n = 2, 3, 5, 10, and 15. (c) and (d) In-

situ DIC analyses for n=15 with the reference image at the start of the cycle.

The images in Figure 6.9 (a) and Figure 6.9 (b) are correlations between the reference

image for n=1 and the first images at the beginning of n = 2, 3, 5, 10, and 15. The images illustrate

the existing morphologies of residual martensitic strain that accrued through cycling. Comparing

the images in Figure 6.9 (a) to image 8 in Figure 6.6 (b), notice that a nearly uniform contour

dominates the ROI for the PBF alloys. Comparing the images in Figure 6.9 (b) to image 9 in Figure

6.7 (b), nearly uniform contours exist within a band for the LDED alloys. The results illustrate

residual martensite accrues through the same martensite conversion pathway. Figure 6.9 (c) and

Figure 6.9 (d) show the in-situ DIC measurement for n = 15. The reference image is the first image

94

for the cycle. The evolutions of the strain contours are comparable to those in Figure 6.6 (b) and

Figure 6.7 (b), respectively. The results corroborate that the underlying micro-scale transformation

pathway is unchanged by cycling.

NiTi SMAs were fabricated using selective laser melting and laser directed energy

deposition additive manufacturing techniques. The compressive shape memory effect behavior was

investigated in the as-built conditions. For PBF, net-shape compression specimens were fabricated

using alloy NiTi powder feedstock (Ni50.09Ti49.91 at.%). For LDED, specimens were micromachined

from build coupons fabricated using elementally blended Ni and Ti powder feedstock (Ni53Ti47

wt.% or Ni47.9Ti52.1 at.%). Despite contrasting build parameters, laser energy densities were about

~60 J/mm3. Transformation temperature ranges were above room temperature. Based on the well-

accepted relationship between Ni content and Ms temperatures [16], the alloys should behave as

near-equiatomic NiTi compositions. Indeed the moduli, critical and yield stresses are equivalent

for the PBF and LDED NiTi alloy compression specimens.

There is a disparity between the feedstock versus build composition for the LDED material.

SME behavior at room temperature substantiates the composition is much closer to the elementally

blended feedstock (47.9 at.% Ni). Accounting for the 1 at. % error associated with EDS, Ni-

enrichment in the NiTi matrix may result for LDED. This is in opposition to previous AM works,

which postulated that the Ni concentration would decrease based on preferential Ni evaporation

during the AM deposition process [54], as Ni has a lower melting temperature (1455 °C) compared

to Ti (1668 °C). As discussed in [16], enthalpy decreases within an increase in Ni content and

substantiate local compositional variation within the LDED build. Frenzel et al [16]. present

empirically established dependencies of Ms temperatures on at.% Ni concentration, Ms ≈ 50 °C for

Ti-rich and near equiatomic NiTi alloy compositions. The Ms temperatures (~75 °C) for LDED

materials are much higher than expected for 51 at.% Ni, whereas ~55 °C for PBF agrees. The Ni

95

concentration for LDED requires investigation on a finer scale with higher accuracy probing

methods, as the EDS compositional analysis produced counterintuitive results.

Multi-scale deformation analysis was employed which utilized an extensometer for macro-

scale strain measurements and digital image correlation for full-field micro-scale deformation

measurements. The PBF and LDED specimens exhibited complete shape memory effect recovery

with concentrated micro-scale strain levels as high as -4% completely recovered during heating.

The more uniform composition and grain structure for PBF alloys can cause the observed more

uniform strain evolution. For the spatially varying LDED alloy microstructure, when stress exceeds

the critical value, localized regions of increasing strain contours are bounded and grow throughout

the ROI as conversion progresses to completion. The prevailing strain localization facilitates a stark

strain hardening-like response compared PBF. Higher stress levels can produce the seemingly

textured martensite observed in the XRD analysis for LDED alloys. a strain hardening-like beyond

𝜎𝑐𝑟𝑀 →𝑀+ and the stress increase is much greater than PBF. The resulting residual martensite is

unstable and SME recovery ensues when heating initiates for the LDED alloy. On the other hand,

for the PBF material, a finite temperature increase (i.e. thermal driving force) is required to facilitate

SME recovery in the relatively homogeneous oriented/textured microstructure.

A martensite stabilization effect was observed after eight cycles, which is confirmed by

DIC analysis. Stabilization of recovery ratios and elastic moduli, as well as 𝜎𝑐𝑟𝑀 →𝑀+ becoming

immeasurable are indicative of the stabilization effect. A stabilized phase transformation pathway

facilitates a lower driving force for martensite reorientation and detwinning, as back stress

associated with the residual martensite can assist the conversion [118]. DIC measurements confirm

localized residual strain accrues with cycling and that the detwinning pathway is stabilized for both

LDED and PBF despite the dissimilar grain structures and composition. The presence of residual

martensite lowers the elastic moduli and critical stresses during cycling. The textured martensite

96

in XRD scans compared to PBF results in moduli approaching values for reoriented/detwinned

martensite and critical stresses becoming immeasurable.

The NiTi compound forming reaction is highly exothermic and the heat of fusion can

facilitate self-heating up 1000 °C [59]. As Ti reactive to oxygen, oxidation can facilitate the

formation of Ti-rich secondary phases and thus facilitate Ni-enrichment [59]. Frenzel suggests that

as the Ti2Ni phases may absorb small fractions of O, thus becoming Ti2NiOx, with the O content of

each Ti2NiOx phase being below the detection limit for EDS [16]. The black features observed in

the LDED material may be a result of pull-out during the polishing procedure, which would remove

an oxide phase and hinder detection of the phase.


The work demonstrates the PBF and LDED materials exhibit equivalent macroscale shape

memory behavior, despite inherent differences between PBF and LDED techniques that can bring

about microstructural contrasts which give rise to the observed contrasting microscale localized

deformation morphologies.

This chapter focused on the shape memory response for Ti-rich AM alloys. Preliminary

characterization of the TIMT in as built coupons confirmed that the LDED fabricated build

exhibited shape memory behavior. Transformation characteristics were constant away from the

substrate, implying that the substrate affects shape memory behavior at small build heights. This

alloy was confirmed to be martensitic at room temperature, which allowed for characterizing the

shape memory effect response. Multi-scale deformation analysis was used to correlate the

microstructure and material response. Additionally, the LDED response was compared to the

response for specimens fabricated using PBF. The LDED and PBF specimens exhibited complete

shape memory effect recovery. Concentrated meso-scale strains as high as -4% were completely

97

recovered during heating. For the spatially varying LDED alloy microstructure, when stress level

exceeded the critical stress, localized regions of increasing strain contours are bounded and grow

throughout the ROI as twinned martensite converts to detwinned martensite. The prevailing strain

localization facilitates a stark strain hardening-like response compared to PBF. The more uniform

composition and grain structure for the PBF alloys result in the more uniform strain evolution that

is observed. After unloading, the resulting residual martensite is unstable. SME recovery ensues

when heating begins for the LDED alloy. On the other hand, for the PBF material, a finite

temperature increase (i.e. thermal driving force) is required to facilitate SME recovery in the

relatively homogeneous oriented/textured microstructure.

98

Chapter 7

CORRELATING MICROSTRUCTURE AND

SUPERELASTICITY OF DIRECTED ENERGY DEPOSITION

ADDITIVE MANUFACTURED Ni-RICH NiTi ALLOYS


This work builds on previous work [67] in which as built LDED AM Ni-rich NiTi alloys

exhibited complete superelastic recovery due to inherent Ni4Ti3 precipitation during processing and

an aging post-deposition heat treatment enhanced the superelastic response. For this work, post-

processing solutionizing and aging heat treatments were employed to alter the microstructure and

improve the superelastic shape memory responses. This work advances understandings of the

interplay between multiple AM microstructure length scales and stress/thermal-induced martensitic

deformation length scales. This is the first report of the heterogeneous characteristic columnar and

equiaxed grain structures in as built AM NiTi alloys. This work demonstrates Ni-rich NiTi alloy

compositions are advantageous for LDED AM as post processing heat treatments can tailor

precipitation and superelasticity, despite the heterogeneous grain structures. This work has been

published in [65].


7.2.1 Microstructure analysis

Typical microstructures observed in the as built and heat treated (950 °C) conditions are

shown in Figure 7.1. Both Ni4Ti3 and Ni3Ti microconstituent phases exist in the as built materials

and remain after a heat treatment of 10 hours at 950 °C (Figure 7.1 (b)). However, heat treatment

99

of 24 hours is sufficient to dissolve these phases (Figure 7.1 (c)). Using XRD analysis,

microconstituent phases are identified at selected locations along the build height in Figure 7.2.

The presence of microconstituent phases in the as built material has been corroborated in previous

work [67]. For the 950 °C, 10 hour heat treatment, Ni3Ti microconstituents remain, along with trace

amounts of Ni4Ti3. The continued presence of these precipitates indicates that an insufficient time

was used. Increasing the duration to 24 hours dissolved the secondary phases and martensitic (B19'

and R-phase) and B2 austenitic structures are predominant.

(a) (b)

(c)

Figure 7.1 Back scatter electron images showing the microconstituent morphologies for (a)

as built alloys and alloys heat treated at 950 °C for (b) 10 h and (c) 24 h durations. In (a)

and (b), lenticular microconstituents are Ni4Ti3 precipitates and Ni3Ti secondary phases

appear globular.

The composition is homogenized by dissolving secondary phases. The composition was

measured in multiple location using EDS, on three specimens: as built, heat treated at 950 °C for

10 hours, and solution treated at 950 °C for 24 hours. The average composition measurements are

100

Figure 7.2 XRD scans with increasing post-deposition heat treatment duration at 950 °C.

The specified locations are along the build height locations. Phases have been identified as

B2 ; B19' ; R-phase ; Ni4Ti3 ; and Ni3Ti .

Figure 7.3 Evolution of composition with increasing post deposition heat treatment

duration at 950 °C. The dashed horizontal line is the input powder feedstock composition.

Circles represent average compositions. Squares represent maxima and minima.

shown in Figure 7.3 with the standard deviation, maximum, and minimum measurements. The

largest composition range (difference between maximum and minimum measurements) and highest

101

standard deviations are observed for the as built material. The large standard deviation measured

for the as built material may be a result of segregation during solidification [115] or non-uniform

precipitation of Ni-rich secondary phases drawing Ni out of the matrix [119]. The most uniform

composition is measured for the 950 °C, 24 hour heat treatment. This temperature and time are

sufficient to dissolve secondary phases and produce a uniform composition. Therefore, the optimal

solutionizing treatment for these alloys is 950 °C, 24 hours.

The grain morphologies for as built and solution treated materials are shown in Figure 7.4

(a) and Figure 7.5 (a), respectively. Representative columnar grains having a high aspect ratio

(length (z) / width (x)) are elongated in the build height direction and identified. These columnar

grains are expected to form nearest the center of each deposited pass, which corresponds to the

direction of the highest thermal gradient [120]. The regions of columnar grains are separated by a

horizontal distance of approximately 1.9 mm, which is equal to the hatch spacing. A region of large

equiaxed grains, which have aspect ratios approaching 1, is also identified between the columnar

grain regions. Regions with finer equiaxed subgrain structures (100 μm to 1.4 mm) also exist.

Larger equiaxed grains result from the remelting in the overlap regions between passes [58,121].

The regions of small equiaxed subgrain structures are separated by a vertical distance of 0.6 mm,

which corresponds to the layer thickness. Thus subgrains form in the layer overlap regions

[115,122].

The corresponding grain orientation maps are shown in Figure 7.4 (b) and Figure 7.5 (b),

and the pole figures are shown in Figure 7.4 (c) and Figure 7.5 (c). For the as built specimen, there

is a slight preferred orientation. Typically, cubic AM alloys exhibit a preferred orientation, due to

the directional solidification of the AM process [120]. For the solutionized alloy, however, there is

a random B2 austenitic grain orientation. The LDED AM fabrication technique produces a

characteristic grain morphology generally without texture. Recrystallization and grain growth are

expected at the high temperatures prevalent in the solution treatment [123,124]. Alternatively, the

102

grain structure can be altered by changing the build plan during the AM fabrication process [58].

Measuring volume fractions of recrystallized equiaxed

(a) (b)

(c)

Figure 7.4 (a) Grain boundary structure, (b) grain orientation map, and (c) pole figure for

an as built Ni-rich compression specimen. The grain orientation is identified as normal the

specimen surface.

substructures or grain boundary reorientation could characterize the degree/extent of

recrystallization [124]. Ultimately this work shows that the solution treatment is unable to alter

those characteristic attributes of columnar grains coexisting with equiaxed and subgrain structures.

103

(a) (b)

(c)

Figure 7.5 (a) Grain boundary structure and (b) grain orientation map for a solution

treated (950 °C, 24 h) specimen. (c) Corresponding pole figure.

Solution treated specimens were subsequently aged, and sizes and area fractions of the

Ni4Ti3 precipitate morphologies were compared to directly aged materials in Figure 7.6. Inset SEM

images in Figure 7.6 (a) illustrate contrasting height dependencies. The directly aged alloy has the

largest variation in average precipitate size and standard deviation at each height. The solution

treated and aged material, on the other hand, displayed a more homogeneous Ni4Ti3 precipitate

morphology. In contrast to the aged microstructures, the morphology is isotropic with equivalent

average sizes and inconsequential standard deviations with changes in height. Precipitation is

104

(a)

(b)

Figure 7.6 Spatial resolution of the lenticular Ni4Ti3 precipitate morphology. Build height

(in the z-direction) dependence of the Ni4Ti3 precipitate morphology as a result of

solutionizing and aging (SL+Aged) versus directly aging as built material (Aged). (a) Area

fraction with selected SEM images inset and (b) Length along the major axis of the

lenticular precipitate.

105

governed by presence of grain boundaries or defects and the degree of supersaturation of Ni [37].

The 53 at.% Ni composition is supersaturated and uniform for the solution treated material. Due to

the high Ni content, relatively homogenous precipitation of Ni4Ti3 occurs despite the complex

heterogeneous grain structure in the AM alloys. The precipitate morphology length scale, with

precipitate sizes less than 5 μm and interparticle spacing of 0.5 ±0.1 μm, is finer than the grain

structure, with grain sizes ranging from 100 μm to 1.4 mm. Thermo-mechanical characterization is

carried out to correlate shape memory behavior with the various precipitate morphologies.

Consequences of aging the as built alloy versus aging after solution treatment are elucidated.

7.2.2 Thermal-induced and stress-inducted martensitic transformation behavior

Differential scanning calorimetry (DSC) analysis provides insights into the thermal-

induced MT without external load biasing the phase evolution. Contrasting DSC thermo-grams in

Figure 7.7 expose differential underlying TIMT after aging the solution treated material compared

to that observed in the as built alloy. The homogeneous/isotropic microstructure for the

solutionized and aged material produces distinct cooling peaks (Figure 7.7 (a)) and a single heating

peak (Figure 7.7 (b)). Conventionally processed NiTi alloys which have undergone a thermo-

mechanical aging treatment typically exhibit a homogeneous distribution of Ni4Ti3 precipitates and

show similar responses to the responses observed here [101]. It is well accepted that the two

separable cooling peaks are indicative of B2 → R → B19' and the single heating peak represents

B19' → B2 [1,101]. For the aged material, the cooling (Figure 7.7 (a)) and heating (Figure 7.7 (b))

thermo-grams exhibit amalgamated peaks over broad forward and reverse phase transformation

temperature ranges. The marked difference from the solution treated and aged material response is

due to microstructural heterogeneities including composition gradients and heterogeneous Ni4Ti3

106

morphologies [101,103]. Overlapping peaks can be attributed to concomitant spatially varying

thermodynamic transformation driving forces. Likewise, the differential precipitate morphologies

are expected to produce varying SIMTs.

(a) (b)

Figure 7.7 (a) Forward and (b) reverse thermal-induced martensitic transformations for

directly aged (550 °C, 3h) and solutionized and aged (950 °C, 24 h followed by 550 °C, 3h)

alloys.

Comparing Figure 7.8 (a), Figure 7.9 (a), Figure 7.10 (a), and Figure 7.11 (a) reveals

dissimilar SIMTs underpinning the compressive stress-strain responses across the as built,

solutionized, directly aged, and solution treated and aged conditions. Austenite elastic deformation

is distinguished by the elastic modulus 𝐸𝐴, and the critical transformation stresses 𝜎𝐴→𝑀

characterizes the onset of the SIMT. In the absence of precipitates, the highest 𝜎𝐴→𝑀 and 𝐸𝐴 are

measured for the solution treated material (Figure 7.9 (a)). For Ni-rich concentrations, it is well

known that Ms temperatures decrease with increasing Ni concentration [16], thus the Ms

temperature is estimated to be below -200 °C for the 53 at.% Ni concentration of the solution treated

material. It is reasonable that the test temperature exceeds the superelastic window [91] for this Ni-

rich concentration. Hence, martensite is strain-induced so that superelastic recovery is not observed.

For as built (Figure 7.8 (a)), aged (Figure 7.10 (a)), and solution treated and aged (Figure 7.11 (a))

specimens, the presence of Ni4Ti3 precipitates shifts the superelastic window to include the test

temperature. Internal stress fields of precipitates can promote the SIMT and decrease 𝜎𝐴→𝑀 and

𝐸𝐴 [125,126]. The stress-strain curves for aged (Figure 7.10 (a)) and solution treated and aged

107

(a)

εaxial =

-0.1% -0.2 -0.5 -0.7 -1.0 -1.4 -2.1 -2.6 -3.0 -3.3

-3.1 -2.8 -2.5 -2.1 -1.4 -1.0 -0.6 -0.3 -0.2

(b)

Figure 7.8 (a) Compressive stress-strain response for an as built alloy. The austenitic elastic

modulus EA is defined on the loading curve. The 0.2% offset critical transformation stress

(σA→M) and stress at the onset of elastic unloading are illustrated and specified along the

ordinate axis. The stress hysteresis (∆𝜎) is the differential between σA→M and that onset stress.

(b) Full-field axial strain contour images showing meso-scale strain evolutions. Above each

image are macro-scale strain (%) values, which are identified by the points along the stress-

strain curve in (a).

(Figure 7.11 (a)) materials exhibit a distinct “knee” near the 𝜎𝐴→𝑀, which typically precedes that

stress plateau of the flag shape superelastic stress-strain responses [37,39]. The stress hysteresis ∆𝜎

and 𝜎𝐴→𝑀 measurements are comparable for the as built (Figure 7.8 (a)) and aged (Figure 7.10 (a))

alloys. More uniform precipitate morphologies in the solution treated and aged specimen (Figure

7.6) bring about the smallest ∆𝜎 and 𝜎𝐴→𝑀. Differential macro-scale property measurements

108

underscore that the SIMT pathway depends on whether an as built alloy is aged directly or

following a solution treatment.

(a)

εaxial =

-0.3% -0.6 -1.0 -1.3 -1.6 -2.2 -2.9 -3.3 -3.8 -4.2

-4.1 -3.9 -3.7 -3.6 -3.4 -3.1 -2.9 -2.7 -2.3

(b)

Figure 7.9 (a) Compressive stress-strain response for a solutionized (950 °C, 24 h) Ni-rich

alloy. The austenitic elastic modulus EA, critical transformation stress (σA→M) and stress at

the onset of elastic unloading are illustrated and specified along the ordinate axis. The stress

hysteresis (∆𝜎) is the differential between σA→M and that onset stress. (b) Full-field strain

contour images showing meso-scale strain evolutions. Above each image are macro-scale

strain (%) values, which are identified by the points along the stress-strain curve in (a).

109

(a)

εaxial =

-0.1% -0.2 -0.6 -0.9 -1.2 -2.0 -2.5 -3.2 -3.7 -4.0

-3.8 -3.2 -2.4 -2.0 -1.3 -0.7 -0.4 -0.2 0

(b)

Figure 7.10 (a) Compressive stress-strain response for an aged (550 °C, 3 h) alloy. The

austenitic elastic modulus EA is defined on the loading curve. The 0.2% offset critical

transformation stress (σA→M) and stress at the onset of elastic unloading are illustrated and

specified along the ordinate axis. The stress hysteresis (∆𝜎) is the differential between σA→M

and that onset stress. (b) Full-field strain contour images showing meso-scale strain

evolutions. Above each image are macro-scale strain (%) values, which are identified by the

points along the stress-strain curve in (a).

110

(a)

εaxial =

-0.1% -0.2 -0.5 -0.8 -1.1 -1.8 -2.5 -2.9 -3.3 -4.0

-3.9 -3.5 -3.0 -2.0 -1.2 -1.0 -0.8 -0.4 -0.2

(b)

Figure 7.11 (a) Compressive stress-strain response for a solutionized and aged (950 °C, 24 h

followed by 550 °C, 3 h) alloy. The austenitic elastic modulus EA, critical transformation stress

(σA→M) and stress at the onset of elastic unloading are illustrated and specified along the

ordinate axis. The stress hysteresis (∆𝜎) is the differential between σA→M and that onset stress.

(b) Full-field strain contour images showing meso-scale strain evolutions. Above each image

are macro-scale strain (%) values, which are identified by the points along the stress-strain

curve in (a).

Finer scale full field deformation analyses are employed in order to discern underlying

SIMTs in the variable Ni4Ti3 microstructures. DIC strain contours in Figure 7.8 (b), Figure 7.9 (b),

Figure 7.10 (b), and Figure 7.11 (b) depict the evolution of the phase transformation as the growing

volume fraction of material undergoing the SIMT interacts with the AM microstructure. Typically,

distinct boundaries are observed between austenite and the transformed martensite [39,127,128].

However, for the solution treated material in Figure 7.9 (b), the phase transformation evolves with

111

εaxial =

-0.1% -0.2 -0.5 -0.7 -1.0 -1.4 -2.1 -2.6 -3.0 -3.3

-3.1 -2.8 -2.5 -2.1 -1.4 -1.0 -0.6 -0.3 -0.2

Figure 7.12 Full-field transverse strain contour images showing meso-scale strain evolutions

for the as built Ni-rich alloy. Above each image are macro-scale axial strain (%) values, which

are identified by the points along the stress-strain curve in Figure 7.8 (a).

εaxial =

-0.3% -0.6 -1.0 -1.3 -1.6 -2.2 -2.9 -3.3 -3.8 -4.2

-4.1 -3.9 -3.7 -3.6 -3.4 -3.1 -2.9 -2.7 -2.3


for the solution treated alloy (950 °C, 24 h). Above each image are macro-scale axial strain

(%) values, which are identified by the points along the stress-strain curve in Figure 7.9 (a).

large areas, represented by encircled regions, exhibiting diffuse contours. Within diffuse regions,

strains increase in a seemingly uniform manner, i.e. phase transformation fronts are

indistinguishable. Localized strain concentrations are marked within the diffuse contours. The

concentrations appear to grow during loading and shrink during unloading. Moreover, austenite

regions remain untransformed throughout deformation, which can be attributed to the phase

112

transformation being strain-induced. For the as built alloy, the interaction between the SIMT and

the anisotropic/heterogeneous Ni4Ti3 microstructure causes strain contours to evolve in a diffuse

manner as illustrated in Figure 7.8 (b). Similar strain contours are observed for the transverse strain

response, shown in Figure 7.12 and Figure 7.13 for the as built and solutionized alloys, respectively.

εaxial =

-0.1% -0.2 -0.6 -0.9 -1.2 -2.0 -2.5 -3.2 -3.7 -4.0

-3.8 -3.2 -2.4 -2.0 -1.3 -0.7 -0.4 -0.2 0


for the directly aged specimen (550 °C, 3 h). Above each image are macro-scale axial strain

(%) values, which are identified by the points along the stress-strain curve in Figure 7.10 (a).

εaxial =

-0.1% -0.2 -0.5 -0.7 -1.0 -1.4 -2.1 -2.6 -3.0 -3.3

-3.9 -3.5 -3.0 -2.0 -1.2 -1.0 -0.8 -0.4 -0.2


for the solution treated and aged specimen (950 °C, 24 h followed by 550 °C, 3 h). Above each

image are macro-scale axial strain (%) values, which are identified by the points along the

stress-strain curve in Figure 7.11 (a).

113

Transformation fronts arise when the SIMT takes place in the aged microstructures. For

the directly aged alloy, the SIMT produces multiple transformation fronts delineated (dotted lines)

in the image at the -0.9% strain during loading in Figure 7.10 (b). The reverse transformation

transpires with fronts identified at -2.0 and -1.3 % strains during unloading. The transverse strain

contours are similar (Figure 7.14). On the other hand, the axial and transverse contours for the

solution treated and aged material reveal a single transformation front in Figure 7.11 (b) and Figure

7.15. An interface moves across the ROI and the area fraction of transformed martensite increases.

At the maximum macro-scale strain of -4%, the interface has traversed the ROI. The strain gradient

behind the interface exhibits the largest micro-scale value of -6%. The single interface/front SIMT

facilitates the lowest critical stress level, which is advantageous for ensuring austenite transforms

prior to slip. The single front motion/path underlying the reverse MT during unloading is the

reverse of the forward path during loading.

The single front can be attributed to the SIMT in the more homogeneous precipitate

morphology. Single fronts are typically observed for NiTi SMAs deformed in tension [128–130]

and is commonly described as Lüders-like deformation [129]. The SMAs are typically textured

extruded bars, sheets, or drawn wires [127,131–135] with oriented and refined grain structures

begetting preferential martensite orientation. The nature of transformation front evolutions is

typically rationalized considering the impact texture has on the SIMT [127,128]. If grain structure

controlled the SIMT, the heterogeneous/untextured nature of the AM structure would not be

expected to produce a single planar transformation front. For the current alloys, the precipitate

length scale is much closer, than the grain scale, to the tens of nanometer scale for martensite plate

widths [1,37]. Thus, it is reasonable to assume local stress fields associated with precipitates will

preferentially orient martensite and that the precipitate morphology is the microstructure length

scale dictating the SIMT and shape memory behavior for the aged LDED AM alloys. In the most

homogeneous Ni4Ti3 morphology of the solutionized and aged material, uniform internal stress

114

fields should arise that preferentially orient martensite [43]. Note the inclination of the planar front

with respect to the loading axis matches the commonly observed 55 ° angle [128,129]. It has been

postulated that the angular measure minimizes the strain incompatibility at the macroscopic

interface between elastically deformed austenite and the transformed martensite [129]. Even finer

scale and higher magnification full-field deformation analysis can expose the SIMT interactions

with grain and precipitate length scales to further investigate this hypothesis.

The material properties for the as built, solutionized, directly aged, and solutionized and

aged specimens are listed in Table 7.1. The compressive material properties for other NiTi alloys,

fabricated using conventional methods (extruded, hot rolled, cold-drawn) or other net-net shape

processes (cast, powder metallurgy) are also listed. The conventionally processed NiTi alloys have

undergone a thermo-mechanical processing step, producing an optimized microstructure and shape

memory response. These materials have the highest macro-scale recovery strains; AM alloys cannot

compare to conventionally processed NiTi alloys; rather AM is seen as an alternative near-net shape

fabrication process. Cast alloys have a spatially varying grain structure consisting of elongated

grains and equiaxed grains [15,83,124,136]; this microstructure is the most similar to the observed

AM microstructure. Comparing recovery strains for all of the AM alloys to that of the cast alloy,

the AM alloys are able to recover larger superelastic strains. Additive manufacturing of NiTi thus

shows promise for fabricating alloys which exhibit good superelastic behavior and the two-step

heat treatment process (solutionizing then aging) shows promise for being able to tailor the

microstructure and thus the superelastic response.

115

Table 7.1 Calculated material and shape memory properties for the as built, directly aged, solutionizied, and solution treated and

aged compressive responses.

Young’s

modulus

Poisson’s

ratio

Critical

transformation stress

Stress

hysteresis

Macro-scale

recovered strain

Meso-scale

recovered strain

EA (MPa) νZX σA→M (MPa) Δσ (MPa) (%) (%)

As built 87 0.42 470 280 -3.1 -4.2

Directly Aged

(550 °C, 3 h) 83 0.49 530 370 -4.0 -6.0

Solutionized

(950 °C, 24 h) 100 0.46 890 - - -

Solution treated + Aged

(950 °C, 24 h + 550 °C, 3 h) 78 0.49 270 160 -3.9 -6.0

Extruded (Ni50.8Ti49.2 at.%) 62 - 375 -5.4 -

Hot Rolled [83] 87 - 500 320 -3.7 -

Cold-drawn [83] 52 - 550 520 -3.7 -

Cast [83] 37 - 200 180 -1.0 -

Powder metallurgy

(press & sinter) [117] 71 - 370 140 -1.3 -

116


A micro-scale deformation analysis was employed to investigate the influence of

solutionizing and aging heat treatments on the superelastic behavior and martensitic transformation

morphologies in AM fabricated NiTi alloys. Due to the complex thermal histories caused by the

layer by layer deposition process, spatially varying grain and precipitate structures across a range

of size scales exist in the as built material. It was found that a solutionizing heat treatment of 24

hours was required to homogenize composition and dissolve undesired secondary phases otherwise

present in the as built material. Macro-scale stress-strain responses were insufficient to explain the

consequences of the SIMT taking place in the different post deposition heat treated microstructures.

Full-field deformation analysis, however, exposed the importance of solution treating prior to aging

in obtaining superelastic behavior and the martensitic transformation response. A characteristic

grain structure consisting of equiaxed and columnar grains, which correspond to the layer thickness

and hatch spacing used in the AM build plan, exists in the as built condition. A solutionizing heat

treatment at 950 °C for 24 hours produces a uniform B2 microstructure. However, the solution

treatment is not able to noticeably alter the as built grain structure. Aging a properly solutionized

alloy with supersaturation of Ni produces a relatively uniform Ni4Ti3 precipitate morphology; thus

overcoming the otherwise spatially varying morphology expected from aging the inherently

heterogeneous grain structure produced by LDED. In contrast to the directly aged case, the

relatively uniform Ni4Ti3 morphologies produced by solutionizing and aging results in improved

shape memory responses: distinct DSC peaks corresponding to B2 → R → B19' and B19' → B2

and reversible interface motion underpinning the superelastic response.

117

Chapter 8

Ni-CONCENTRATION DEPENDENCE OF DIRECTED ENERGY

DEPOSITED NiTi ALLOY MICROSTRUCTURES


During the laser-based additive manufacturing process, localized re-melting and re-heating

in regions between overlapping passes and layers drive localized microstructure development. Our

previous work, in references [65,67], began to characterize the spatial variation in the precipitate

and grain structure for NiTi alloys fabricated using LDED. This work advances understandings of

the microstructure which results from the layer-by-layer deposition process of AM. Three

orthogonal planes were analyzed, providing a three-dimensional analysis of the microstructure

resulting from overlapping passes, overlapping layers, and a combination of overlapping passes

and layers. The findings support the idea that similar grain morphologies result within regions

where previously deposited material is re-melted and re-heated as new material is deposited, for

both Ti-rich and Ni-rich feedstock compositions. This is the first report in which the microstructure

of LDED NiTi alloys is probed in interfacial regions formed by overlapping passes or layers and

outside of these regions. This work was published in [137].


Samples for microstructure characterization were sectioned, as shown in Figure 8.1 (a),

Figure 8.1 (b), and Figure 8.1 (c), to provide a 3-dimension visualization of the microstructure. The

grain and precipitate morphologies were characterized in interfacial regions and away from the

interface (i.e. within the bulk) regions. Interfacial regions are defined as areas of previously

deposited material that are re-melted or re-heated as new material is deposited. For the X-Y plane,

118

which spanned by the pass-build directions, the overlap that produces the interfacial regions is

illustrated in Figure 8.1 (a). Interfacial regions on the Y-Z plane correspond to the build-height

directions and capture the interlayer regions (Figure 8.1 (b)). For the X-Z plane, corresponding to

the pass-height directions, interfacial regions in Figure 8.1 (c) represent both pass overlaps and

interlayers.

(a)

(b)

(c)

Figure 8.1 Schematics of interfacial regions formed by overlapping passes and/or layers.

(a) Build coupon XY plane spanned by pass and build directions with periodic interfacial

regions formed by overlapping passes. (b) Build coupon YZ plane spanned by build and

height directions with periodic interfacial regions formed by overlapping layers. (c) Build

coupon XZ plane spanned by pass and height directions with interfacial regions formed by

overlapping passes and layers.

Figure 8.2 and Figure 8.3 respectively show the grain structures for the Ni-rich and Ti-rich

alloys. Figure 8.2 (XY) and Figure 8.3 (XY) correspond to overlapping passes. The region labeled

(a) consists of elongated grains. Grains are elongated in the highest heat flow direction, which

corresponds to the pass direction, which is noted as the x-direction in the figures [74,120]. The

region labeled (b) consists of equiaxed grains. Grain sizes are approximately 0.2 mm for both

119

compositions. Regions of equiaxed grains are alternated with regions of elongated grains. The

distances in the x-direction between the regions of equiaxed grains is similar to the hatch spacing,

which is approximately 1.8 mm for the Ni-rich alloy and 1.7 mm for the Ti-rich alloy. The small

equiaxed grains are situated on the edges of the pass [138], which corresponds to areas with the

fastest cooling rates [61,74,115,139].

Figure 8.2 XY, YZ, and XZ cross-sectional views of the Ni-rich NiTi alloy build. (XY)

Elongated/columnar grains are identified in the (a) bulk and equiaxed grains exist in the

(b) interpass/interfacial region. (YZ) Elongated grains exist in the region (a) and away from

the interlayer/interfacial regions. (XZ) Columnar grains exist within and adjacent to

interfacial (both interpass and interlayer) regions (b)

Grain structures resulting from overlapping layers and the resulting interlayers are shown

in Figure 8.2 (ZY) and Figure 8.3 (ZY). Fusion lines are observable and delineate the lower

boundary of a deposited layer, and the distance between fusion lines corresponds to the 0.7 mm

layer height. The uniform elongated grain structure is oriented along the build height and consistent

X

Y

(a)

(b)

1 mm

Elongated

Equiaxed

Elongated

1 mm

(a)(b)

Z

Y

1 mmZ

X

Columnar

Equiaxed(b)

(a)

120

with the direction of highest heat flow [74,120]. The corresponding major axis lengths, in the z-

direction, are nearly 1.0 mm for the Ni-rich and Ti-rich alloys. Elongated grains can span multiple

layers, suggesting that there is some epitaxial growth upon the previously solidified layer as new

material is deposited [74].

Figure 8.3 XY, YZ, and XZ cross-sectional views of the Ti-rich NiTi alloy build. (XY)

Elongated/columnar grains are identified in the (a) bulk and equiaxed grains exist in the

(b) interpass/interfacial region. (YZ) Elongated grains exist in the region (a) and away from

the interlayer/interfacial regions. (XZ) Columnar grains exist within and adjacent to

interfacial (both interpass and interlayer) regions (b)

Pass overlap and interlayers produce the grain structures shown in Figure 8.2 (ZX) and

Figure 8.3 (ZX). The characteristic semi-circle shape of pass boundaries [19,20] is outlined. Note

that distances in the x-direction between the bottom vertices is marked. For the Ni-rich and Ti-rich

compositions, the distances are similar to the hatch spacing (1.3 mm), whereas the z-distance (0.7

(a)

(b)

1 mm

Elongated

Equiaxed

X

Y

Elongated

1 mm

(a)

(b)

Z

Y

Columnar

Equiaxed

1 mmZ

X

(a)

(b)

121

mm) between these vertices is comparable to the layer thickness. Away from the boundaries, larger

columnar grains can be distinguished at locations (a). Smaller equiaxed exist along and near the

boundaries marked with (b). The length of the columnar grains is about 0.7 mm and they are

oriented normal to the boundaries. Orientations generally align toward the highest heat flow

direction.

(a) (b)

(c)

Figure 8.4 SEM images showing precipitate morphologies for the Ni-rich alloy (a) within

an interfacial region and (b) in the bulk. (c) SEM images of the Ti-rich alloy

Interfacial regions exhibit the representative precipitate morphologies shown in Figure 8.4

(a) and Figure 8.4 (b) for Ni-rich compositions. In general, precipitates do not readily form in the

additive manufactured Ti-rich materials, and none is observed here (Figure 8.4 (c)). Very small

Ni4Ti3 precipitates exist in Ni-rich alloys at locations within and nearest the interfacial regions, e.g.

1.0 μm in Figure 8.4 (a). Precipitates increase in size away from the boundary in the bulk, e.g. 2.3

μm in Figure 8.4 (b). Oliveira et al. [68] observed Ni4Ti3 precipitates in a Ni-rich alloy within a

122

heat affected zone (HAZ) and contrasted the HAZ with the center of the pass, which did not contain

precipitates. The variation in the secondary phase formation was attributed to temperature and rapid

solidification rates at the center of the pass precluding precipitation.


The additive manufacturing process results in localized regions which undergo re-melting

and re-heating. These interfacial regions are formed when passes and layers are overlapped.

Sectioning the build coupons in three orthogonal cross sections provided a 3-dimensional view of

the grain and precipitate morphologies in the interpass and interlayer regions. Within the interfacial

regions, where the material is re-melted and re-heated, the grain and precipitates sizes are smaller

relative to the bulk. While these columnar grains showed preferential orientation in the highest heat

flow direction corresponding to the build height, the fine equiaxed grain structures in the interfacial

regions showed no orientation preferences. In addition to these changes in grain structures,

precipitation of second phases is also impacted in these regions. The remelting and reheating

characteristic of the interfacial regions produces precipitates in the Ni-rich alloys with fine oriented

morphologies different from those observed in the bulk regions. The interfacial region

microstructure is different from the bulk microstructure. As interfacial regions can constitute nearly

25% of an AM build, the important finding of microstructure variation between interfacial and bulk

regions will directly impact the shape memory response.

123

Chapter 9

SUMMARY AND CONCLUSIONS

This dissertation characterizes multiple microstructure length scales and employs a multi-

scale strain analysis approach in order to characterize the shape memory response of additive

manufactured NiTi shape memory alloys. Elementally blended powders were utilized in

conjunction with laser-based directed energy deposition to fabricate build coupons. The build

coupons were sectioned and the microstructure and shape memory behavior were spatially

resolved. From this work, the following conclusions can be drawn:

• The LDED layer-by-layer deposition process produced a spatially varying grain and

precipitate structure which resulted in a spatially varying shape memory response. The

finer grain and precipitate structures are present within the interfacial regions relative to

the bulk regions, for both Ti-rich and Ni-rich compositions.

• Post-deposition heat treatments improved the precipitate microstructure and in turn the

superelastic response. Precipitation of Ni-rich secondary phases readily occurs within the

as built Ni-rich alloy, though the precipitate size and area fraction vary significantly with

build height. Post-depositions heat treatments were able to tailor the precipitate

microstructure for the Ni-rich builds, with solutionizing and aging treatments producing

the most uniform precipitate morphologies, despite the complex grain structure. The

precipitate morphology appears to control the behavior.

124

Chapter 10

RECOMMENDATIONS FOR FUTURE WORK

The current work has made progress towards developing relationships between the

fabrication process, microstructure, and shape memory behavior for additive manufactured NiTi

shape memory alloys. To further advance the understanding of AM NiTi, the following research

directions are proposed:

10.1 Designing microstructure via process parameter control

Additive manufacturing has been envisioned as a tool for material design based on the idea

that process parameters which are variable in AM (but not in traditional fabrication) can be

leveraged for achieving a desired microstructure. The proposed future work would focus on

understanding how the grain structure and precipitate morphology are altered or controlled by

varying the additive manufacturing process parameters. A limitation of this current dissertation

work is that all build coupons were fabricated using the same process parameters and the same

build plan. The processing parameters of interest in the proposed research would be the parameters

related to energy density (laser power, scan speed, hatch spacing, layer thickness) and the scan

strategy. Other existing works have started correlating the process parameters to the microstructure,

primarily focusing on changing the laser power or the scan speed. Controlling the laser power and

scan speed controls the melt pool size and temperature [28]. The trend of increasing energy density

resulting in increased grain size is reported for both LDED [55,60,61,113] and PBF [57,62]

fabricated alloys. This relationship has been attributed to the increased energy input resulting in

either a higher temperature or a slower solidification rate, which allows grains to coarsen

[55,60,62]. This phenomena extends to NiTi alloys fabricated using micro direct metal deposition

125

[63]; the fine grain size was attributed to the relatively small energy input resulting in a small melt

pool size. The hatch spacing and layer thickness parameters would affect the degree of overlap and

the interfacial region dimensions of the build. A typical degree of overlap is 25% [28]. This

dissertation work, which is the first to probe the microstructure within the interfacial regions and

the bulk regions, found that there is a microstructure variation between the interfacial and bulk

regions. Specifically, that finer grain sizes and precipitate morphologies (for the Ni-rich alloy) are

observed within the interfacial regions. A study which characterizes the microstructure resulting

from varied processing parameters would significantly advance the understanding of

microstructures resulting from the additive manufacturing fabrication process.

10.2 Tension-compression asymmetric response of AM NiTi alloys

The asymmetric response between the tension and compression stress states is known to

exist for NiTi shape memory alloys [41–43,134,140,141]. This asymmetric response arises as

different martensite variants are activated or preferred for different stress states [40,43]. However,

this dissertation work does not characterize the tension-compression asymmetry. As the loading

axes of the tension specimens (build-/y-direction) are not the same as the loading axes of the

compression specimens (height-/z-direction), this comparison should not be made for this work.

Additionally, the tensions specimens fractured at low strains, and the superelastic response

typically could not be characterized (as fracture occurred, rather than strain recovery). Post-

deposition heat treatments, which improved the strain recovery in compression specimens, may aid

in increasing strain recovery in the tension specimens.

Recently, progress has been made by the Multifunctional and Adaptive Materials

Laboratory on characterizing shape memory response for tension and compression specimens

oriented in the build-/y-direction. These specimens were extracted from Ni-rich build coupons

126

fabricated using 3 pass, 14 layers, as shown in Figure 10.1. This study will fill a current gap in a

characteristic phenomenon of NiTi shape memory alloys. A proposed future work would continue

this asymmetry investigation, while additionally studying the effectiveness of post-deposition heat

treatment in altering the degree of asymmetry.

(a) (b)

Figure 10.1. Plan for extracting tension and compression specimens with the same loading

axis direction (a) Tension specimens with the width of the gage section parallel to the layers.

(b) Tension specimens with the width of the gage section perpendicular to the layers (Build

coupon IDs: B49, B50).

10.3 Precipitate morphology microstructure length scale dictating SIMT

From this dissertation work, different strain contour evolutions were observed for the

solutionized and aged specimen, which had an isotropic precipitate morphology, relative to the as

built and directly aged specimens, which had anisotropic precipitate morphologies. Additionally,

the characteristic grain features were not altered by heat treatment. These findings led to the

hypothesis that the Ni4Ti3 precipitate morphology, rather than the grain structure, dominates the

MT morphology and shape memory behavior. To test this hypothesis, one can analyze AM alloys

which have different grain structures and similar precipitate morphologies. The similar precipitate

morphologies can be achieved through solutionizing and aging heat treatments.

As the grain structures for the pass-build (XY) plane, pass-height (XZ) plane, and build-

height (YZ) plane have different grain structures, orthogonal surfaces of a compression specimen

127

will also have different grain structures. In this dissertation work, planes were examined

individually. However, the strain contour evolutions for different grain structures can be

simultaneously investigated using a 2-camera 2D DIC set-up. In this setup, two separate machine

vision systems are set-up to capture images of orthogonal planes during the same thermo-

mechanical experiment. By examining orthogonal planes simultaneously, the effect of the different

grain structures on the superelastic behavior can be investigated. If the precipitate morphology

dominates the SIMT, then the complex grain structures inherent to AM would have minimal effect

on the response. Thus, future research could focus on tuning the precipitate morphologies to

achieve the desired behavior and optimize the superelastic response.

10.4 Functionally Graded NiTi Structures

An advantage of using LDED AM systems is the ability to change the feedstock

composition during deposition, to fabricate functionally graded materials. The purpose of

fabricating functionally graded NiTi materials is to engineer the composition and microstructure in

order to enable spatially on demand superelastic behavior. Functionally graded builds have been

created in previous works, as briefly described in this section, but have yet to be characterized. The

objectives for this previous work were to successfully fabricate functionally graded NiTi structures

(which has been completed) and investigate the local shape memory response and microstructure

within these structures, and at the interfaces between the different materials (which has not yet been

completed).

The first step in this work was fabricating functionally graded NiTi materials. These

structures were fabricated using the high power-high deposition (HPHD) laser-based directed

energy deposition system, shown in Figure 10.2. The deposition head is shown in Figure 10.2 (b).

128

The deposition head has four powder nozzles (nozzle diameter: 1.65 mm) which are angled at 30 °

with respect to the vertical, as shown in Figure 10.2 (c).

(a)

(b) (c)

Figure 10.2 (a) Image shows the High Power-High Deposition (HPHD) additive

manufacturing system. The HPDP system is within the glove box enclosure. (b) Deposition

head with four coaxially-fed powder nozzles. The hot plate and substrate and the build

coordinate axes are shown. (c) Close-up of the four powder nozzles.

In a manner similar to this dissertation work, the processing parameters were optimized for

the deposition system and powders. Three powders were used in this work: elemental Ti powder,

elementally blended Ni53Ti47 wt.%, and Ni58Ti42 wt.%. The mass flow rates were measured, as

129

shown in Figure 10.3. The NiTi powders had similar flow rates for the same voltage, whereas the

Ti powder flowed at a slower rate.

Figure 10.3 Measurement of powder mass flow rate of the feedstock powder, based on the

voltage (dial) reading of the powder feeder for the elemental Ti and elementally blended

NiTi powders.

The processing parameters for the depositions are listed in Table 10.1. The width and

thickness of a single pass was measured for the three powder compositions, to establish the layer

thickness and hatch spacing. The layer height and pass width were similar for the different powders

(heights measured as 0.5 mm for Ti powder and 0.6 mm for NiTi powders; widths measured as 3.2

mm for Ti powder, 3.1 for Ni53Ti47 wt.% powder, and 3.2 mm for Ni58Ti42 wt.% powder). The layer

thickness was set equal to the layer height and the hatch spacing was set equal to half of the pass

width. The same hatch spacings and layer thicknesses were used for depositing all powder

compositions. The travel path of the laser was unidirectional within a layer (i.e. started on the same

side) and alternating layers were deposited in opposite directions. The scan velocity was set as 10.1

mm/s, except for the last pass in the layer, which was set at 8.4 mm/s. This was based on the

suggestion of CIMP-3D staff, as a slower velocity on the last pass of a layer tends to result in

straighter build coupons.

Builds were fabricated using either a single powder composition or two powder

compositions. The builds fabricated using a single powder composition were fabricated to use as

130


the high power-high deposition (HPHD) laser-based directed energy deposition system.

Parameter Value

(Laser) Energy

Source

Parameters

Laser wavelength 1070 nm

Laser power 1000 W

Spot size (at part) 4 mm

Position of focal point 212 mm

Material

Parameters

Base Material Alloy CP Ti

Filler Material Alloy elemental Ti powder;

elementally blended Ni and Ti powders

(Ni53Ti47 wt.% or Ni58Ti42 wt.%)

Feedrate Ti powder: 8.5 g/min.

NiTi powders: 10 g/min.

Powder characteristics described in 2.2.1

Environmental

Parameters

Chamber Gas Ar

Supplemental Gas/Flow rate Ar; 9.4 L/min.

Vacuum Level 2.2 mbar

Manipulation

and Toolpath

Travel speed 10.1 mm/s.

Layer height 0.5 mm

Hatch spacing 3.2 mm

Other Temperature 250 °C (set on hot plate beneath substrate)

Dwell Times 0 s between passes, 0 s between layers

test coupons for obtaining representative microstructure and behavior. The builds fabricated using

two powder compositions were fabricated either in an A-B pattern (powder composition B

deposited on top of powder composition A) or in an A-B-A pattern (powder A deposited on powder

B deposited on powder A). The A-B builds are identified as the functionally graded structures and

the A-B-A builds are identified as compliant mechanism structures. Using the parameters listed in

Table 10.1, functionally graded and compliant mechanism builds were fabricated. These builds are

shown in Figure 10.5. The build coupons remained attached to the substrate after deposition.

131

(a)

(b)

Figure 10.4 Deposition of compliant mechanism structures using the high power-high

deposition system. (a) During deposition, powder is directed into the melt pool formed by

the laser. (b) After deposition, the build is still hot at the top of the build as heat is conducted

through the previously deposited material.

132

(b)

(a) (c)

Figure 10.5 Functionally graded and compliant mechanism build coupons fabricated using

elemental Ti and elementally blended NiTi powder using the high power-high deposition

system.

The builds were sectioned from the Ti substrate. During the sectioning, a few build coupons

delaminated from the substrate. The only build coupons which delaminated from the substrate were

fabricated with the Ti-rich powder composition (Ni53Ti42 wt.%) deposited first. The pre-heating

was insufficient to mitigate residual stresses and prevent delamination of these larger build coupons

from the substrate. Therefore, if larger structures are fabricated from NiTi powder compositions

additional measures may need to be taken to prevent delamination. Additionally, the effect of

residual stress on microstructure, and how that can be mitigated through heat treatment, must be

characterized.

The proposed future work would characterize the microstructure of these fabricated

functionally graded depositions. It is anticipated that this characterization would enable meaningful

prediction of the shape memory behavior by leveraging the microstructure-MT-shape memory

behavior relationships developed in this dissertation work. However, it would be valuable to verify

this claim with thermo-mechanical experimentation on the functionally graded samples. It is

possible that functionally graded materials exhibit unique behaviors which could not be anticipated

by the relationships developed in this dissertation work.

133

Appendix A

LDED NiTi summary table

Summary table for works which characterize the microstructure or shape memory behavior of additive manufactured NiTi shape

memory alloys using powder fed AM techniques. Powder feedstock compositions were either blended elemental powders (denoted

as Ni:Ti (at.%)) or alloyed powder (denoted as Ni##Ti## (at.%)). Energy density is calculated as

𝒍𝒂𝒔𝒆𝒓 𝒑𝒐𝒘𝒆𝒓 (𝒉𝒂𝒕𝒄𝒉 𝒔𝒑𝒂𝒄𝒊𝒏𝒈 ∗ 𝒔𝒄𝒂𝒏 𝒔𝒑𝒆𝒆𝒅 ∗ 𝒍𝒂𝒚𝒆𝒓 𝒕𝒉𝒊𝒄𝒌𝒏𝒆𝒔𝒔)⁄ . The post-processing denotes whether the characterization was

completed on the as built (AB) or after a post-processing heat treatment (SL for solution treated; HT for (precipitation) aging heat

treatment). Microstructure and shape memory behavior was characterized. Shape memory behavior characterization included

determining the transformation temperatures (TT), or characterizing the superelastic (SE) or shape memory effect (SME) response.

Powder

composition (at.%)

[Ref.]

Energy

Density

(J/mm3)

Size of

Deposits

(mm3) Post-processing

Microstructure

Shape Memory

Behavior

Ph

ase

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TT

SE

SM

E

40:60 [47] - - AB X X X

40:60 [142] 15000 - AB X X X

44.9:55.1 [47] - - - X X X

44.9:55.1 [142] 15000 - HT X X X

Ni44.9Ti55.1 [143] - - AB X X

47.9:52.1 [48] 71 7280 AB X

47.9:52.1 [79] 71 7280 AB X X X X X X

49.9:50.1 [142] - 15000 - X X X

49.9:50.1 [142] - - - X X X

Ni50Ti50 [45] 22-39 577 AB X X X X X

Ni50Ti50 [64] - 1693 AB X X X

Ni50Ti50 [144] 6-20 - AB X X X

134

Powder

composition (at.%)

[Ref.]

Energy

Density

(J/mm3)

Size of

Deposits


Ph

ase

Co

mp

./

Imp

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Gra

in S

ize

Gra

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TT

SE

SM

E

Ni50Ti50 [145] 77 - AB X X X X X

50:50 to 63:37 [46] 155 499 AB X X

Ni50.1Ti49.9 [55] - 497 AB X X X

Ni50.1Ti49.9 [113] - 4524 AB X X

Ni50.1Ti49.9 [60] - 503 HT

Ni50.4Ti49.6 -

Ni50.9Ti49.1 [44] - 1287 SL X X X

Ni50.7Ti48.6 [146] - 4021 AB X X

HT X

Ni50.8Ti49.2 [46] 155 499 AB X X X

Ni50.8Ti49.2 [147] - 18 AB X X X

Ni52Ti48 [61] - - HT X X X

53:47 [48] 71 7280 AB X

HT X

53:47 [67] 71 7280 AB; HT X X X

53:47 [65] 71 7280

AB X X X X X X X X

HT X X

SL X X X X X X X X

SL+HT X X X X X X

55:45 [46] 155 499 SL X X X

57:43 [148] 160 1995 SL+HT X X X X X

- [149] - - AB X X X

- [150] - 15000 HT X X X

135

Appendix B

PBF NiTi summary table

Summary table for works which characterize the microstructure or shape memory behavior of additive manufactured NiTi shape

memory alloys using powder BED AM techniques. Alloyed powder compositions were used (denoted as Ni##Ti## (at.%)). Energy

density is calculated as 𝒍𝒂𝒔𝒆𝒓 𝒑𝒐𝒘𝒆𝒓 (𝒉𝒂𝒕𝒄𝒉 𝒔𝒑𝒂𝒄𝒊𝒏𝒈 ∗ 𝒔𝒄𝒂𝒏 𝒔𝒑𝒆𝒆𝒅 ∗ 𝒍𝒂𝒚𝒆𝒓 𝒕𝒉𝒊𝒄𝒌𝒏𝒆𝒔𝒔)⁄ . The post-processing denotes whether

the characterization was completed on the as built (AB) or after a post-processing heat treatment (SL for solution treated; HT for

(precipitation) aging heat treatment). Microstructure and shape memory behavior was characterized. Shape memory behavior

characterization included determining the transformation temperatures (TT), or characterizing the superelastic (SE) or shape

memory effect (SME) response.

Powder

composition (at.%)

[Ref.]

Energy

Density

(J/mm3)

Size of

Deposits


Microstructure Shape Memory

Behavior

Ph

ase

Co

mp

./

Imp

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Gra

in S

ize

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TT

SE

SM

E

Ni49.3Ti50.7 [120] - - AB X X

Ni49.7Ti50.3 [54] 83 288 AB X X

Ni49.7Ti50.3 [151] 83 318 AB X

Ni49.7Ti50.3 [152] 45-195 - AB X X

Ni49.7Ti50.3 [17] 16-59 827 AB; SL X X

Ni49.9Ti50.1 [49] - 8400 AB

Ni49.9Ti50.1 [153] - - - X

Ni49.9Ti50.1 [154] 24-139 200 AB

Ni49.9Ti50.1 [155] - - AB X X

Ni49.9Ti50.1 [156] 20-45 - AB X

Ni50Ti50 [157] - - AB

Ni50Ti50 [158] - 1080 AB; HT X

Ni50Ti50 [159] - - AB X X

Ni50Ti50 [160] - - - X X

Ni50Ti50 [161] 85 923 HT X X

136

Powder

composition (at.%)

[Ref.]

Energy

Density

(J/mm3)

Size of

Deposits


Ph

ase

Co

mp

./

Imp

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SE

SM

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Ni50Ti50 [162] 24-174 500 AB X X X

Ni50.09Ti49.91 [163] - 159 AB X X

Ni50.09Ti49.91 [53] 44-185 159 AB X X X X

Ni50.09Ti49.91 [79] 56 128 AB X X X X X X X

Ni50.1Ti49.9 [164] 111-126 -

Ni50.1Ti49.9 [165] 95-126 3318 AB X X X X

Ni50.1Ti49.9 [166] 56 608 AB X X

Ni50.1Ti49.9 [167] 55.5 364 AB X X X X

Ni50.12Ti49.9 [168] 111-126 891 AB X X X

Ni50.16i49.84 [56] - - AB X X X

Ni50.2Ti49.8 [169] 83 827 AB X X X X

Ni50.2Ti49.8 [170] - - AB

Ni50.2Ti49.8 [152] 45-195 - AB; HT X X X

Ni50.2Ti49.8 [17] 16-59 827 AB X X X

Ni50.2Ti49.8 [171] 126 432 AB; SL X X X X X X

Ni50.5Ti49.5 [17] 16-59 827 AB X X

Ni50.6Ti49.4 [172] 100-111 191 SL X X X

Ni50.6Ti49.4 [173] 110 191 AB X X

Ni50.6Ti49.4 [174] 65 1045 AB X X X

Ni50.7Ti49.3 [152] 45-195 - AB; HT X X X

Ni50.7Ti49.3 [17] 16-59 827 AB; SL+HT X X X

Ni50.8Ti49.2 [52] 85 - SL; SL+HT X X

Ni50.8Ti49.2 [170] - - AB

Ni50.8Ti49.2 [175] 55.5 603 AB; SL; SL+HT X X X

Ni50.8Ti49.2 [176] 56 603 AB; SL X X X X X

Ni50.8Ti49.2 [177] 55.5 603 AB; SL+HT X X

Ni50.8Ti49.2 [66] 55.5 432 AB; HT X X X X X

Ni50.8Ti49.2 [178] 56 353 AB X

137

Powder

composition (at.%)

[Ref.]

Energy

Density

(J/mm3)

Size of

Deposits


Ph

ase

Co

mp

./

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sity

TT

SE

SM

E

Ni50.8Ti49.2 [179] 55.5 729 AB X X

Ni50.8Ti49.2 [180] 55.5 512 SL+HT X X

Ni50.8Ti49.2 [181] 37-83 159 AB X X X

Ni50.81Ti49.19 [182] 55-278 126 AB X

Ni50.81Ti49.19 [12] - 512 AB X X X

Ni50.9Ti49.1 [183] 58-100 500 HT X

Ni50.9Ti49.1 [184] - 308 HT X

Ni50.9Ti49.1 [58] 45-125 539 AB X X X X X

Ni51Ti49 [57] 58-100 577 AB X X X X

Ni51.2Ti48.8 [185] - 14074 AB X X X

138

Appendix C

NiTi build coupon inventory

Summary table of the build coupons fabricated using the LARS system at CIMP-3D.

Build ID Powder composition (wt.%) # passes # layers substrate temp. (°C) Notes

B1 Ni53Ti47 1 1 25 Parameter development

B2 Ni53Ti47 1 1 25 Parameter development - Figure 2.6




















B22 Ni53Ti47 1 12 245 DSC

B23 Ni53Ti47 3 6 245 DSC

B24 Ni53Ti47 3 12 245

B25 Ni53Ti47 3 12 245 DSC

B26 Ni53Ti47 3 14 245

139

Build ID Powder composition (wt.%) # passes # layers substrate temp. (°C) Notes

B27 Ni53Ti47 3 14 235 DSC

B28 Ni53Ti47 3 14 250

B29 Ni53Ti47 3 14 270

B30 Ni53Ti47 1 14 385 DSC

B31 Ni53Ti47 1 14 350 DSC

B32 Ni53Ti47 1 14 320

B33 Ni53Ti47 1 14 320

B34 Ni53Ti47 6 14 225 DSC, microscopy

B35 Ni53Ti47 6 14 305 sectioned into tension specimens

B36 Ni53Ti47 6 14 330 nozzle clogged during deposition

B37 Ni53Ti47 6 14 305 microscopy

B38 Ni53Ti47 6 14 355 sectioned into compression specimens

B39 Ni53Ti47 1 6 270 DSC

B40 Ni53Ti47 1 6 325

B41 Ni53Ti47 1 6 350

B42 Ni53Ti47 1 6 365

B43 Ni58Ti42 1 14 250

B44 Ni58Ti42 1 14 350

B45 Ni58Ti42 1 14 360

B46 Ni58Ti42 1 14 370

B47 Ni58Ti42 1 14 370

B48 Ni58Ti42 1 14 355

B49 Ni58Ti42 3 14 225 sectioned into tension and compression specimens

B50 Ni58Ti42 3 14 335 sectioned into tension and compression specimens







B57 Ni58Ti42 6 14 390 microscopy

140

Appendix D

Sample and Specimen Extraction Locations

Summary table of specimens, which were extracted from the AM build coupons. Specimens locations are identified, based on the

origin of each build coupon. Units are mm. The material conditions are: as built; directly aged (DA); heat treated (HT);

solutionized (SL); or solutionized then aged (SL+HT).

Build ID (composition) Specimen ID Loading axis

Specimen location

Material Condition X Y Z

B35 (Ni53Ti47)

B35T1 Y 6.5 28 2.5 as built

B35T2 Y 6.5 28 3.6 as built

B35T3 Y 6.5 28 4.7 as built

B35T4 Y 6.5 28 5.8 as built

B35T5 Y 6.5 28 6.9 as built

B38 (Ni53Ti47)

B38C1 Z 4.5 12.2 5 as built

B38C2 Z 4.5 16.4 5 as built

B38C3 Z 4.5 20.6 5 as built

B38C4 Z 4.5 24.8 5 as built

B38C5 Z 4.5 29 5 as built

B38C6 Z 4.5 33.2 5 as built

B38C7 Z 4.5 37.4 5 as built

B38C8 Z 4.5 41.6 5 as built

B382C1 Z 8.6 12.2 5 as built

B382C2 Z 8.6 16.4 5 as built

B382C3 Z 8.6 20.6 5 as built

B382C4 Z 8.6 24.8 5 as built

B382C5 Z 8.6 29 5 as built

B382C6 Z 8.6 33.2 5 as built

B382C7 Z 8.6 37.4 5 as built

B382C8 Z 8.6 41.6 5 as built

141

Build ID Specimen ID Load Dir. X Y Z Material Condition

B49 (Ni58Ti42)

B49T1 Y 4 24 1.5 SL (1000 °C, 36 h)

B49T2 Y 4 24 2.6 SL (1000 °C, 36 h)

B49T3 Y 4 24 3.7

B49T4 Y 4 24 4.8

B49T5 Y 4 24 5.9

B49T6 Y 4 24 7

B49C1 Y 4 41 2.5

B49C2 Y 4 41 6.5

B50 (Ni58Ti42)

B50T1 Y 2 17 4.5

B50T2 Y 3.1 17 4.5 SL (1000 °C, 36 h)

B50T3 Y 4.2 17 4.5

B50T4 Y 5.3 17 4.5 SL (1000 °C, 36 h)

B50C1 Y 4 42 2.5

B50C2 Y 4 42 6.5

B50C3 Y 4 33 2.5

B50C4 Y 4 33 6.5

B51 (Ni58Ti42)

B51C1 Y 4 6 4

B51C2 Y 4 16 4

B51C3 Y 4 26 4

B51C4 Y 4 34 4 as built

B51C5 Y 4 46 4 as built

B52 (Ni58Ti42)

B52C1 Z 4 3 5

B52C2 Z 4 8.5 5

B52C3 Z 4 14 5

B52C4 Z 4 19.5 5

B52C5 Z 4 25 5

B52C6 Z 4 30.5 5

B52C7 Z 4 36 5

B52C8 Z 4 41.5 5 as built

B52C9 Z 4 47 5 as built

142


B53 (Ni58Ti42)

B53CZ7 Z 4.5 7 5 HT (1000 °C, 36 h)

B53CZ12 Z 4.5 12 5

B53CZ17 Z 4.5 17 5

B53CZ22 Z 4.5 22 5

B53CZ27 Z 4.5 27 5 HT (1000 °C, 36 h)

B53CZ32 Z 4.5 32 5

B53CZ37 Z 4.5 37 5

B53CZ42 Z 4.5 42 5 HT (1000 °C, 36 h)

B53CZ47 Z 4.5 47 5 HT (950 °C, 24 h)

B53CY7CS Y 8.6 7 2.5

B53CY7FS Y 8.6 7 6.5

B53CY16CS Y 8.6 16 2.5

B53CY16FS Y 8.6 16 6.5

B53CY25CS Y 8.6 25 2.5

B53CY25FS Y 8.6 25 6.5

B53CY34CS Y 8.6 34 2.5

B53CY34FS Y 8.6 34 6.5

B53CY43CS Y 8.6 43 2.5

B53CY43FS Y 8.6 43 6.5

B54 (Ni58Ti42)

B54T1 Y 6.5 28 1.5 as built

B54T2 Y 6.5 28 2.6 as built

B54T3 Y 6.5 28 3.7 -

B54T4 Y 6.5 28 4.8 as built

B54B5 Y 6.5 28 5.9 DA (450 °C, 3h)

B54T6 Y 6.5 28 7 DA (550 °C, 3h)

B54T7 Y 6.5 28 8.1 as built

143


B55 (Ni58Ti42)

B55T1 Y 6.5 28 1.2 DA (450 °C, 3h)

B55T2 Y 6.5 28 2.3 -

B55T3 Y 6.5 28 3.4 as built

B55T4 Y 6.5 28 4.5 DA (550 °C, 3h)

B55T5 Y 6.5 28 5.6 DA (450 °C, 3h)

B55T6 Y 6.5 28 6.7 as built

B56 (Ni58Ti42)

B56C1 Z 4.5 14.2 5 as built

B56C2 Z 4.5 18.4 5 SL (950 °C, 24h)

B56C3 Z 4.5 22.6 5 as built

B56C4 Z 4.5 26.8 5 DA (550 °C, 3h)

B56C5 Z 4.5 31 5 DA (450 °C, 3h)

B56C6 Z 4.5 35.2 5 DA (450 °C, 3h)

B56C7 Z 4.5 39.4 5 DA (450 °C, 3h)

B56C8 Z 4.5 43.6 5 DA (450 °C, 3h)

B562C1 Z 8.6 14.2 5 DA (550 °C, 3h)

B562C2 Z 8.6 18.4 5 DA (550 °C, 3h)

B562C3 Z 8.6 22.6 5 HT (950 °C, 24h)

B562C4 Z 8.6 26.8 5 -

B562C5 Z 8.6 31 5 SL (950 °C, 24h) + HT (550 °C, 3h)

B562C6 Z 8.6 35.2 5 HT (950 °C, 10h)

B562C7 Z 8.6 39.4 5 -

B562C8 Z 8.6 43.6 5 SL (950 °C, 24h)

144

Appendix E

Nontechnical Abstract

Additive manufacturing (AM) offers the capability to design material microstructures on

multiple length scales. This allows for the tailoring of properties and shape memory response of

NiTi shape memory alloys (SMAs). The layer-by-layer deposition process of AM produces

localized microstructures which are distinct from those found in conventionally processed alloys.

This work characterizes the grain and precipitate microstructures resulting from the interfaces that

build up during the layer-by-layer deposition process. These microstructures are correlated to

deformation mechanisms using multi-scale thermo-mechanical experimentation. By developing

SMAs for the laser-based directed energy deposition (LDED) AM technique, structures with

engineered microstructure, composition, and geometry can be fabricated to achieve on-demand

shape memory behavior. The current work fabricated build coupons from two feedstock

compositions using LDED. The AM microstructure and shape memory responses were spatially

resolved by extracting specimens from the AM build coupons and employing a full-field, multi-

scale deformation analysis technique. Results show that the AM microstructure is inherently

spatially varying. Finer grain and precipitate morphologies are distinctive of the interfacial regions

between layers and passes. The additive manufactured alloys exhibit shape memory effect and

superelastic shape memory behaviors in the as built condition. Post-deposition heat treatments are

shown to refine the precipitate morphology, providing a means to enhance shape memory recovery.

145

Appendix F

Publications and Presentations

REFERRED JOURNAL PUBLICATIONS

Bimber, Beth A., Reginald F. Hamilton, and Todd A. Palmer. “Ni-concentration dependence of

directed energy deposited NiTi alloy microstructures.” Shape Memory & Superelasticity. In

Publication.

Hamilton, Reginald F., Beth A. Bimber, Todd A. Palmer. 2018. “Correlating microstructure and

superelasticity of directed energy deposition additive manufactured Ni-rich NiTi alloys.”

Journal of Alloys and Compounds. 739: 712-722.

Bimber, Beth A., Reginald F. Hamilton, Jayme Keist, and Todd A. Palmer. 2016. “Anisotropic

Microstructure and Superelasticity of Additive Manufactured NiTi Alloy Bulk Builds Using

Laser Directed Energy Deposition.” Materials Science and Engineering A 674: 125–34.

Hamilton, Reginald F., Beth A. Bimber, Mohsen Taheri Andani, Mohammad H. Elahinia. 2017.

“Multi-scale shape memory effect recovery in NiTi alloys additive manufactured by selective

laser melting and laser directed energy deposition.” Journal of Materials Processing

Technology 250 55–64.

Elahinia, Mohammad H., Narges Shayesteh Moghaddam, Mohsen Taheri Andani, Amir

Amerinatanzi, Beth A. Bimber, and Reginald F. Hamilton. 2016. “Fabrication of NiTi through

Additive Manufacturing: A Review.” Progress in Materials Science 83: 630–63.

Hamilton, Reginald F., Todd A. Palmer, and Beth A. Bimber. 2015. “Spatial Characterization of

the Thermal-Induced Phase Transformation throughout As-Deposited Additive Manufactured

NiTi Bulk Builds.” Scripta Materialia 101: 56–59.

PRESENTATIONS

Bimber, Beth A., Reginald, F. Hamilton, and Todd A. Palmer. “Additive Manufacturing Shape

Memory Alloys: Heat Treating to Tune Microstructure” Poster Presentation at Materials Day

2017, University Park, PA October 17, 2017.

Bimber, Beth A., Reginald F. Hamilton. “Laser Directed Energy Deposition Additive

Manufactured NiTi SMAs: Heat Treated Material Microstructures and Superelasticity”

Presentation at Materials Science and Technology Conference 2017, Pittsburgh, PA, October 8-

12, 2017.

Bimber, Beth A., Reginald F. Hamilton, Todd A. Palmer. “Additive Manufacturing for Tuning

Shape Memory Behavior” Poster Presentation at Center for Innovative Materials Processing

through Direct Digital Deposition AM Day and Open House, University Park, PA, June 15,

2017.

Bimber, Beth A., Reginald F. Hamilton, Todd A. Palmer. “Additive Manufacturing for Tuning

Shape Memory Behavior” Poster Presentation at Penn State College of Engineering

IndustryXchange, University Park, PA, May 24, 2017.

146

Bimber, Beth A., Reginald F. Hamilton, Jayme Keist, Todd A. Palmer. “Anisotropic

Microstructure and Superelasticity of Additive Manufactured NiTi Alloy Bulk Builds Using

Laser Directed Energy Deposition” Poster Presentation at Research Penn State 2016, University

Park, PA October 5-6, 2016.

Bimber, Beth A., Reginald F. Hamilton. “Characterization of Bulk NiTi Components Fabricated

using Laser Based Directed Energy Deposition” Presentation at the Materials Science and

Technology Conference 2015, Columbus, OH, October 4-8, 2015.

Bimber, Beth A., Reginald F. Hamilton. “Powder Bed Fusion and Directed Energy Deposition

Additive Manufacturing Techniques for Fabrication of NiTi Shape Memory Alloys”

Presentation at the Materials Science and Technology Conference 2015, Columbus, OH,

October 4-8, 2015.

Bimber, Beth A., Reginald F. Hamilton. “Shape Memory Response in NiTi Fabricated Using

Laser-based Directed Energy Deposition” Presentation at the Materials Science and Technology

Conference 2014, Pittsburgh, PA, October 27-31, 2014.

Bimber, Beth A., Reginald F. Hamilton. “Development of Novel Laser-based Digital Direct

Manufactured Shape Memory Alloys” Presentation at the Materials Science and Technology

Conference 2013, Montreal, Canada, October 27-31, 2013.

147

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169

VITA

Beth (Bimber) Last graduated in 2011 with a B.S. degree in Engineering Science and

Mechanics from the Pennsylvania State University. She obtained a M.S. degree in Engineering

Science in 2013, also from the Pennsylvania State University. These degrees were attained under

the guidance of Dr. Barbara Shaw, where she developed, fabricated, and optimized a three-

electrode device for measuring in vivo corrosion rates of novel Mg-Ti thin films. She later pursued

a Ph.D. at the Pennsylvania State University, under the guidance of Dr. Reginald F. Hamilton.

During the doctoral studies, Beth investigated the microstructure and shape memory behavior of

additive manufactured NiTi shape memory alloys.

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