Document [original]

Characterization of the fatigue behavior of additively

manufactured β-metastable titanium alloy

Ti-5Al-5V-5Mo-3Cr on different length scales

vorgelegt von

Dipl.-Ing.

Erika Gabriele Alves Alcântara

ORCID: 0000-0002-7736-6794

von der Fakultät III – Prozesswissenschaften

der Technischen Universität Berlin

zur Erlangung des akademischen Grades

Doktor der Ingenieurwissenschaften

- Dr.-Ing. -

genehmigte Dissertation

Promotionsausschuss:

Vorsitzende: Prof. Dr. Isabella Gallino

Gutachterin: Prof. Dr.-Ing. Claudia Fleck

Gutachter: Prof. Dr. Michael de Wild

Tag der wissenschaftlichen Aussprache: 5. Februar 2024

Berlin 2024

To my ancestors, for everything they went

through so that life could reach me and I could

be who I am.

To the bravery of those who dare to make a

difference in a foreign country, weather the

storms, and come out stronger.

To myself, for pushing through, and finishing

what I started.

"A Ph.D. is not only a degree or about learning a subject; it's about mastering the art

of overcoming challenges, with dedication, perseverance, and the willingness to

embrace uncertainty"

- Erika Alcântara -

III

Alves Alcântara, Erika Gabriele. Characterization of the fatigue behavior of

additively manufactured β-metastable titanium alloy Ti-5Al-5V-5Mo-3Cr on

different length scales.

Technische Universität Berlin, Institute of Materials Science and Technology, Chair of

Materials Science and Engineering

Strasse des 17. Juni 135

D-10623 Berlin

Abstract

The β-metastable titanium alloy Ti-5Al-5V-5Mo-3Cr (Ti-5553) has recently found

growing interest as medical implant material due to its advantageous mechanical

properties when compared to the up-to-date standard alloys. Besides biocompatibility,

implant materials need to exhibit sufficient fatigue resistance. This thesis investigates

the fatigue and cyclic deformation behavior on different length scales of Ti-5553, made

by laser powder bed fusion of metals (LPBF-M), with distinct microstructures resulting

from different heat treatments. The nano-scale investigation reveals how the α-

precipitates orientation and distribution influence the deformation behavior of Ti-5553

in the (α+β)-solution annealed state (ST) during cyclic nanoindentation. In addition,

quasi-static and cyclic nanoindentation tests were conducted to investigate the

mechanical properties of LPBF-M Ti-5553 in the “as-built”, (α+β)-solution annealed and

aged (STA) and β-annealed, slowly cooled, and aged (BASCA) conditions.

Microstructural variations notably influence the nanohardness, reduced elastic

modulus, plasticity, and cyclic deformation behavior of Ti-5553. The macro-scale

perspective summarizes the cyclic deformation behavior of shot-peened Ti-5553

specimens in Hanks’ balanced salt solution (HBSS). Changes in the free corrosion

potential and the corrosion current are highly sensitive indicators for fatigue-induced

damage on a rough surface, which was correlated to the microscopic examination,

fracture surface features, and fatigue crack development. The gained deeper

understanding of the mechanical performance of this promising β-metastable alloy

across different microstructures and different length scales helps establish a

foundation for predicting fatigue characteristics and optimizing LPBF-M Ti-5553 alloy

performance, especially in the context of implant materials.

Keywords: additively manufactured Ti-5Al-5Mo-5V-3Cr; ST heat treatment; STA heat

treatment; BASCA heat treatment; cyclic nanoindentation; corrosion-fatigue; cyclic

deformation behavior; high cycle fatigue; fatigue-induced surface damage; implant

materials

Alves Alcântara, Erika Gabriele. Charakterisierung des Ermüdungsverhaltens

additiv gefertigter β-metastabile Titanlegierung Ti-5Al-5V-5Mo-3Cr auf

verschiedenen Längenskalen.

Technische Universität Berlin, Institut für Werkstoffwissenschaften und -technologien,

Fachgebiet Werkstofftechnik

Strasse des 17. Juni 135

D-10623 Berlin

Kurzfassung

Die β-metastabile Titanlegierung Ti-5Al-5V-5Mo-3Cr (Ti-5553) hat aufgrund ihrer im

Vergleich zu den aktuellen Standardlegierungen vorteilhaften mechanischen

Eigenschaften in letzter Zeit wachsendes Interesse als Material für medizinische

Implantate gefunden. Neben der Biokompatibilität müssen Implantatmaterialien eine

ausreichende Ermüdungsbeständigkeit aufweisen. Diese Arbeit untersucht das

Ermüdungs- und zyklische Verformungsverhalten auf verschiedenen Längenskalen

der Ti-5553, hergestellt durch Laser-Pulverbettschmelzen von Metallen (engl. LPBF-

M), mit unterschiedlichen Mikrostrukturen, die aus verschiedenen

Wärmebehandlungen resultieren. Die Untersuchung in der Nanoskala zeigt, wie die

Ausrichtung und Verteilung der α-Ausscheidungen das Verformungsverhalten von Ti-

5553 im (α+β)-lösungsgeglühten Zustand (engl. ST) während der zyklischen

Nanoindentation beeinflusst. Darüber hinaus wurden quasistatische und zyklische

Nanoindentationstests durchgeführt, um die mechanischen Eigenschaften von LPBF-

M Ti-5553 in den konturnahen (engl. „as-built“), (α+β)-lösungsgeglühten und

gealterten (engl. STA) und β-geglühten, langsam abgekühlten und gealterten (engl.

BASCA) Zustände zu untersuchen. Mikrostrukturelle Variationen beeinflussen

insbesondere die Nanohärte, den reduced Elastizitätsmodul, die Plastizität und das

zyklische Verformungsverhalten von Ti-5553. Die Makroperspektive fasst das

zyklische Verformungsverhalten von kugelgestrahlten Ti-5553-Proben in Hanks‘

ausgewogener Salzlösung (engl. HBSS) zusammen. Änderungen des freien

Korrosionspotentials und des Korrosionsstroms sind hochempfindliche Indikatoren für

ermüdungsbedingte Schäden an einer rauen Oberfläche, die mit der mikroskopischen

Untersuchung, den Merkmalen der Bruchoberfläche und der Entwicklung von

Ermüdungsrissen korreliert wurden. Das gewonnene tiefere Verständnis der

mechanischen Leistung dieser vielversprechenden β-metastabilen Legierung über

verschiedene Mikrostrukturen und verschiedene Längenskalen trägt dazu bei, eine

Grundlage für die Vorhersage von Ermüdungseigenschaften und die Optimierung der

Leistung der LPBF-M Ti-5553-Legierung zu schaffen, insbesondere im

Zusammenhang mit Implantatmaterialien.

Schlüsselwörter: additiv gefertigtes Ti-5Al-5Mo-5V-3Cr; ST-Wärmebehandlung;

STA-Wärmebehandlung; BASCA-Wärmebehandlung; zyklische Nanoindentation;

Korrosionsermüdung; zyklisches Verformungsverhalten; Ermüdung bei hohem Zyklus;

ermüdungsbedingte Oberflächenschäden; Implantatmaterialien

Eidesstattliche Erklärung

Hiermit versichere ich an Eides statt, dass ich die Dissertation selbständig verfasst

habe. Alle benutzten Hilfsmittel und Quellen sind aufgeführt. Weiter erkläre ich, dass

ich nicht schon anderweitig einmal die Promotionsabsicht angemeldet oder ein

Promotions-eröffnungsverfahren beantragt habe. Veröffentlichungen von

irgendwelchen Teilen der vorliegenden Dissertation sind von mir vorgenommen

worden.

Berlin, den 25.02.2024

Erika Gabriele Alves Alcântara

Acknowledgments

This dissertation is based on the work conducted at the Technische Universität

Berlin (TU Berlin) under the collaboration of the Chair of Materials Science and

Engineering (Faculty III of Process Sciences, Institute of Material Sciences and

Technology) and Chair of Machine Tools and Manufacturing Technology (Faculty V of

Mechanical Engineering and Transport Systems, Institute of Machine Tools and

Factory Management). I am grateful to Prof. Dr.-Ing. Claudia Fleck for the supervision

of the thesis, for her guidance, encouragement, and scientific advice, which

enlightened my way till the end, and especially for the great scientific independence

she gave to me. I would like also to thank Prof. Dr. h. c. Dr.-Ing. Eckart Uhlmann for

providing me with the materials that were analyzed in the context of this work. Without

it, I would not be able to conduct this research.

Special thanks go to all of my colleagues, especially Dr. rer. nat. Reinhard

(René) Meinke, for supporting me in facing the technical challenges encountered

during the development of my dissertation. I would like to thank Martina Schaube for

the kind metallographic preparation of the specimens; Ralf Engelmayer for technical

support; Dr. Anke Märten and M.Sc. Merle Schmahl for the support with the

nanoindentation technique; Kai Winkler and Dr.-Ing. Daniela Hübler for the support

with Python programming used to evaluate most of the nanoindentation data; and

Central Electron Microscopy Unit (ZELMI) from TU Berlin for the support with the high-

resolution scanning electron microscopic (HR-SEM) images, electron backscatter

diffraction measurements (EBSD), focused ion beam (FIB) preparation, and

transmission electron microscopy (TEM) investigations.

My sincere thanks go to all the students who developed their bachelor (BA),

master (MA), and diploma (DI) final works under my co-supervision: Ana Luiza de

Paiva Melo (DI), Frederik Johannes Schwartz (BA), Ilaria Nigro (MA), Juliana Sofia

Fonseca Camargos (DI), Larrisa Mara Batista Duarte (DI), Lukas Joachim Wagner (BA

and MA), Marina Fernandes Jamar (DI), Moritz Pelk (BA), Nawar Yusfi (MA), and Suad

Beganovic (MA). You contributed significantly to my further professional and personal

development.

I also would like to express my gratitude to the evaluation committee — special

thanks to Prof. Dr. Michael de Wild for his willingness to evaluate this doctoral thesis.

VII

Huge and special gratitude goes to my beloved husband Markus, my son Liam,

and my fofulino Basti for their patience, for encouraging me to be the best person I can

be, and for always giving me the strength and unconditional love often needed.

Last but not least, the greatest thanks and appreciation go to my parents

Tarcísio and Elzi, my sisters Rafaella and Cindy, and our lindeza Kurt for turning me

into the person I am today, for always believing in me and supporting me in my life

choices that brought me to achieve this goal of completing my doctoral study.

VIII

List of symbols and abbreviations

Symbols

projected indent area

Ap-u

pile-up area

Percentage of total elongation at fracture

Da,p

plastic indentation depth amplitude

depth of the contact circle

displacement associated with the elastic recovery during unloading

Dmax

maximum depth

Dmin

minimum depth

depth of the residual impression

elastic modulus

frequency

hp-u,max

maximum pile-up height

Icorr

free corrosion current

KIC

fracture toughness

martensite start temperature

number of cycles

Nmax

maximum cycle number

force amplitude

mean force

Pmax

maximum force

Pmín

minimum force

load ratio

linear correlation coefficients

arithmetic mean roughness

tensile strength

average maximum surface height

stress magnitude

Tβ

β-transus temperature

Ucorr

free corrosion potential

Vp-u

pile-up volume

alpha titanium phase

α'

alpha prime titanium phase

α’’

alpha double prime titanium phase

αp

primary α-phase

beta titanium phase

βc

critical minimum level of β-stabilizing elements

βs

similar intersection point of the β-transus temperature

ΔDmin

change of plastic deformation between cycles

strain

εa,p

plastic strain amplitude

2θ

angle between transmitted beam and reflected beam by X-ray diffraction

Kα

specific wavelength of the X-ray spectrum

σ0.2

yield stress

σ107

high cycle fatigue strength

σa

stress amplitude

omega titanium phase

Abbreviations

Additive manufacturing; additive manufactured

LPBF-M

Laser powder bed fusion of metals

ABF

Annular bright field

ASTM

American society for testing and materials

Bachelor final work

BASCA

Slow cooling and aging

bcc

Body-centered cubic structure

Bright field

BSE

Backscattered electron

CAT

Constant-amplitude test

Counter electrode

CP-Ti

Commercially pure titanium alloy

DED

Directed energy deposition

Diploma final work

DIN

Deutsches Institut für Normung

EBSD

Electron backscatter diffraction

EDX

Energy-dispersive X-ray spectroscopy

Europäische Norm

FDA

U.S. food and drug administration

FIB

Focused ion beam

HBSS

Hanks' balanced salt solution

HCF

High cycle fatigue

hcp

Hexagonal closed-packed structure

HR-SEM

High-resolution scanning electron microscopic

Hardness according to Vickers

Ingot metallurgy

ISO

International organization for standardization

LIT

Load-controlled axial stepwise load increase test

Laser melted

Master final work

Nanoindentation

Recrystallization anneal

Reference electrode

SFE

Stacking fault energy

SLM

Selective laser melting

Solution treatment

STA

Solution treatment and aging

STOA

Solution treated plus over aged in the β-field

TEM

Transmission electron microscopy

Titanium

Ti-1023

Ti-10V-2Fe-3Al alloy

Ti-5553

Ti-5Al-5V-5Mo-3Cr

Ti-62222

Ti-6Al-2Sn-2Zr-2Mo-2Cr alloy

Ti-64

Ti-6Al-4V

TU Berlin

Technische Universität Berlin

UTS

Ultimate tensile strength

VED

Volumetric energy density

VT22

Ti-5Al-5V-5Mo-1Cr-1Fe alloy

Working electrode

XII

XRD

X-ray diffraction

β-C

Ti-3Al-8V-6Cr-4Zr-4Mo alloy

β-CEZ

Ti-5Al-2Sn-2Cr-4Mo-4Zr-1Fe alloy

CONTENT

ABSTRACT ................................................................................................................III

ZUSAMMENFASSUNG .............................................................................................IV

EIDESSTATTLICHE ERKLÄRUNG ............................................................................V

ACKNOWLEDGEMENTS ..........................................................................................VI

LIST OF SYMBOLS AND ABBREVIATIONS ……...……………………….………....VIII

1 INTRODUCTION.................................................................................................. 3

1.1 Background and motivation of the study ..................................................... 3

1.2 Importance of studying cyclic deformation behavior on different length

scales......................................................................................................................... 5

1.3 Overview and outline of the thesis ................................................................ 5

2 STATE OF THE ART ........................................................................................... 7

2.1 Titanium and its alloys ................................................................................... 7

2.1.1 The β-metastable Ti-5553 alloy ................................................................... 13

2.2 Heat treatments for β-metastable titanium alloys ...................................... 16

2.3 Titanium and its alloys as implant materials .............................................. 19

2.4 Additive manufacturing of titanium alloys ................................................. 20

2.4.1 Laser Powder Bed Fusion of Metals (LPBF-M) ........................................... 23

2.4.2 Processability of Ti-5553 via LPBF-M .......................................................... 26

2.5 Fatigue ........................................................................................................... 27

2.5.1 Corrosion-fatigue ......................................................................................... 28

2.5.2 Nano-fatigue ................................................................................................ 31

3 MATERIAL ........................................................................................................ 35

4 METHODS ......................................................................................................... 39

4.1 Pre-loading microstructural investigation .................................................. 39

4.2 Nano-scale experiments............................................................................... 40

4.2.1 Quasi-static nanoindentation ....................................................................... 40

4.2.2 Cyclic nanoindentation ................................................................................ 41

4.3 Macro-scale experiments ............................................................................. 44

5 RESULTS .......................................................................................................... 47

5.1 Microstructure ............................................................................................... 47

5.2 Nano-scale investigation: ST condition ...................................................... 56

5.2.1 Cyclic deformation and creep behavior ....................................................... 56

5.2.2 Cyclic nanoindent morphology..................................................................... 60

5.2.3 TEM investigation ........................................................................................ 62

5.3 Nano-scale investigation: comparison between “as-built”, STA and

BASCA conditions .................................................................................................. 64

5.3.1 Quasi-static nanomechanical properties ...................................................... 64

5.3.2 Cyclic deformation behavior ........................................................................ 66

5.4 Macro-scale investigation: STA condition .................................................. 72

5.4.1 Estimation of the endurance limit in load increase tests .............................. 72

5.4.2 Cyclic deformation behavior in simulated physiological media .................... 73

5.4.3 Characterization of fatigue-induced surface damage in CAT ...................... 75

6 DISCUSSION ..................................................................................................... 79

6.1 Influence of α-precipitate orientation and distribution on the deformation

behavior of Ti-5553 in the ST condition ................................................................ 79

6.1.1 Microstructure .............................................................................................. 79

6.1.2 Cyclic deformation and creep behavior ....................................................... 80

6.1.3 Fatigue-induced indent characteristics ........................................................ 82

6.1.4 Summarizing model mechanism.................................................................. 83

6.2 Effect of STA and BASCA heat treatments on the microstructure and

nanomechanical properties of Ti-5553 ................................................................. 84

6.2.1 Microstructure .............................................................................................. 85

6.2.2 Quasi-static mechanical properties .............................................................. 86

6.2.3 Nanofatigue behavior .................................................................................. 88

6.3 Cyclic deformation behavior of Ti-5553 in quasi-physiological medium . 92

6.3.1 Fatigue-induced surface damage ................................................................ 94

7 SUMMARY AND ADDRESSED OPEN QUESTIONS ....................................... 97

8 CONCLUSION AND OUTLOOK ..................................................................... 100

9 REFERENCES ................................................................................................ 102

10 SUPPORTING INFORMATION ....................................................................... 121

1 INTRODUCTION

1.1 Background and motivation of the study

In view of the rapidly increasing of aged population ratio of representative

countries after the year 2030 together with the still increasing life expectancy of the

population [1], a great number of revision surgeries will require the replacement of

failed bone tissue. Wear/corrosion, fibrous encapsulation, inflammation, low fracture

toughness/low fatigue strength, and mismatch in modulus of elasticity (E) between

bone and implant (see Fig. 1) are possible causes for implant failure that can lead to

interventions [2].

Fig. 1: Possible causes for implant failure that can lead to interventions and their

effects on the body (Reprinted from [2], with permission of the copyright owner).

Fracture, wear, and corrosion are the most often failure modes in implant

materials, such as long-term orthopedic and cardiovascular implants, along with short-

term implants including bone screws and plates [3]. Due to cyclic loading in the body,

fatigue-induced implant fractures are common [4,5]. Hence, fatigue failure has been

investigated for several years as an important cause of implant failure [6–9]. In recent

years, the raised demand from patients for customized implants with enhanced long-

term stability, together with the transition from conventional to rapidly and directly from

computer-aided fabrication, has also driven research interests in fatigue toward

additive manufactured implant materials [10–16]. Additive manufacturing (AM) consists

of fabricating a part layer-by-layer by utilizing a combination of energy delivery and

material deposition. One process, namely laser powder bed fusion of metals (LPBF-

M), currently offers new design possibilities for the development of customized medical

implants, for example for complex bone reconstructions [17–19]. The conventional

manufacturing process of implants can take days or weeks, being very costly and still

not suitable for every patient, whilst AM can provide more affordable customized

implants that can be fabricated in only several hours.

Titanium alloys (Ti-alloys) are in common use for implant applications because of

their excellent combination of mechanical strength, corrosion resistance, and

biocompatibility [2,20]. Among the different Ti-alloys, (α+β)-alloy Ti-6Al-4V (Ti-64) is

one of the most used Ti-based materials for orthopedics implant fabrication because

of its high strength and fatigue properties at room temperature, together with its higher

resistance to corrosion compared with other biomedical alloys [21]. However, there is

a mismatch in stiffness between the Ti-64 implants and the surrounding bone. This

leads to the so-called “stress shielding effect” [22], which causes bone resorption, and,

eventually, loosening of the implant. As a consequence, the implant needs to be

exchanged, either due to pain or, even worse, following fatigue failure of the implant.

These complications can be avoided by implants with a lower structural stiffness which

may be achieved by porous bulk materials and surfaces allowing the ingrowth of bone,

and/or by implant materials with a lower E, such as β-Ti-alloys [1,23–25]. Recently the

β-metastable Ti-5Al-5V-5Mo-3Cr (Ti-5553) has found growing interest as biomaterial,

due to its advantageous mechanical properties. As compared to Ti-64, Ti-5553 exhibits

a better combination of ductility, toughness, and E [26].

Understanding the fatigue and cyclic deformation behavior of promising implant

materials is crucial for improving the durability and reliability of implants, ultimately

leading to better patient outcomes. Therefore, the purpose of this work is to

characterize the fatigue behavior on different length scales of AM β-metastable Ti-5553

alloy with different microstructures. On the nano-scale, quasi-static and cyclic

nanoindentation tests combined with high-resolution electron microscopy were

performed to investigate the influence of the microstructures on the mechanical

properties of Ti-5553. The cyclic deformation behavior on the macro-scale was

evaluated in simulated physiological media, including the examination of fatigue-

induced surface damage by means of corrosion potential and current measurements

correlated to microstructural investigations.

1.2 Importance of studying cyclic deformation behavior on different length

scales

Characterizing the cyclic deformation behavior of materials at different length

scales is crucial to gain a comprehensive understanding of the fatigue response of

materials and to predict their performance under cyclic loading conditions in silico [27].

Macro-fatigue tests provide an average response over the microstructural constituents,

aiming to understand the fatigue life and the mechanisms leading to failure [28]. On

the other hand, micro-scale fatigue testing can delve into the local interactions of grains

or phases with deformation mechanisms like dislocations and twinning, allowing us to

extract the influence of local structural inhomogeneities on the cyclic deformation

response of materials [29]. Furthermore, fatigue testing at the micro-scale offers an

alternative mean to determine the fatigue properties and to study the cyclic deformation

mechanisms, especially for materials where conventional fatigue tests are impractical.

This present work investigates how AM β-metastable Ti-5553 alloy responds to

cyclic loading over time across different length scales approaches, encompassing

localized microstructural aspects (e.g. defects, phase distribution) and macrostructural

features (e.g. surface roughness, interactions with the environment). This knowledge

may contribute to the development of safer and more reliable implant materials capable

of withstanding the stresses and strains experienced within the human body, ultimately

reducing the risk of patient injury.

1.3 Overview and outline of the thesis

This thesis characterizes the cyclic deformation behavior on different length

scales of AM β-metastable Ti-5553 alloy with distinct microstructures resulting from

different heat treatments. The nano-scale investigation reveals how the α-precipitates

orientation and distribution influence the deformation behavior of Ti-5553 in the (α+β)-

solution annealed state (ST) during cyclic nanoindentation. Additionally, the impact of

(α+β)-solution annealing and aging (STA) and β-annealing, slow cooling, and aging

(BASCA) heat treatments on the microstructure and nanomechanical properties of Ti-

5553 was comparatively investigated. The macro-scale perspective summarizes the

cyclic deformation behavior of Ti-5553 in simulated physiological media, correlating

microstructural investigations with corrosion potential and current measurements to

investigate surface damage caused by fatigue.

The next chapter focuses on the state of the art regarding Ti and its alloys,

including their classifications, a summary of the commonly performed heat treatments

for β-metastableTi-alloys with their resulting microstructure, and a brief overview on Ti

and its alloys as implant materials. This chapter also provides background information

and context for the fatigue testing conducted at both nano and macro scales, as well

as a concise explanation of the fundamental functionality of the AM variante known as

LPBF-M that was used to produce the specimens analyzed in the context of this work.

The chemical composition of the powder, and the LPBF-M parameter settings

used to manufacture the specimens are described in Chapter 3.

Chapter 4 provides detailed information about the methods used to characterize

the studied material under different conditions and at different length scales, which led

to the results described in Chapter 5.

Chapter 6 discusses the influence of α-precipitate orientation and distribution on

the deformation behavior of AM-Ti-5553 in the ST condition, along with the effects of

STA and BASCA heat treatments on the microstructure and nanomechanical

properties of this alloy. Additionally, the cyclic deformation behavior of Ti-5553 in a

quasi-physiological medium is discussed.

Chapter 7 offers a summary of the results and the open questions addressed in

this thesis. The conclusions related to the characterization of the fatigue behavior of

AM Ti-5553 on different length scales, as well as suggestions for future research

directions are provided in Chapter 8.

2 STATE OF THE ART

2.1 Titanium and its alloys

Ti, with its atomic number 22 in the periodic table of elements, is a transition

element in the category of light metals with a density of 4.51 g/cm³ [30]. As can be

seen in Fig. 2, this metal exhibits a very good strength-to-density ratio and, depending

on its alloying, it is significantly higher than that of aluminum or magnesium, and about

twice as high as that of low alloy steels. Another good characteristic of Ti is its good

corrosion resistance and biocompatibility in oxidizing environments, which is based on

the high-speed spontaneous formation of a dense and very stable inert oxide layer on

the surface even at low oxygen partial pressures [20].

Fig. 2: Ashby diagram in terms of density (ρ) and strength (σf). The guidelines of

constant σf/ρ, σf2/3/ρ, and σf1/2/ρ are used in minimum weight, yield-limited, design.

(Reprinted from [31], with permission of the copyright owner).

With a melting point of around 1678°C, Ti undergoes an allotropic transformation

in the solid state at 882°C, also known as α→β transition temperature or β-transus

temperature (Tβ), changing from the α-phase (hexagonal closed-packed structure, hcp)

to a β one (body-centered cubic structure, bcc) [21]. The atomic models of hcp and bcc

structures can be seen in Fig. 3. The Tβ can be modified by adding α-stabilizer

elements (Al, C, O, N, B) or β-stabilizer elements (Fe, Mo, V, Cr, Ni, Cu, W, Mn, Co,

Nb, Ta, Si, H), while the neutral elements (Sn, Zr) are not able to modify it, although

they can act as α or β-stabilizer in the presence of other α or β-stabilizers [32,33]. To

qualitatively represent the effects of varying amounts of α and β stabilizers and the

neutral elements in terms of stabilizing the α and β phases in Ti-alloys, it is usual to

convert them to equivalent Al and Mo concentrations, respectively, as follows [33,34]:

[Al]eq = [Al] + [Zr]/6 + [Sn]/3 + 10[O+C+2N], and

[Mo]eq = [Mo] + [Ta]/5 + [Nb]/3.6 + [W]/2.5 + [V]/1.5 + 1.25[Cr] + 1.25[Ni] +

1.7[Mn] + 1.7[Co] + 2.5[Fe]

where [X] is the concentration of element “X” in wt.% in the alloy.

Fig. 3: Atomic models of the hcp (a) and the bcc (b) structures of Ti. (Reprinted from

[32], with permission of the copyright owner).

Thus, depending on the expressions of allotropy, Ti-alloys can consist fully of α

or β or a mixture of α and β phases in equilibrium at room temperature, thereby being

classified as commercially pure α- and near-α-alloys, β-alloys and (α+β)-alloys. When,

however, the β-phase can still be retained at room temperature in metastable

equilibrium by rapid quenching of alloys in which the martensite start temperature (Ms)

is below room temperature, they are classified as metastable β-alloys [34]. The general

classification of the alloying elements can be seen in Fig. 4 and two schematic phase

diagrams depiction commonly used to describe the individual and combined effects of

alloying elements are shown in Fig. 5.

Fig. 4: Overview of the phase diagrams of the neutral, as well as of α and β stabilizing

alloy elements (Reprinted from [32], with permission of the copyright owner).

Fig. 5: a) Schematic pseudo-binary phase diagram for alloys with β-stabilizer and b)

three-dimensional phase diagram depiction of alloys with α and β stabilizers.

(Reprinted from [34], with permission of the copyright owner).

The α- and near-α-Ti-alloys are composed of single α-phase, presenting

relatively low tensile strengths, although their high thermal stability leads to reasonable

creep strengths in high temperatures [33]. The presence of Al, the principal α-stabilizer,

enhances the tensile strength and the E [35]. The formability of α-Ti-alloys is restricted

during hot working at temperatures below Tβ due to their hcp structure, causing them

to exhibit a high rate of strain hardening [33]. The single-phase α-Ti-alloys are

extensively used in applications that are not particularly demanding in terms of strength

but focus more on their attractive corrosion resistance [21,32]. The several grades of

commercially pure titanium (CP-Ti) and the ternary composition Ti–5Al–2.5Sn are the

most commonly used α-Ti-alloys [33]. Near α-Ti-alloys are classified as (fully) α-Ti-

alloys that contain small amounts (about 1.0 to 2.0 wt.%) of β-stabilizers, stabilizing

around 5 to 10% of the β-phase into the structure at room temperature. The presence

of the β-phase enables two-phase strengthening via control over the scale of the two

phases, their morphologies, and distributions [35]. However, these additions of β-

stabilizing elements are too small, making normal strengthening through the

decomposition of retained β also insignificantly small. Therefore, the improvement in

mechanical properties in these alloys arises mainly from the formation of martensitic

phases and from the manipulation of precipitated microstructures [33]. The near α-Ti-

alloys present higher tensile strength at room temperature than do the (fully) α-Ti-

alloys, and they show the greatest creep resistance of all Ti-alloys at temperatures

above approximately 400 °C [33].

The (α+β)-titanium alloys contain more β-phase than the near α-Ti-alloys. To

achieve the biphasic microstructure, α-stabilizers are added to stabilize and strengthen

the α-phase, together with 4–6% of β-stabilizing elements that allow substantial

amounts of β-phase to be retained from the β or (α+β)-phase fields [33]. Compared

with the α-phase, the β-phase induces a lower resistance to plastic deformation and a

significant anisotropy of physical and mechanical properties [21]. Therefore, due to

their biphasic microstructure, (α+β)-Ti-alloys offer a range of better combinations of

strength, toughness, ductility, and high-temperature properties when compared to the

near α-Ti-alloys [32,35]. Heat treatment and thermo-mechanical processing have been

used for a long time to improve the strength properties of (α+β)-Ti-alloys to achieve

high levels of tensile strength at room temperature [36–41]. Currently, the (α+β)-Ti-

alloys have the greatest commercial importance with one composition, namely Ti-64,

also known as IMI 318, comprising more than half the sales of Ti-alloys both in Europe

and the United States of America [33,42]. Originally developed for aircraft structural

applications in the 1950s, the lightweight and yet strong Ti-64 is nowadays established

as the most popular Ti-alloy widely used in highly loaded structures in the aerospace

industry for being extremely suitable for jet engines, gas turbines, and many airframe

components [35,43,44]. One of the principal uses of this alloy was for the forged

components in the Lockheed/Boeing F22 military aircraft [33]. However, due its

excellent properties such as high strength, low density, high fracture toughness,

excellent corrosion resistance, and superior biocompatibility [42], this (α+β)-Ti-alloy,

thereby, has also been used in the marine and chemical industries [45–47], besides

being the most popular choice for application as bioprosthetic and implant materials

[1,2,20]. Particularly for biomedical applications, the Ti-64 ELI is widely used in the

biomedical industry [48,49]. Ti-64 ELI (Grade 23) is very similar to Ti-64 (Grade 5),

except that Ti-64 ELI contains reduced levels of O, N, C, and Fe. ELI is short for “Extra

Low Interstitials”, and these lower interstitials provide improved ductility and better

fracture toughness for the Ti-64 ELI material [50]. Ti–64 ELI has been standardized in

the American Society for Testing and Materials (ASTM) among various Ti-alloys to be

used as surgical implants (ASTM F 136: 2013 [51]), being the most representative Ti-

alloy used as an alternative artificial metallic material for failed hard tissue [52].

The β-Ti-alloys are defined as those containing higher [Mo]eq and lower [Al]eq,

enabling the β-phase to be retained in either a metastable or a stable condition after

cooling to room temperature during a thermal treatment. Alloy compositions referred

to as β-metastable Ti-alloys are those placed between the critical minimum level of β-

stabilizing elements (βc) and the similar intersection point of the Tβ line (βs), as shown

in Fig. 6, thus being able to precipitate a second phase upon aging treatment. In turn,

stable β-Ti-alloys could not be hardened by heat treatments due to more highly alloyed

compositions to the right side of βs, where the β-phase remains stable even after rapid

cooling, preventing the occurrence of martensitic transformation. Thereby, when

compared with the other types α, near-α and (α+β)-Ti-alloys, the β-Ti-alloys exhibit

higher strength at moderate temperatures due to their possible strengthening by the

precipitation of fine α-phase from a super-saturated β-solid-solution. Among the β-Ti-

alloys with their higher β-phase stabilization, TIMETAL 15-3 and TIMETAL 21S are

good examples, both having better formability and being hence mainly used for sheet

and foil applications [34]. Therefore, the β-Ti-alloys can satisfy diverse requirements

of very high strength with adequate toughness and fatigue resistance for use in in

airframe applications [53–55] as well as to meet requirements of low E,

biocompatibility, and fatigue strength for use in biomedical applications [1,2,23,25].

Fig. 6: Pseudo-binary β-isomorphous phase diagram showing locations of metastable

and stable β-Ti-alloys. (Reprinted from [33], with permission of the copyright owner).

In addition to the α and β phase transformations that can be controlled by

moderately rapid diffusion or can occur without change in composition and very quickly,

it is also possible to have non-diffusion transformations that generally lead to non-

equilibrium phases. The decomposition of the β phase in Ti-alloys can occur by

martensitic transformations, which also often happens in (α+β)-alloys. The transition

from β to martensite is responsible for an acicular structure in cooled or cooled and

aged Ti-alloys. Martensitic transformation takes place through the sliding movement of

atoms, which provides a homogeneous transformation from β-phase to martensite α-

phase [32]. There are several types of martensites formed on Ti-alloys, but the two

main ones are α' (alpha prime), which appears as an acicular phase with an hcp

structure, but is similar in the microstructural appearance of the acicular α; and α’’

(alpha double prime), a supersaturated orthorhombic phase. Both are formed by

cooling and decompose in successive aging in α and β phases [56]. The α' typically

arises when the material is cooled from above its Tβ to below its Ms. The formation of

α' is characterized by a transformation from the β to a metastable hcp structure. This

transformation involves a shear deformation mechanism, resulting in a distortion of the

crystal lattice. Hence, α' is often responsible for the improvement in mechanical

properties, such as increased yield strength and hardness, observed in quenched Ti-

alloys. The α", in turn, is a secondary martensitic phase that can form typically during

further cooling below the temperature at which α' forms. This results in a further

transformation of the martensitic structure, leading to the formation of a more stable

orthorhombic phase. Like α', α" forms especially in alloys where the β-phase elements

are in high quantity, due to athermal slip or deformation in the plane of {110} [57].

Martensite α’’ formed in this way may precipitate primarily along the boundaries of β-

grain [58]. α" possesses unique mechanical properties compared to α'. While it may

exhibit lower strength and hardness than α', α" often provides improved ductility and

toughness. This may be advantageous in applications requiring a combination of

strength and toughness [35,56,59].

Another non-equilibrium phase, a non-compact hexagonal phase called omega

(ω), is metastable in nature and forms in Ti-alloys with elements that tend to stabilize

the β-phase. Depending on the conditions, the formation of the ω-phase can occur

during quenching from high temperatures in the β-phase region or during aging of the

quenched alloy. The addition of α-stabilizer elements such as Al and neutral elements

such as Zr or Sn, the latter acting as an α-stabilizer in the presence of other α-

stabilizers, suppress the formation of the ω-phase to promote the formation of the α-

phase [33].

2.1.1 The β-metastable Ti-5553 alloy

The Ti-5553 alloy has been recently developed with excellent mechanical

properties to improve processability. This alloy is a variation of the Russian VT22 alloy

(Ti-5Al-5V-5Mo-1Cr-1Fe) [60] and is an alternative to replace the Ti-10V-2Fe-3Al (Ti-

1023) in some aircraft structural components [61]. Ti-5553 exhibits excellent strength

characteristics, even higher than the currently used (α+β)-Ti-alloy grades, such as Ti–

64. Depending on processing, the minimum yield strength of Ti-5553 can vary between

about 800 and 1200 MPa, in comparison with that of Ti-64 varying between about 850

and 1050 MPa [34].

The position of Ti-5553 compared to other Ti-alloys based on their [Mo]eq and

[Al]eq contents and its position on the pseudobinary diagram for the Ti-β-stabilizer

element system is given in Fig. 7.

Fig. 7: a) Classification of Ti-5553 compared to other usual titanium alloys based on

their [Mo]eq and [Al]eq contents [62]; b) Pseudobinary diagram for the Ti-β-stabilizer

element system with the indication of the most used commercial titanium alloys [34].

(Reprinted with permission of the copyright owners).

The decomposition of the fully β structure, which in turn can be retained on

quenching into a bimodal microstructure α+β during aging, makes Ti-5553 an age-

hardenable alloy. The slow precipitation kinetics of the α-phase gives this alloy an

advantage over other β-Ti-alloys. Typical microstructures obtained when aging the Ti-

5553 alloy after homogenization above the Tβ are presented in Fig. 8. Both the amount

and morphology of the α-phase are modified by changing the aging temperature. As

well explained by Clément et al. [62], the nucleation of α-phase is scarce at 800°C

(i.e., about 70-60°C below the transus temperature) and the α-precipitates remain quite

equiaxed, as can be seen in Fig. 8a. By decreasing the temperature to 700°C, an α-

film can be observed covering prior β-grain boundaries almost completely (Fig. 8b)

with many parallel plates growing toward the grain interior together with intragranular

α-precipitates within the β-matrix. Down to 500°C, an entanglement of residual β and

fine α precipitates homogeneously nucleated within the whole grains replace the now

completely transformed β-matrix (Fig. 8c).

Fig. 8: Scanning-electron micrographs of solution-treated and aged Ti-5553 alloy at

(a) 800°C, (b) 700°C, and (c) 500°C. (Reprinted from [62], with permission of the

The relevant mechanical properties of the β-annealed and bimodal

microstructures of this alloy, discussed by Bücher et al., 2007, as cited in Lütjering and

Williams [60], are shown in Table 1. It can be seen that the yield strength (σ0.2) and

the ultimate tensile strength (UTS) are quite similar for the β-annealed structure and

the bimodal structure. The tensile elongation as well as the high cycle fatigue (HCF)

strength (σ107) is higher for the bimodal structure as compared to the β-annealed

structure. The fracture toughness is about the same for both microstructures. Values

of E of β and α phases, Poisson's ratio, and Tβ temperature of α ↔ β transformation

are also given in Table 2.

Table 1: Mechanical properties of Ti-5553 with β annealed and bimodal

microstructures [60].

Microstructure

σ0.2

(MPa)

UTS

(MPa)

Elongation

(%)

KIC

(MPa.mm1/2)

σ107

(MPa)

β-annealed

1100

1145

6.4

66.1

500

bimodal

1090

1150

13.4

65.8

575

Table 2: Elastic modulus of β and α phases, Poisson's ratio and Tβ temperature of the

transformation α ↔ β [63].

Parameter

Value

Eβ

~ 70 − 90 GPa

Eα

~ 115 − 125 GPa

~ 0.297 − 0.3

Tβ

845 – 860 °C

2.2 Heat treatments for β-metastable titanium alloys

Ti-alloys can be heat-treated in either the single β-phase field or the two-phase

α+β field, therefore giving wide scope in altering their microstructure and directly

optimizing their mechanical properties. There are several reasons for the heat

treatment of Ti-alloys: (i) to reduce residual stresses developed during fabrication

(stress relieving); (ii) to produce the most acceptable combination of ductility,

machinability, and dimensional and structural stability, especially in (α+β)-alloys

(annealing); (iii) to increase strength by STA; and (iv) to optimize special properties,

such as fracture toughness, fatigue strength and high-temperature creep strength [56].

Generally, the α and near α-alloys are either non-heat-treatable or cannot have

a significant change in the microstructure by heat treatment. Heat treatment of the α-

alloy (e.g., CP-Ti) from the β-phase region results in a microstructure that contains

other phases along with the hcp α-phase [33]. Stress relief and annealing are most

likely for this Ti-alloy classification. As for the (α+β)-alloys, they can have their phase

compositions, sizes, and distributions manipulated within certain limits by several heat

treatments [56,59], for example, mill-annealing [60] and β-solution treatment plus over

aging (STOA) [64].

The heat treatment approaches for β-alloys, in turn, are significantly different

from those for the α and near α-alloys, and (α+β)-alloys. The β-metastable Ti-alloys

not only can be stress relieved or annealed, but also can be solution treated and aged.

The STA steps promotes precipitation hardening, which can be applied at various

temperatures and times depending on the chemical structure of the material and the

desired mechanical properties [20,59]. The precipitation hardening treatment applied

to β-metastable Ti-alloys aims to increase ductility, fracture toughness, yield and

tensile resistances, as well as creep, impact and corrosion resistances [65].

The Table 3 and Fig. 9 summarize the heat treatments commonly applied to β-

metastable Ti-alloys and their characteristic resulting microstructure.

Table 3: Several heat treatments for Ti-alloys and their characteristics [59,65].

Thermal treatment

Characteristics

Stress relief

This is a low-temperature heat treatment generally

conducted in the temperature range of 370 – 815 °C. It is

generally used after prior fabrication steps or prior heat

treatment operations to remove residual stresses, which

could cause distortion during machining or after a forming

operation. It would not change the microstructure.

STA

β-metastable alloys will generally be solution-treated above

the Tβ. The solution treatment will generally be followed with

rapid cooling, sometimes using air cooling depending on

the alloy. The aging of these alloys will result in α-phase

precipitation. For Ti-5553, the STA condition consists of 10–

20 vol.% globular αpi in a matrix of retained β-phase. The

scale of the aged β is quite fine and the individual α-

precipitates are generally not visible by optical microscopy.

STA-Ti-5553 is characterized by high strength and good

fatigue behavior (Fig. 9a).

BASCA

This heat treatment results in a high-volume fraction of

lamellar α, which results from controlled cooling from β-field

temperatures. The BASCA-Ti-5553 alloy is characterized

by fracture toughness values double that of the STA

condition (KIC > 65 MPa·m1/2) and intermediate strength

properties (Fig. 9b).

Cryogenic treatment

Cryogenic treatment involves subjecting the alloy to

extremely low temperatures (typically below -150°C). This

treatment can enhance the stability of the β-phase and

improve wear resistance, fatigue strength, and creep

resistance. The α-precipitates are formed in alloys that

were cooled below the martensitic temperature during

cooling to cryogenic temperatures. The α-phase exhibits a

needle-like structure (see Fig. 9c).

Cold working and aging

This process is used for Ti-3Al-8V-6Cr-4Zr-4Mo (β-C)

spring wire. Before the last drawing operation, the material

is solution-treated above the Tβ, resulting in the

microstructure shown in Fig. 9d. It is then given the ﬁnal

drawing pass with something on the order of 30% cold

work. After forming the spring, the alloy is then aged.

Fig. 9: a) Micrographs of Ti-5553 in STA condition [61]; b) Light optical image of Ti-

5553 in the BASCA condition showing coarse α-platelets (dark) and minor amounts of

retained β (light) [54]; c) SEM micrographs of TB8 alloy (Ti-15Mo-3Al-2.7Nb-0.2Si)

treated by deep cryogenic treatment before aging treatment [66]; d) Light optical

micrograph of β-C in the STA condition showing precipitation of fine α-phase (light) in

β-grains (dark) after cold working and aging [54]. (Reprinted with permission of the

STA and BASCA are the processes most often used to improve the strength of

AM Ti-5553 [67–69]. The two-stage STA treatment primarily leads to an increase in

strength. Solution annealing just below the β-transus temperature (~620 °C) results in

globular primary α-phase particles dispersed in the β-matrix, with diameters in the

micrometer range. The subsequent aging treatment, which is followed by rapid cooling,

produces a transformed β-phase with fine, homogeneously distributed, needle-like

secondary α-phase particles with diameters in the nanometer range [70]. The BASCA

heat treatment aims to achieve maximum fatigue and fracture toughness [71,72]. The

alloy is annealed in the β-field, then slowly cooled, aged in the (α+β)-range, and again

slowly cooled down to room temperature. The resulting microstructure contains an

acicular, plate-like α-phase with a basket-weave structure in a β-matrix. The single

platelets are several micrometers long and around 500 nm thick [59].

For conventionally manufactured Ti-5553, STA, and BASCA heat treatments

have resulted in substantially superior static and fatigue properties relative to other

titanium alloys [73,74]. Ultimate tensile strength values exceeding 1240 MPa Ti-5553

[75,76] and fatigue endurance strengths in the range of 650-700 MPa were achieved

[77]. Smaller secondary α-precipitates in STA-Ti-5553 alloy enhanced yield and tensile

strength but reduced ductility, nevertheless overall improving fatigue strength [73]. Ti-

5553 heat treated via BASCA exhibited high Vickers hardness [74]. STA and BASCA

treatments also effectively influence the mechanical properties of AM-Ti-5553.

Whereas "as-built" samples showed 830 MPa tensile strength post-LPBF-M, BASCA

increased yield strength and elongation by 219% and 121%, respectively, and STA

treatment achieved 78% higher ultimate tensile strength [78].

2.3 Titanium and its alloys as implant materials

From a purely mechanical perspective, metal is the most commonly used

material variant for implants due to its superior properties compared to polymers and

ceramics. While polymers have limited mechanical strength and are only suitable for

certain situations, ceramics, although they have some advantages, are too brittle and

have poor fracture toughness [31]. However, the mechanical requirements are not the

only criterion for selecting the proper implant material. In order to find a place in medical

applications, materials must have some specific properties. In the case of implants,

one of the most important requirements for a material is biocompatibility, followed by

corrosion resistance. According to DIN EN ISO 10993-1 [79], materials that have no

negative impact on their metabolism in direct contact with living tissues are

biocompatible. For implants, this means that they must not provoke a defense reaction

of the immune system.

Ti and its alloys are hence in common use for implants because of their excellent

combination of mechanical strength, corrosion resistance, and biocompatibility [2,20].

The first relevant applications of Ti as a biomaterial for dental and surgical devices go

back to the end of World War II, following the Ti manufacturing progress [80]. CP-Ti is

widely considered to be the most compatible metal for the human body because of the

presence of a stable and an inert oxide layer, which spontaneously forms on its surface

[30]. Among the Ti-alloys, Ti-64 is currently the most widely used titanium alloy for

bone implants. This (α+β)-alloy combines excellent high specific tensile strength and

fatigue strength with high chemical stability (corrosion resistance) and biocompatibility

under in vivo conditions [48,56,60]. Even though Ti-64 has a relatively low E (E ≈ 110

GPa) compared to the long-used stainless steel 316L (E ≈ 210 GPa) and Co-Cr-alloys

(E ≈ 240 GPa), it is still significantly higher compared to trabecular and cortical bone

[81] due its fraction of α-phase [22]. This considerable difference between the E of

bone and the implant will result in a high mismatch with their stiffness, leading to the

so-called “stress shielding effect”, where certain bone regions are shielded from

loading [2,22]. These lower or unloaded bone regions are then resorbed, eventually

leading to loosening of the implant. As a consequence, the implant needs to be

exchanged, either due to pain or, even worse, following fatigue failure. In order to

reduce the probability of occurrence of these postoperative conditions, a lower E can

be achieved by an increase in the fraction of the β-phase in the microstructures of Ti-

alloys. Therefore, compositions of alloys with higher levels of β-stabilizing elements

are being considered and further developed to optimize the mechanical properties of

bone implants [1,22,23,48,55]. One of these Ti-alloys is the β-metastable Ti-5553,

which not only presents a lower E, but also an excellent processability by AM such as

LPBF-M [82–84]. This alloy demonstrates ease to be additively manufactured and

shows potentially favorable mechanical properties for orthopedic applications [26].

Heat treatment may increase the Ti-5553 alloy strength to values comparable to (α+β)-

Ti alloys without sacrificing the low stiffness achieved by β-stabilizers alloying [85,86],

which can also be beneficial for implant applications. The applicability of AM Ti-5553

as implant material has been recently reported [87]. After manufacturing, the alloy was

mechanically suitable for medical implants, and the biological compatibility was tested

through saos-2 (human bone) cell culture to determine the level of biological activity

on the surface of the alloy. The researchers concluded that the Ti-5553 specimens

performed very similarly to the established Ti-64 alloy, suggesting a strong potential

for Ti-5553 as a bone replacement implant.

2.4 Additive manufacturing of titanium alloys

This section provides a brief overview of AM, as this manufacturing approach is

often used to produce Ti-alloy implants. However, it is important to keep in mind that

this section does not provide an extensive examination of the AM technique, as the

current work is not centered on its development or assessment. This description aims

to provide rather a concise understanding of AM concerning the production of Ti-alloys.

AM represents a material-saving alternative to conventional processes such as

casting, forging, or subtractive manufacturing processes. It is based on the principle of

slicing a digital solid model in multiple layers to create a toolpath, uploading this data

in the machine, and building the part up layer by layer following the sliced model data

using a heat source (laser, electron beam or electric arc) and feedstock (metal powder

or wire) [42,88] This technology results in the decisive advantage in producing

completely new shapes and filigree, which enables significant weight savings and

improves the functionality of the component. Although AM is not yet suitable for series

production in many areas due to high costs, there are already a large number of

different rapid technologies that are suitable for many interesting areas of application.

Most of the studies reveal that the mechanical properties of AM material are as good

as or better than those of conventionally fabricated [88–90].

The AM technology is particularly suitable for medical applications, as it enables

the manufacture of patient-specific products that can be precisely adapted to the

respective stress situation. Fig. 10 shows a brief schematic development cycle to

additively manufacture an implant material.

Fig. 10: Brief schematic diagram of the product development cycle to additively

manufacture an implant (Reprinted from [91], with permission of the copyright owner).

AM researchers are currently developing a wide range of biocompatible

feedstock material and processing systems for medical devices, such as hip, knee, or

articular cartilage joints [92–94]. Ti-alloys stand out as the most used AM materials

when it comes to biomedical applications (see Table 4).

Table 4: FDA

cleared AM-materials for biomedical applications [94].

Material

Application

CP-Ti

Ti-64

Ti-6Al-7Nb

Ti-5Al-2.5Fe

Ti-15 Zr-4Nb-2Ta-0.2Pd

Ti-29Nb-13Ta-4.6Zr

83%–87%Ti-13%–17%Zr

(Roxolid)

316L

Co-Cr-Mo, Co-Ni-Cr-Mo

NiTi

PMMA, PE,

PEEK

SiO2/CaO/Na2O/P2O5

Zirconia

Al2O3

Ca5(PO4)3(OH)

Screw and abutment

Artificial valve, Stent, Bone fixation

Crowns, Knee joint, Hip joint

Spinal implant

Crown, Bridges, Dentures, Implants

Crown, Bridges, Dentures

Knee joint, Hip joint, Surgical tools, Screw

Artificial valve, Plates, Bolts, Crowns, Knee joint, Hip

joint

Catheters, stents

Dental bridges, articular cartilage, Hip joint femoral

surface, Knee, Joint bearing surface, Scaffolds

Bones, Dental implants, orthopedic implants

Porous implants, Dental implants

Dental implants

Implant coating material

The U.S. Food and Drug Administration (FDA) is responsible for protecting the public health by

ensuring the safety, efficacy, and security of human and veterinary drugs, biological products, and

medical devices; and by ensuring the safety of our nation's food supply, cosmetics, and products that

emit radiation (www.fda.gov).

For the processing of metallic materials such as Ti-alloys, LPBF-M, initially also

known worldwide as selective laser melting (SLM)

, has been established in recent

years, which will be discussed in more detail hereafter. This section will first look at the

functionality and some parameters that are important for the process. Research

studies focusing on the LPBF-M process to manufacture Ti-alloys are then

subsequently described to underline the relations between the AM process and

material microstructures and properties.

2.4.1 Laser Powder Bed Fusion of Metals (LPBF-M)

LPBF-M is a technology based on the principle of laying down an amount of

metal powder on the build platform and scanning the bed of powder with a laser as a

heat source that either partially or completely melts the powder in the path of the beam

and re-solidify and bind it together as it cools off (ASTM specification F2924-12a and

13 for Ti-64 and Ti-64 ELI grade respectively) [88]. The manufacturing accuracy of the

process is largely dependent on the size of the powder particles used [17,42,84]. After

the melt solidifies, a solid layer of material remains, which corresponds to the first layer

of the component. The base plate is then automatically lowered by one layer thickness

and a new powder layer is applied. This process is repeated continuously. The layer

thickness can vary from 15 to 150 µm [21]. Fig. 11 shows the schematic representation

of the LPBF-M machine.

The LPBF-M process can be compared to a continuous welding process, where

numerous welds are repetitively performed and stacked together, resulting in the

creation of a three-dimensional geometric space. Although similar to welding, the

fundamental distinction lies in the objective of the welds: in welding, they connect

multiple distinct geometries, while in LPBF-M, the welds themselves form the

geometry. Moreover, the molten pool dimensions in LPBF-M are approximately

0.1mm³, indicative of the smaller scale in which this process operates compared to

welding [97].

SLM is according Trademark entry 30094322, Deutsches Patent- und Markenamt (engl. German

Patent and Trademark Office) a word mark of the SLM Solutions Group [95]. In the literature, SLM is

used as an abbreviation for the process of selective laser melting (e.g. [14,17,21,96])

Fig. 11: Schematic representation of a LPBF-M machine. (Reprinted from [21], with

permission of the copyright owner).

Depending on the requirements, the laser powder bed fusioned component can

then be post-processed or used directly. Post-processing of the component surface is

advisable in many cases, as it is mostly rough, similar to that of a product made by

sand casting [10,98,99]. The effect of surface roughness on the fatigue performance

of LPBF metals has been investigated in recent years. For example, Uhlmann et al.

[10] showed that the bad surface quality of LPBF 316-L stainless steel parts is a key

problem of this technology, not only for functional surfaces but also for the fatigue life.

Turned specimens showed the best fatigue behavior while as-built specimens showed

the worst. Sanaei and Fatemi [98] reviewed the effect of post-processing conditions on

the surface of LPBF-M Ti-64 reporting that the surface roughness effect is detrimental

and primary on the fatigue performance even in the presence of relatively large internal

defects and for various microstructures. The work from Greitemeier et al. [99]

corroborates the findings from the other two investigations cited above. The authors

concluded that the fatigue properties of AM Ti-64 samples are dominated by surface

roughness effects. Regarding Ti-5553, various post-processing processes (vibratory

finishing, shot peening, electrochemical polishing, and turning) have different effects

on the surface quality and topography [100]: (i) Shot peening only removed the melted

powder particles on the surface, maintaining the surface characteristics after the LPBF-

M process. The profile ordinate Ra was reduced by 33% and the maximum height of

the profile average maximum surface height (Rz) was roughly halved. (ii)

Electrochemical polishing showed comparable surface parameters to the surface after

shot peening, although the surface was processed more homogeneously in

comparison to shot peening. (iii) Finishing by vibratory grinding was able to significantly

improve the surface quality, with relatively little material being removed due to the

process. The surface parameters were reduced to Ra = 0.2 µm or Rz = 1.65 µm. (iv)

Turning achieved surface qualities are up to 50% better than vibratory grinding. Only

turning increased the modulus of elasticity, while no influence on the tensile strength

could be determined for vibratory grinding finishing, shot peening, or electrochemical

polishing. In contrast, the influence of all these post-processing methods on the

resistance to cyclic stress was observed. While the influence of shot peening and

electrochemical polishing only became clear at lower stress amplitudes, the resistance

to cyclic stress could be increased over all stress amplitude ranges during turning and

vibratory finishing.

Heat treatment is also usually performed after LPBF-M to optimize the

mechanical properties of the components, substantially affecting their evolution of

microstructure and porosity [101]. Different heat treatments on the microstructural and

mechanical aspects of Ti components produced by the LPBF-M process demonstrated

the be a positive influence on their failure response and strength [96,102]. After heat

treatment of porous Ti-64, fine acicular α′ martensitic microstructure was transformed

to a mixture of α and β phases, hence changing its ductility as well as some mechanical

properties such as hardness, plateau stress, and first maximum stress [103]. Gerlitzky

[100] investigated three different heat treatment strategies for Ti-5553 manufactured

by LPBF-M: ST, STA, and BASCA. The not-heat treated state presented

Rm = 820 MPa, At=14 % and HV=323. The heat treatment STA did lead to an increase

in the tensile strength (Rm) by around 80% and in the E, but only a low elongation

capacity with a percentage of total elongation at fracture of 1.5%. The material shows

a very brittle material behavior. The measured Vickers hardness increases to over 536

HV. After ST, the material properties only improved to a very small extent change. The

tensile strength increased to Rm = 880 MPa, with the total elongation at break

decreasing to At = 12%. The Vickers hardness HV also remained almost unchanged.

BASCA heat treatment increased Rm to approx. 30% (Rm = 1085 MPa), while the total

elongation at break was slightly reduced to At = 13.6%. The Vickers hardness was

increased to 404 HV. However, ST, STA, and BASCA have no significant influence

under cyclic loading. Minor differences could only be observed with increased stress

amplitudes (σa).

2.4.2 Processability of Ti-5553 via LPBF-M

This section explores the recent literature considering the processability of

Ti-5553 via LPBF-M, its resulting microstructure, and its properties. Table 5

summarizes the process parameters to produce Ti-5553 parts using this AM approach.

Relatively good values obtained from different processing conditions confirmed the

processability of Ti-5553 [82,83]. The mechanical behavior analysis of cubic samples

provided a UTS of about 800 MPa, strain up to 14%, and density near 99.95% [82].

The resulting microstructure after LPBF-M consisted of hcp α-phase embedded in the

bcc β-phase, with no indications of a third phase [83,100]. The high amount of β-phase

was attributed to suppress the possibility of martensitic α'-phase formation [104]. The

micrographs exhibited elongated grains that grew throughout the layers, unlike other

Ti-alloys where cellular solidification or segregation typically occurs [97]. Large β-

grains elongated in the building direction containing dispersed α-phase precipitates

were observed for all processing conditions [82,83]. The α-phase exhibited morphology

of small plates sizing around 1μm, which acted as reinforcement particles by

enhancing the mechanical properties of LPBF-M Ti-5553 [100].

The processability of the Ti-5553 using LPBF-M was also investigated based on

different particle size distributions [84] and on the effect of volumetric energy density

(VED) on microstructure, density, roughness, and hardness [105]. As the powder size

increased, the surface roughness also increased because of particles still adhered to

the surface. Stable process parameters for a reference grain fraction of 10–63μm were

developed by keeping the laser power and layer thickness constant at 45 and 30 µm,

respectively [84]. For low and high VED values, lack of fusion and keyhole pores was

observed, respectively, although a density of 99.92%, the highest density value

reported so far for LPBF-fabricated Ti-5553 parts, was achieved. A smoother surface

presenting roughness of less than 12μm was obtained for the specimen manufactured

at VED of 112 J/mm3, while the presence of ω-phase was revealed in the specimens

manufactured at medium VED. At high VED values, the presence of α-phase needles

was observed in the specimens, resulting in a higher but nonuniform hardness

compared with the specimens manufactured at low and medium VEDs [105].

Table 5: LPBF-M parameters for manufacturing Ti-5553.

Preheating

Power

Laser

spot

Layer

thickness

Hatch

spacing

Scanning

speed

Reference

(°C)

(W)

(μm)

(mm/s)

N/A

60-140

N/A

100-180

[82], [83]

200

400

80-120

N/A

[84]

N/A

250

N/A

120

1000

[100]

N/A

100-275

N/A

350-1080

[105]

N/A

100

N/A

600

[104]

2.5 Fatigue

Fatigue is a process of material deformation and damage accumulation due to

repeated loading and unloading leading to failure of the material at a stress much lower

than its tensile or yield strength [106]. Fatigue failures are the most common types of

fractures in components and machines and probably make up about 90% of all

fractures [107]. Therefore, the failure of metals subjected to alternating loads has been

a major concern for engineers and material scientists for at least 170 years. Jean Victor

Poncelet, a designer of cast iron axles for mill wheels, officially used the term "fatigue"

for the first time in 1839 [108]. Poncelet postulated that the axles became tired, or

fatigued, after a period of use before breaking. These failures became a serious

problem during the 19th century, as machines with rotating or vibrating parts began to

be used more and more. This represented a major challenge for engineers, as failures

occurred unexpectedly in seemingly safe loading that would not cause static failures.

The failure of the Versailles train in 1842 is considered the milestone that started the

path to understand the mechanism of cyclic fatigue [109]. August Wöhler, who worked

as an engineer for the Prussian rail system, was the first to carry out systematic fatigue

studies. His work was a direct result of the investigation of the Versailles train failure

and lasted about 12 years, being the first scientific investigation of fatigue failure.

Wöhler conducted his experiments under rotating bending conditions, testing axes in

the laboratory to failure under alternate loading. The most important result of his work

was the discovery of a well-defined stress amplitude (now called the fatigue limit) below

which failures would not occur. He published his findings in 1870 [110]. Since then the

performance of materials in fatigue is usually characterized by the "S-N curve", also

known as "Wöhler curve", which is a graph of stress magnitude (S) by the number of

cycles (N). A curious fact is that Wöhler never presented his data in graphic form, but

generally only in extensive tables [111].

For the next thirty years, the microscopic reasons that cause fatigue failure were

not further systematically studied, remaining unclear until 1903, when Ewing and

Humfrey [112] published their work on the fracture of metals under repeated

alternations of stress. They observed slip bands in which microcracks had developed

on the surface of fatigued steel, coming to the main conclusion that fatigue damage

occurs due to the accumulation of various irreversible cyclic plastic micro-

deformations. Today, there is a very detailed understanding of the microstructural

mechanisms that lead to fatigue damage [113]. Although the traditional technique of

uniaxial mechanical testing, which gives us information on the macro-scale mechanical

properties of fatigue, has proven itself, it is quite expensive both in time and in money

and requires a large number of test samples [114]. In the last decades, the

development of alternative non-destructive testing techniques that enable the use of

smaller quantities of test materials, ease of conduct, and high resolutions of

mechanical properties measurement have increased to characterize the fatigue and

cyclic deformation behaviors in a short term.

2.5.1 Corrosion-fatigue

Numerous factors can influence the fatigue behavior of AM implant materials,

including stress levels, the presence of defects or impurities, and environmental

conditions. One particularly significant degradation process is corrosion-fatigue, which

combines the effects of corrosion and cyclic loading to accelerate material failure [115].

Corrosion is a chemical process that occurs as a result of the interaction between a

material and its surrounding environment, leading to the formation of corrosion

products and a gradual loss of material mass [116]. On the other hand, fatigue is a

mechanical process influenced by cyclic loading, which induces local plastic

deformation and the initiation and propagation of cracks [113]. When corrosion and

fatigue interact synergistically, the resulting corrosion-fatigue behavior can be more

detrimental than the sum of their individual effects. This is not only due to the

accelerated degradation rate but also the complex interaction mechanisms between

corrosion and fatigue processes, which distinguishes it from pure fatigue that occurs

in the absence of a corrosive environment such as dry inert gas or vacuum (see Fig.

12). Corrosion fatigue is hence a process dependent on the interactions among

loading, environmental and metallurgical factors. Other variables such as chemical

composition, heat treatment, and microstructure further contribute to this complex

phenomenon [116,117].

Fig. 12: Distinction between corrosion-fatigue and pure mechanical fatigue.

(Reprinted from [116], with permission of the copyright owner).

The effects of corrosive environments on the fatigue behavior of metals were

studied as early as the 1920s, although the first paper to describe corrosion-fatigue

experiments was published by Haig in 1917 [118]. No further attention was paid to the

subject until about 1925 when investigations on this subject gained more attention

[115]. Since then the investigation of the combined effects of both corrosion and fatigue

damage mechanisms became necessary to better understand the corrosion fatigue

phenomenon. The investigations have revealed that damage contribution of corrosion

to fatigue crack growth primarily occurs during crack extension in the load-increasing

part of the load cycles. The loading rate in this part of the cycle is hence important. As

a consequence, the load frequency proves to be determining [28].

Implant materials are subjected to cyclic stress, and encounter complex

environments during their service life. As a result, their effectiveness relies heavily on

their ability to withstand corrosion-fatigue. Hence, it is crucial to research how the body

environment affects the fatigue and cyclic behavior of implant materials. Numerous

investigations have focused on the complex interaction between fatigue and corrosion

of Ti-alloys in quasi-physiological media. The number of cycles to failure of CP-Ti

decreased under fully reversed, uniaxial loading in 0.9 wt.% NaCl solution and in serum

at 37 °C as compared to loading under ambient conditions [119,120]. Axial tests

revealed a lower fatigue life in oxygen-deficient as compared to oxygen-saturated

Ringer's solution, and the fatigue strength decreased significantly with increasing

maximum numbers of cycles [121]. For (α+β)-Ti-alloys, crack growth rates were

significantly higher in 3.5 wt.% NaCl and Ringer's solutions compared to air [122,123].

Furthermore, AM (α+β)-Ti-alloys produced by LPBF-M and directed energy deposition

(DED) techniques experienced a clear reduction in their fatigue life by around 50% in

Ringer’s solution and simulated body fluid (SBF) in comparison with air [124,125], and

the crack growth rate increased significantly in 3.5 wt.% NaCl solution in comparison

with air, artificial saliva, and Ringer’s solution [126]. β and β-metastable Ti-alloys

demonstrated an enhanced resistance to corrosion and wear when exposed to a NaCl

solution [127]. However, when subjected to a synthetic cell culture medium (Eagle's

solution), their fatigue strength depended on their specific chemical composition [128].

To the best of the author's knowledge, no research has been conducted on the

corrosion-fatigue behavior of neither conventionally nor additively manufactured Ti-

5553.

Repeated mechanical stress cycles in quasi-physiological media can result in

corrosion-induced fatigue damage. Microcracks, pits, and other forms of damage can

happen on the implant surface, which eventually can lead to corrosion [129]. Based on

simple electrochemical measurements, it is possible to detect the surface damage

onset in environmental-assisted fatigue tests by using an open circuit setup that allows

the simultaneous monitoring of the corrosion current increase and/or potential

decrease signals. It is important to note that the measurements are performed on

passivating materials. The free corrosion potential (Ucorr) is measured between the

working electrode (WE, specimen) and a reference electrode (RE, usually Argenthal

Ag/AgCl, +207 mV to normal hydrogen electrode), while the free corrosion current (icorr)

is determined by connecting the WE with a counter electrode (CE) via a low noise

multimeter. The CE is a specimen consisting of the same material with the same

surface condition as the WE. As long as the loaded surface of the specimen remains

intact, CE and WE remain on the same electrochemical potential versus the RE. When

detecting considerable changes in current and potential measurements, there is an

indication of modification on the WE surface, such as the formation of slip bands,

intrusions and extrusions, or microcracks. These alterations result in negative potential

peaks due to the alloys' more negative equilibrium potential in the non-oxidized state

compared to the oxidized state. Previous studies have demonstrated that surface

damage can be detected in Ti-alloys undergoing axial or rotating bending fatigue tests

in physiological media by measuring changes in the free corrosion potential and/or

current [130,131]. Fatigue crack up to 50μm length could be detected by a decrease

in corrosion potential for the binary Ti-6Al-7Nb alloy under rotating bending in Ringer’s

and Hanks' solutions. The authors observed that the surface activates and reacts with

surrounding constituents due to localized plastic deformation before fatigue crack

initiates. Microscopic investigations revealed three stages of damage propagation

correlated with electrochemical measurements [130]. The crack propagation rate as

well as the activation rate of the fresh Ti-surface was directly correlated to the corrosion

current response for the binary titanium alloys Ti-64and Ti-6Al-7Nb tested by axial

stress-controlled constant amplitude and load increase tests, as well as in rotating

bending tests in oxygen-saturated Ringer’s solution. Pronounced crack propagation

starts when the critical stress intensity is reached at the crack tip or when microcracks

coalesce into a larger crack [131].

2.5.2 Nano-fatigue

Small-scale depth-sensing indentation techniques, such as nanoindentation

(NI), have become highly valuable for observing and characterizing the mechanical

behavior of materials at a local scale. This versatile technique has been applied to both

thin films and bulk materials to determine various mechanical properties related to

elastic and plastic deformation [132]. The NI technique involves pressing a sharp

indenter with known properties (e.g. a Berkovich, cube-corner, or spherical-shaped

diamond tip) into the surface of the material under a constant load (P) continuously

applied to the indenter, and the displacement (D) of the indenter tip into the specimen

is measured. The sample is then "unloaded" by withdrawing the indenter tip until the

load on the indenter tip reaches 0 after a certain holding period. The well-established

Oliver-Pharr method [133] is utilized to extract reduced modulus and hardness values

from the load-displacement curves (Fig. 13).

Fig. 13: Load-displacement curves from a nanoindentation experiment with maximum

load Pmax and depth beneath the specimen-free surface Dmax. The depth of the contact

circle Dc and slope of the elastic unloading dP/dD allow specimen modulus and

hardness to be calculated. Dr is the depth of the residual impression, and De is the

displacement associated with the elastic recovery during unloading. (Adapted from

[133], reprinted with permission of the copyright owner).

This technique offers several advantages over traditional macroscale

indentation methods, making it a powerful tool in materials characterization. One of the

key advantages of nanoindentation is its ability to measure the mechanical properties

of small volumes of material. Traditional methods, such as tensile testing or bending,

require relatively large samples, which may not always be available for nanoscale

structures or coatings. In contrast, nanoindentation only requires a small portion of the

material, resulting in less material waste and the ability to test localized regions of

interest. Furthermore, materials often exhibit size-dependent mechanical behavior at

small length scales due to various size effects, such as grain boundaries, dislocations,

or interfacial effects. By testing materials at the nanometer scale, researchers can gain

a better understanding of their mechanical behavior and discover new phenomena not

observed at larger scales. NI also offers the advantage of high spatial resolution. The

size and geometry of the indenter tip used in NI experiments enable precise positioning

and control of the applied load. This spatial resolution allows researchers to perform

indentation tests at specific locations of interest, such as nanostructured regions. The

ability to probe specific regions provides valuable information about the local

mechanical properties and helps in designing materials with tailored properties.

In recent years, innovative techniques have been employed to evaluate the cyclic

mechanical properties of a wide range of materials at the nano-scale. One such

technique gaining significant attention is cyclic nanoindentation. Unlike traditional

quasi-static nanoindentation [133], cyclic nanoindentation involves the repetitive

application and removal of the indenter, providing valuable insights into the dynamic

mechanical behavior and response of materials. Usually, the indentation load is

precisely controlled at each cycle, allowing analysis of the displacement resulting of

loading-unloading processes. This cyclic loading process results in a series of

indentation-displacement curves (hysteresis loops), which capture the response of the

material to varying loading conditions. Fundamental phenomena can be then analyzed,

such as cyclic creep (incremental plastic deformation evolution during consecutive

cycles) and cyclic softening/hardening, which are critical for assessing the

performance of the material under repetitive loading conditions.

Cyclic nanoindentation offers several advantages, such as providing a fast and

relatively simple approach to overcome the limitations of conventional macro fatigue

testing, which requires a lot of material and a high number of specimens [132].

Furthermore, the method allows probing differences in the properties of single phases

as well as the influence of the local microstructure in a quasi-non-destructive way [133–

135]. Recently, cyclic nanoindentation has been used to assess the local fatigue

properties of metals. The method has mainly been applied to the characterization of

thin films [136–139]. The low number of publications reporting on bulk materials usually

address the cyclic deformation behavior up to relatively low maximum numbers of

cycles in the range of 10 to 300 only [140–143]. For closed-cell aluminum foams,

cyclically indented six times with increasing loads, uniaxial stress-strain plots were

constructed from the penetration depths and forces of the single cycles, and properties

of individual microstructural phases and their volume fractions were identified [141].

The progression of indentation depth with the number of cycles exhibited two distinct

regions for solution-treated and aged AZ61 Mg-alloy [142] and duplex stainless steel,

cyclically nanoindented in load control up to 300 cycles [143]. A primary or transient

stage where the indentation depth rate decreased with the number of cycles preceded

a secondary “steady-state” region suggesting a balance between cyclic hardening and

softening. In displacement-controlled tests up to a maximum number of cycles of 100,

Ti-64 with a duplex microstructure with more than 90% α-phase exhibited cyclic

softening for both, primary and lamellar α-phase, indicated by a decrease in peak

stress [140]. In turn, there are only two reports so far on cyclic nanoindentation fatigue

of metals up to a much higher maximum cycle number of 105 [144,145]. The tests

performed on cross-sections of struts extracted from A356.0 aluminum alloy open-cell

foam [144] and of Mg–SiC nanocomposites [145] showed significant influences of the

phase composition in a micrometer-sized interaction volume below the indent on the

cyclic deformation behavior.

3 MATERIAL

Cylindrical standard fatigue specimens according to DIN 50113 [146] (Fig. 14a)

and customized specimens for the corrosion-fatigue setup of the Chair of Materials

Science and Engineering (Fig. 14b) were provided by the Chair of Machine Tools and

Production Engineering of TU-Berlin. The specimens were produced by LPBF-M (SLM

Solutions 250H, MTT Technologies GmbH, Germany), built at 90° to the building

platform, from Ti-5553 powder made by plasma atomization (grain size 15 - 45μm; D50

= 34μm; AP&C Advanced Powders & Coatings, Quebec, Canada). Table 6 displays

the LPBF-M parameter settings and the chemical composition of the powder. The

process parameter combination achieved a relative component density of 99.8% [100].

It is important to mention that the optimization of building parameters towards low

porosity is not within the scope of this study.

Fig. 14: Geometries of the cylindrical specimens a) according to DIN 50113 [146] and

b) customized for the corrosion-fatigue setup of the Chair of Materials Science and

Engineering. All dimensions are in mm.

Table 6 displays the LPBF-M parameter settings and the chemical composition

of the powder. The process parameter combination achieved a relative component

density of 99.8% [100]. The relative component density of the additively manufactured

specimens was determined using the Archimedean principle based on the

displacement method. The corresponding evaluation and a schematic representation

of the experimental procedure and setup of the density measurement according to

Archimedes are available in [100]. It is important to mention that the optimization of

building parameters towards low porosity is not within the scope of this study.

Table 6: LPBF-M parameter settings and the chemical composition (wt.%) of the

Ti-5553 powder.

LPBF-M building parameters

Layer

thickness DS

[μm]

Laser power

[W]

Scanning

speed vS

[mm/s]

Hatch

pattern

Hatch

distance ΔS

[mm]

Focus

position

250

1000

Stripes

0.12

Powder composition (wt.%)

Balance

4.86

4.97

2.80

0.35

Fig. 15 shows typical flaws resulting from the LPBF-M process, such as

insufficient material cohesion and pores. Pores are typically round or irregularly shaped

and are randomly distributed throughout the material. The pore size varies, but they

are generally small (up to 20μm). Insufficient material cohesion or bonding defects

occur at interfaces between layers or adjacent regions within the part. These bonding

defects occur at specific interfaces, with their geometry following the layer boundaries.

Fig. 15: Typical flaws resulting from the LPBF-M process (e.g. STA specimen): a,c,d)

insufficient material cohesion and b,d) pores.

Following LPBF-M, the specimens were heat-treated according to ST, STA, and

BASCA protocols [100] to generate a binary (ɑ+β)-microstructure, while some of them

remained in the “as-built” condition. For the macro-fatigue tests, STA-specimens were

subsequently shot-peened. The pre-established heat treatment and shot-peening

parameters, as well as details on the ductility of the material and the roughness

induced by the shot-peening, can be found in a prior study of the specimen provider

[100].

For nanoindentation tests and microstructure investigations, cross- and

longitudinal sections were extracted from the gauge length of one standard fatigue

specimen (Fig. 14a) of each condition, as shown schematically in Fig. 16, using a

precision cutting diamond saw under water cooling to ensure accuracy and prevent

damage to the specimens. The customized specimens (Fig. 14b) were used for the

fatigue tests on the macro-scale.

Fig. 16: Standard cylindrical fatigue specimen and extraction scheme for transverse

and longitudinal sections used for nanoindentation and microstructure investigations.

4 METHODS

4.1 Pre-loading microstructural investigation

Cross- and longitudinal sections (Fig. 16) of “as-built”, ST, STA, and BASCA

conditions were ground on SiC abrasive paper down to 500 grit and polished with

diamond suspension down to a grain size of 9 µm. For the ST condition, grain size and

phase distribution were first evaluated by light microscopy (DMR, Leica, Germany)

after etching with Kroll’s reagent [147]. The microstructures of the cross-sections of all

conditions were evaluated by HR-SEM (Gemini SEM500 NanoVP, Zeiss, Oberkochen,

Germany) in the backscattered electron (BSE) mode at a voltage of 8-15 kV and a

working distance of 7.4-15.8 mm. Further, longitudinal and cross-sections were

polished for 10 minutes with an active oxide polishing solution (OP-S, Struers,

Denmark) buffered with hydrogen peroxide and ammonia to achieve a smooth surface

fit for evaluation of grain orientation and phase distribution by electron backscatter

diffraction (EBSD).

The EBSD measurements were performed with a field emission scanning

electron microscope DSM 982 GEMINI (ZEISS, Oberkochen, Germany) operated at

15 kV with a step size of 100-400 nm. Image analysis to obtain inverse pole figures

and phase fraction maps, as well as grain size distributions based on area fraction,

was done using the software OIM Analysis 6.0 (EDAX/AMETEK, Mahwah, USA). All

electron microscopy work was carried out at ZELMI of TU Berlin.

The phase composition of one specimen in the “as-built” condition was

measured by X-ray diffraction using a Bruker D8 Diffractometer (Bruker Corporation,

Massachusetts, USA) available at the Chair of Advanced Ceramic Materials, TU Berlin,

equipped with a goniometer radius of 300 mm and a Co anode (Kα = 1.79 Å). The

analysis was performed in the 2θ of 10° to 90° with a step size of 0.02° at 35 kV and

40 mA.

Thin foils were extracted perpendicular to the cross-section of one specimen in

the STA condition using FIB technique on FEI Helios NanoLab 600 (Field Electron and

Ion Company, Hillsboro, USA) and analyzed by TEM (Tecnai G² 20 S-TWIN, FEI

Company, Oregon, USA) at an operating voltage of 300 kV in the bright-field (BF) and

annular bright-field (ABF) modes. The chemical composition was also measured by

energy-dispersive X-ray spectroscopy (EDX) during the TEM investigation.

4.2 Nano-scale experiments

4.2.1 Quasi-static nanoindentation

Quasi-static tests were performed on “as-built”, STA and BASCA conditions using a

Hysitron Triboindenter TI950 (Bruker Corporation, Massachusetts, USA) equipped with

a Berkovich tip. The surface was prepared in the same way as for the EBSD

measurements. 100 indents, equally spaced at a distance of 20μm to avoid mutual

interactions, were placed in a matrix on one cross-section specimen of each condition.

A trapezoid load function (10 s loading → 10 s holding → 10 s unloading) was used

with a maximum force Pmax of 3004 µN. Mechanical properties were extracted from the

resulting load-depth curves according to the Oliver and Pharr method [133].

The morphological investigation of indent characteristics was carried out by

using the scanning probe microscopy (SPM) mode of the nanoindenter to obtain

gradient images and, thus, qualitative and quantitative topographic information on

indent and pile-up sizes of at least 5 nanoindents from each condition. The projected

area of the indent (Ap), the outer pile-up area (Ap-u), and the pile-up volume (Vp-u) were

determined after background subtraction using Fiji [148,149], where the best ﬁt for Ap

and Ap-u was found by visual comparison (see typical example in Fig. 17). Vp-u was

determined by the integration of all pixels under the pile-up area. Indent and pile-up

proﬁles were measured using Gwyddion [150], and maximum pile-up height (hp-u,max)

was obtained from the indent and pile-up profiles (Fig. 18).

Fig. 17: Visual inspection of indent morphology using Fiji: a) SPM topography image

of a typical indent; typical selection of b) projected area and c) pile-up area.

Fig. 18 [151]: Indent and pile-up size evaluation: SPM topography images of a typical

indent with profile lines a) through the highest point of the pile-up and oriented parallel

to the edges of the indent; b) through the corners and the tip of the indent, cutting the

opposite indent edges at 90°. Typical topography charts: c) pile-up, measured along

the profile lines in a); d) pile-up and indent shape, measured along the profile lines in

b). Note that additional profile lines (“1”) were placed horizontally, outside the area

affected by the indent and the pile-up, representing the zero-level for the height

calculations.

4.2.2 Cyclic nanoindentation

Nanofatigue experiments were performed on the same nanoindenter used for

the quasi-static tests (section 4.2.1). The surface of a cross-section through a LPBF-

M specimen of each condition was prepared in the same way as described in section

4.1 for the metallographic sections used for the EBSD measurements.

24 indents were placed in a square map, spaced equally at a distance of 30μm

to avoid mutual interactions, as for the quasi-static tests. Four additional indent

positions were selected on the ST condition by visual inspection of etched surfaces;

the regions of interest were distinguished by local variations in ɑ-phase occurrence:

one in which the α-precipitates seem to be more elongated and mostly oriented in a

preferential direction, and another in which the α-precipitates seem smaller with

random orientation (see Fig. 28).

For nanofatigue loading, each site was indented cyclically to a maximum cycle

number of 105 at a frequency, f, of 201 Hz. Pmin of 255 ± 5 μN, and Pmax of

2893 ± 4 μN, corresponding to a Pa of 1574 ± 2 μN and Pm of 1319 ± 4 μN, were set

for the tests (indents 1 to 24). The four positions placed in defined regions of interest

on the ST condition surface (addressed in the following as indents A, B, C, D) were

loaded with Pmin = 456 µN and Pmax = 2946 µN (Pa = 1245 μN, Pm = 1701 µN). The

minimum load ensured constant contact between the tip and the sample surface

throughout the tests in all cases.

Because of limitations in the data acquisition rate and the amount of data that

can be stored, the number of data points available is not high enough in the high-

frequency loading cycles to evaluate the force/indentation depth curves. Therefore,

low-frequency, so-called “measurement cycles” at f = 0.05 Hz (Fig. 19a) were inserted

at regular intervals to yield force/indentation depth hysteresis loops with sufficient data

points (Fig. 19b). Further, for indents 1 to 24, static “holding” segments at 10% of the

maximum load were introduced before and after the measurement cycles for thermal

drift correction (for more details, see [152]). The force/indentation depth-hysteresis

loops were evaluated according to protocols used in classical fatigue testing by

customized code based on Python [153]. The cyclic deformation behavior is

characterized by the development of the plastic indentation depth amplitude, Da,p,

determined as the half-width of the hysteresis loop at mean force. The development of

the ratios of Dmin to Dmax over the cycle number, N, and the change of plastic

deformation between cycles, ΔDmin, induced by repeated application of the force

amplitude over N, and represented by the change in minimum depth (Dmin) between

consecutive measurement cycles, was analyzed to evaluate the cyclic creep behavior.

Fig. 19 [151]: a) Load function used for the nanofatigue tests consisting of alternating

blocks of high-frequency loading at f = 201 Hz (marked in red, “loading cycles”) and of

low-frequency loading at f = 0.05 Hz (marked in black, “measurement cycles''). Note

that the force amplitude and the mean force were the same for both blocks; for clarity,

the loading cycles are only depicted as red lines at mean force, with an indication of

the loading times. b) Schematic representation of a typical force/indentation depth

hysteresis loop from which the maximum depth (Dmax), the minimum depth (Dmin), and

the plastic indentation depth amplitude (Da,p) are determined.

For the ST condition, the four nanoindents “A” to “D”, placed in defined regions

of interest, were imaged in the same HR-SEM described in section 4.1., also in the

BSE mode at a voltage of 8 kV, at a working distance of 7.6 mm. Microstructure,

dislocation structure, and density in the volume directly beneath these indents were

investigated using the same TEM and FIB machines mentioned in section 4.1. The

TEM was used at an operating voltage of 200 kV in the BF mode. The FIB technique

was used to prepare thin foils that were analyzed by TEM. One long edge of the foil

was oriented parallel to the indentation direction, and the section was placed as

precisely as possible through the tip of each nanoindent. To protect the foil against

plastic deformation and the sample surface from the gallium ions, the sample was

covered with a thin platinum layer before ion beam milling. Although it cannot be

assured that the TEM foils are placed right through the tip, it is certain that they pass

very close to it, given that the maximum indentation depth after fatigue loading

observed from the triangular profile in the TEM micrographs (compare Fig. 28)

matches the range of depths of the four indents “A” to “D” (182 to 261 nm). All SEM,

FIB, and TEM work was performed at ZELMI of TU Berlin.

The qualitative and quantitative evaluation of indent and pile-up morphology and

size was also made for cyclic nanoindents, as described for the quasi-static

nanoindents in section 4.2.1.

4.3 Macro-scale experiments

Load-controlled axial stepwise load increase (LIT) and constant-amplitude

(CAT) corrosion-fatigue tests were performed at f = 5 Hz with a load ratio of R = -1 on

4 and 5 specimens, respectively, in Hanks' balanced salt solution in distilled water

(HBSS, Table 7) at 37 °C using a Schenck servohydraulic testing system (PSA 100,

Schenck-Instron, Darmstadt, Germany) equipped with a 100 kN load cell. Fig. 14b

presents the geometry and size of the tested specimens. The pH value was set to 7.4

at the start of the tests. Cyclic loading was started after passivating the specimen in

HBSS in the load-free state for about 2 h until both the Ucorr and icorr curves exhibited

minimal fluctuations (at a nanoscale range) over time. Precise alignment of the

hydraulic grips before the tests ensured that the specimens experienced no bending

moments. For the LIT, the initial stress amplitude of 25 MPa was increased by 10 MPa

every 20,000 cycles over 1,000 cycles. To test a large number of cycles, particularly in

the HCF region, one σa of 210 MPa was chosen for the CAT based on the endurance

limit estimated by the LIT.

Table 7 : Required components to prepare of HBSS [154].

Component

Molar weight

Amount

Concentration

(g/mol)

(mg/L)

(M/L)

NaCl

58.44

8000

0.1400

KCl

74.55

400

0.0050

CaCl2

110.98

140

0.0010

MgSO4-7H2O

246.47

100

0.0004

MgCl2-6H2O

203.30

100

0.0005

Na2HPO4-2H2O

177.99

0.0003

KH2PO4

136.09

0.0004

D-Glucose (Dextrose)

180.16

0.0060

NaHCO3

84.01

350

0.0040

The cyclic deformation behavior was evaluated by the evolution of the plastic

strain amplitude (εa,p) versus N. The strain was measured by an inductive strain gauge

with a resolution of 0.05 µm (eddyNCDT 3300, Micro-Epsilon Messtechnik GmbH &

Co. KG, Ortenburg, Germany) between the regions of the heads of the specimen

directly adjacent to the radius (see Fig. 20). Although this leads to less precise strain

measurements relative to measurements in the gauge length, this approach prevented

any potential damage to the surface exposed to HBSS which might be caused by the

clamping of an extensometer.

Fatigue-induced surface modifications were identified using an electrochemical

open circuit setup allowing the simultaneous detection of the free corrosion potential

and current, schematically shown in Fig. 20. Ucorr was measured between the WE

(specimen) and RE (Argenthal Ag/AgCl, +207 mV to normal hydrogen electrode), while

icorr was determined by connecting the WE with a CE via a low noise multimeter. The

CE was a specimen consisting of the same material with the same surface condition

as the WE. One CAT was interrupted when a significant increase in icorr and a decrease

in Ucorr indicated changes in the surface state. The surface modifications were

examined by SEM as described above (section 2.2) to correlate the damage with the

electrochemical measurements. Fracture surfaces of selected specimens were also

investigated by SEM.

Fig. 20: Schematic drawing of the experimental setup for the corrosion-fatigue tests.

5 RESULTS

5.1 Microstructure

The “as-built” material exhibits a single-phase microstructure consisting entirely

of β-phase (Fig. 21a). The β-grains are recognized due to their different gray values.

For the STA condition (Fig. 21b-d), primary globular α-phase precipitates (αp)

are observed inside β-grains (Fig. 21b,c). These αp-precipitates in the grains are

significantly smaller than other longitudinal ones, arranged like a “rope of pearls” at the

β-grain boundaries and surrounded by approximately 2 to 4 µm wide, αp-free regions

along the boundaries. Further, a large number of very fine secondary α-phase

precipitates (αs) with a needle-like shape and a length of up to only 250 nm are seen

(Fig. 21d). After BASCA heat treatment (Fig. 21e,f), αp-precipitates at the β-grain

boundaries, and αs-laths divided into colonies with different orientation are observed.

Both αp-precipitates and αs-laths differ in size and thickness, and they are much thicker

than the precipitates seen in STA. The αp-precipitates are arranged along the β-grains

boundaries, often forming continuous bands (Fig. 21f). The length of the αs-laths is in

the micrometer range.

Digital image analysis of EBSD data gave phase fractions of β and α for STA

and BASCA conditions (Fig 22a,b). The STA condition exhibits 81% of β-phase and

19% of α-phase (Fig. 22a), whereas 37% of β-phase and 63% of α-phase are found

for BASCA (Fig. 22b). For the “as-built” condition, the presence of only β-phase was

confirmed by the X-ray diffraction (XRD) measurement (Fig. 23).

Fig. 21: HR-SEM images of the microstructure of Ti-5553 (cross-sections): a) “as-

built”: single-phase microstructure. b-d) STA: globular αp-precipitates within a

transformed β-matrix (b,c) with fine, very small αs (d). e,f) BASCA: αp-precipitates at

the β-grain boundaries (e), and αs-laths divided into colonies inside the β-grains (f),

both differing in size and thickness.

Fig. 22: Phase fractions of β and α phases obtained from digital image analysis of

EBSD data of the heat-treated conditions: a) STA with 0.81% of β and 0.19 of α, and

b) BASCA with 0.37% of β and 0.63% of α.

Fig. 23: Phase composition of the as-built condition, measured by XRD.

EBSD inverse pole figures (IPF) (Fig. 24) reveal that the differences in gray

values observed in Fig. 21a are indeed due to different crystal orientations. The “as-

built” condition (Fig. 24a) shows β-grains with a square shape in the cross-section,

orthogonal to the LPBF-M build direction. A crystallographic texture in the (001)

direction is noticeable for the STA condition (Fig. 24c). While the shape of the β-grains

in the heat-treated variants remains similar to that of the "as-built" condition, the grain

size distributions based on area fraction measurements from EBSD quality images (QI)

of cross-sections (Fig. 25) show a general shift towards increasing grain sizes from

the “as-built” to the STA, clearly indicating the grain growth. Conversely, the BASCA

condition shows smaller grain sizes, indicating a process of grain refinement. The

dominant grain size diameter in the "as-built" condition is 64 µm, while 79 µm and

28 µm are observed for STA and BASCA, respectively. In the longitudinal sections

(Fig. 24b,e,h), the β-grains exhibit an elongated morphology, parallel to the AM build-

direction. The diameter of these grains is up to 100 µm micrometers. The images in

Fig. 24d,g validate the observation in Fig. 21. α-precipitates are seen inside and along

the boundaries of the β-grains. The STA-microstructure presents a β-phase matrix with

a smaller α-phase (Fig. 24d), whereas, in the microstructure of BASCA, a higher

amount of α-phase is observed, which is also much bigger (Fig. 24g). In both

conditions, the α-precipitates are distributed within the β-grains in portions (colonies)

with preferred orientations.

Fig. 24: a-h) EBSD inverse pole figures (IPF). β-Ti measurements on a,c,f) cross-

sections prepared orthogonal and b,e,h) longitudinal sections prepared parallel to the

LPBF-M build direction. d,g) α-Ti measurements on cross-sections of d) STA, g)

BASCA.

Fig. 25: Grain size distributions based on area fraction from EBSD quality images (row

below) of cross-sections.

ABF and BF-TEM micrographs in Fig. 26 show the STA-microstructure in more

detail. A large number of αs-precipitates are seen and they are closely packed and

oriented perpendicular to each other. The αs-precipitates are significantly smaller than

the αp, exhibiting length up to 250 nm and thickness up to 32 nm. Fig. 27 displays the

EDX chemical composition mapping of the specimen. The elemental mapping was

obtained by two-dimensionally scanning an electron beam over an area where the

elements exist (specimen surface), and the X-ray spectra generated from each pixel

were acquired. The spectrums measured from the pixels on the mapping region

resulted in strong characteristic X-ray peaks of the detected elements. The integrated

intensities of the characteristic X-rays of the elements or the concentrations of the

elements were plotted pixel by pixel to provide the elemental map. When the intensity

or concentration is low, the pixels are depicted black on the elemental map, while high

intensity or concentration is shown by the brighter, colored pixels.

Fig. 26: a) ABF-TEM micrograph revealing the β-matrix with a large number of very

fine αs-precipitates (dark) with a needle-like shape, significantly smaller than the αp

(bright). b) BF-TEM micrograph showing the αs-precipitates in more detail.

Fig. 27: EDX mapping of atom distribution together with the weight amounts of the

STA-specimen.

The ST condition has a biphasic microstructure with small αp-precipitates highly

dispersed within the grains and at the grain boundaries of the retained β-matrix (Fig.

28 and 29e,f). In low-magnification light micrographs (Fig. 28a-d), the grain boundaries

appear as bright stripes, some of which contain black lines. Higher magnifications (Fig.

28e,f) reveal that these bright stripes consist of β-phase and that the black lines are

αp-particles arranged like a rope of pearls approximately in the center of the β-stripe.

Digital image analysis of EBSD data gave phase fractions of about 88% and 12% for

β and αp, respectively, for both transverse and longitudinal sections (Fig. 29a,b). The

β-grains have square cross-sections orthogonal to the LPBF-M build direction, and a

slightly elongated shape in the longitudinal orientation, parallel to the build direction.

EBSD inverse pole figures (IPF) (Fig. 29c,d) show no crystallographic texture.

By light microscopy, lighter and darker regions are visible in both cross- and

longitudinal sections of the ST condition. Each of these regions comprises several

grains (Fig. 28a,b). EBSD-IPF views (Fig. 29e) and higher magnified SEM

micrographs (Fig. 29f) reveal that the differences in gray value are due to different

densities, shapes and amounts of sections of αp-precipitates in the field of view, as

shown exemplarily in Fig. 30. The different views reveal that the precipitates are

acicular, with their diameters varying along their length. Hence, the differences in the

appearance of αp come from their different orientation and, thus, the appearance of the

exposed sectioned plane. Even though the surface comprises regions where the αp-

precipitates exhibit a preferential orientation, EBSD measurements revealed no overall

preferred crystal orientation and no orientation relationship between the α-phase and

the surrounding β-matrix.

Fig. 28 [151]: a-f) Light micrographs of polished and etched cross- (a,c,e) and

longitudinal (b,d,f) sections reveal a biphasic microstructure of αp (dark gray) finely

distributed in retained β-matrix (bright gray). Two regions with different gray values are

observed: a lighter one in which the α-precipitates seem to be more elongated and

mostly oriented in a preferential direction (a-d), and a darker one in which the α-

precipitates seem smaller with random orientation (e,f). At smaller magnification (a,b),

the grain boundaries appear as bright stripes, as if consisting only of β-matrix. Higher

magnifications reveal darker areas within these stripes, which are α-precipitates

arranged like a rope of pearls.

Fig. 29 [151]: a,b) EBSD phase distribution measurements on transverse (a) and

longitudinal (b) sections confirming the existence of a biphasic microstructure with a

phase fraction of about 88% β-phase and 12% α-phase. c-d) EBSD-IPF results of

cross- (c) and longitudinal (d) sections showing the structure of the shape and

orientation of the β-grains. The grains are elongated in the LPBF-M build direction with

a clearly preferred orientation, yet without crystallographic texture. e-f) Higher

magnified EBSD-IPF (e) and SEM (f) images showing the α-precipitates differing in

orientation and sectioning to the surface.

Fig. 30 [151]: Schematic representation of the crystallographic orientation of αp based

on the EBSD-IPF maps: differences in the appearance of αp come from their different

orientations and therefore different sectioning to the surface.

5.2 Nano-scale investigation: ST condition

5.2.1 Cyclic deformation and creep behavior

The cyclic deformation behavior of the ST condition is displayed in Fig. 31 by

plots of the plastic displacement amplitude Da,p versus N. On a macroscopic view,

stages of hardening, saturation, and softening are observed. However, within each of

these three stages, fluctuations of Da,p occur. The indents differ by the extent of the

stages relative to each other, and by the size of the fluctuations. The first stage

comprises the first ten cycles. Here, most nanoindents present overall hardening,

visible by a net decrease in Da,p. This stage is characterized by moderate alterations

between hardening and softening for more than 50% of these indents, by soft

alterations for a quarter of the indents, and by strong alterations for about a fifth.

Indents 3 and 16 (Fig. 31b) are examples of moderate and soft alterations,

respectively. Only about 3% of the nanoindents, for instance indent 1, exhibit softening

following pronounced hardening in the first cycle, and subsequent strong fluctuations.

For most nanoindents, the initial hardening or softening is followed by weak further

softening for loading to up to 104 cycles: more than half of the indents (54%) show an

increase in Da,p with pronounced fluctuations, 18% with small fluctuations and 7%

without fluctuations (compare indents 1 and 16 in Fig. 31b). Only about a fifth of the

nanoindents, e.g. indent 3 (Fig. 31b), exhibit a low amount of further hardening: 18%

show a decrease in Da,p with fluctuations, and 3% without fluctuations. During further

loading, to the maximum number of cycles of 105, the curves of most indents (75%)

decrease with fluctuations, 22% increase with fluctuations, and 3% of them present a

pronounced increase with fluctuations.

Fig. 31 [151]: Cyclic deformation curves, Da,p over N, for a) all indents and b) typical

examples, representing the different progression types (nanoindents 1, 3, and 16).

The progressions of the ratios of Dmin to Dmax and of ΔDmin over the number of

cycles were analyzed to further evaluate the amount of plastic deformation per cycle

and the cyclic creep behavior, respectively. The curves for Dmin/Dmax over N (Fig. 32)

progress in similar ways, with an overall increase in Dmin/Dmax, but they differ in the

extent of the increase: some curves progress to higher and others to lower values.

Fig. 32 [151]: a) Dmin/Dmax over N for all indents. b) Examples of indents that differ in

their extent of curve progression.

ΔDmin represents the incremental, non-reversible deformation, induced by

repeated loading, over the entire loading history. As the number of cycles between the

“measurement cycles” is not constant, and because small differences in the minimum

load are not avoidable, ΔDmin was normalized by the number of cycles over which the

change occurred and by the minimum load, which gives us ΔDmin-norm (Fig. 33). An

overall decrease in ΔDmin-norm over N is seen for all indents (Fig. 33a-c), however, with

high fluctuations from cycle to cycle at the beginning of loading, up to N = 10. Some

indents even present negative ΔDmin-norm values. Up to the 100th cycle, ΔDmin-norm

approaches values between 10-4 and 3x10-4 nm/μN. Over the further course of loading

(104 ≤ N ≤ 105), an overall decrease in ΔDmin-norm with slight alterations below 2x10-6

nm/μN is observed. Note, however, that the ΔDmin-norm values are averaged over

increasing numbers of cycles with increasing N, which is expected to add to a smoother

progression. The three typical progressions (Fig. 33d) highlight the pronounced

differences in the extent of the fluctuations seen over the first 10 cycles. Further, for

each of the indents, ΔDmin-norm approaches distinguishable levels over the further

course of loading (Fig. 33e,f).

Fig. 33 [151]: Development of ΔDmin-norm over a,d) 1 ≤ N ≤ 103, b,e) 102 ≤ N ≤ 104 and

c, f) 104 ≤ N ≤ 105. a-c) Curves for all nanoindents, and d-f) typical examples

(nanoindents 1, 3, and 16) of the different behaviors observed.

5.2.2 Cyclic nanoindent morphology

Nanoindent and pile-up size and morphology were evaluated quantitatively

based on the SPM images. Maximum indentation depth after fatigue loading (Dmax at

the maximum number of cycles, N = 105), projected indent area (Ap), pile-up area

(Ap-u), pile-up volume (Vp-u), and maximum pile-up height (hp-u,max) are shown for all

nanoindents in Table 8.

Table 8 [151]: Results of the quantitative analysis of indent and pile-up size: projected

indent area (Ap), pile-up area (Ap-u), pile-up volume (Vp-u), and maximum pile-up height

(hp-u,max), together with the maximum indentation depth after fatigue loading (Dmax(N =

105)) (n = indent number). The smallest and greatest values for each parameter are

underlined and printed in bold, respectively.

Indent

Pile-up

Dmax (N=105)

Ap-u

hp-u,max

Vp-u

(nm)

(10-3 µm2)

(µm)

(10-4 µm3)

272.72

4.77

11.25

0.19

1.91

231.28

3.81

6.81

0.15

1.82

223.43

3.76

7.11

0.01

2.84

222.48

4.18

6.79

0.11

1.47

217.56

3.39

6.51

0.12

0.36

213.44

3.61

7.47

0.14

0.88

214.1

3.84

10.77

0.10

1.48

205.19

3.57

10.62

0.11

1.25

204.76

2.96

2.52

0.13

0.36

196.16

2.91

5.27

0.07

0.77

208.03

3.14

9.85

0.08

0.83

197.49

2.67

4.07

0.15

0.60

209.54

3.91

9.68

0.06

1.12

213.55

3.73

7.41

0.1

0.86

211.97

3.12

5.28

0.12

0.41

198.81

3.01

4.12

0.09

0.54

195.63

3.19

6.48

0.09

0.31

200.66

3.38

6.14

0.13

0.92

194.13

2.49

3.10

0.06

0.30

220.36

3.74

6.53

0.08

0.70

212.55

4.19

7.14

0.13

0.96

217.21

3.14

3.98

0.11

0.63

216.72

3.82

6.84

0.12

0.89

210.74

3.08

5.57

0.13

0.86

The linear correlation coefficients (R*) for the different parameters are

summarized in Table 9. The strongest correlations are seen between Dmax (N = 105)

and Ap (R* = 0.78), and between Ap and Ap-u (R* = 0.71), whereas the weakest

correlations were found for Ap-u and hp-u,max (R* = 0.11), and for hp-u,max and

Vp-u (R* = 0.29).

Table 9 [151]: Linear correlation coefficients (R) between the indent and pile-up

parameters given in Table 2. The two weakest values of R are underlined, and the

strongest values are given in bold font.

Parameter 1

Parameter 2

Dmax

0.78

Ap-u

0.71

hp-u,max

0.37

Vp-u

0.65

Ap-u

Dmax

0.49

hp-u,max

0.11

Vp-u

0.54

hp-u,max

Dmax

0.60

Vp-u

0.29

Vp-u

Dmax

0.61

Fig. 34 displays SPM topography images and 3D surface plots of typical

nanoindents with different cyclic deformation behaviors, as described in section 3.2.

Indent 1 (Fig. 34a) is located in a surface region comprising αp-precipitates oriented

parallel to each other. It exhibits the largest indent and pile-up sizes, regarding area

(Ap, Ap-u) and maximum height (hp-u,max), but an intermediate pile-up volume (Vp-u). In

the surface region surrounding indent 3, the αp-precipitates are transversely located

(Fig. 34b). This indent is about 21% smaller than indent 1 (Ap). Its lower deformation

compared to indent 1 is also represented by smaller values of Ap-u and hp-u,max.

However, the indent 3 has the highest Vp-u. Indent 16 (Fig. 34c) is even smaller (Ap)

with very little pile-up, as measured by Vp-u, although the surface microstructure

surrounding indent 16 is very similar to that of indent 1.

Indents “A” to “D” were placed in selected areas of the specimen surface at

positions with differences in the local distribution and surface appearance of the ɑ-

phase (see HR-SEM insets in Fig. 35). Indent A is located in a surface region

consisting of β-phase only. It has a large size and pile-up, mainly on its left edge.

Indents B to D were placed in regions with αp-precipitates visible on the surface. While

indent B was made in a region containing precipitates with a preferred orientation,

indents C and D sit in surface regions with a relatively high content of αp-precipitates,

oriented at different angles to each other. Both indents C and D hit αp-precipitates,

while indent B sits in between the precipitates. Intermediate nanoindent and pile-up

sizes are observed for indent B, and both values are smaller than for indent A. For

indent C, the precipitates appear to be agglomerated, with a small or no distance

between them. This nanoindent has an indent area similar to indent A, however with a

smaller pile-up area. Indent D exhibits the smallest indent area among the four indents,

without significant pile-up.

Fig. 34 [151]: 3D surface plots and SPM topography images (insets in the upper right

corners) of cyclic nanoindents with different, typical morphologies: a) indent 1, with a

large projected area, large pile-up height, and high indentation depth; b) indent 3, with

intermediate projected area, pile-up height, and indentation depth values; c) indent 16,

with a small pile-up height, projected area, and indentation depth. Note that for better

visibility of the nanoindent in the 3D surface plots, a different scale of the z-axis (depth

of the indent) as compared to the lateral scale was chosen.

5.2.3 TEM investigation

TEM (Fig. 35) was used to evaluate the microstructure and the dislocation

density in volumes below the indents “A” to “D”. Fig. 35a shows β-phase regions region

in the volume beneath indent A, and only one αp-precipitate is visible in the plane of

the TEM foil. High dislocation density is observed along the grain boundaries. Further,

twins are seen near the indent impression. For indent B (Fig. 35b), a region with a high

content of αp-precipitates is observed, and high dislocation density is seen mainly

along grain or phase boundaries. A high number of αp precipitates with different

orientations is also revealed in the volume directly beneath indent C (Fig. 35c). In this

case, a higher dislocation density appears to be trapped between adjacent,

transversely positioned α-phase precipitates. In the volume beneath indent D (Fig.

35d), three αp-precipitates are oriented parallel to each other, at an acute angle of

about 135° to the surface. Around these particles, a high dislocation density is

observed.

Fig. 35 [151]: a-d) Bright-field TEM micrographs of the volumes beneath the indents

“A” to “D” together with HR-SEM images of the tested surface (insets, lower left corner):

a) Indent A was made in a surface region without α-precipitates, and only one α-

precipitate is visible below the indent in the plane of the TEM foil; high dislocation

densities exist along the grain boundaries, and twinning (red arrow) is also observed.

b) Indent B sits in a region with a high content of α-precipitates, both on the surface

and in the volume surrounding the indent in the plane of the TEM foil. Regions of high

dislocation density are seen, mainly along grain or phase boundaries. c) Indent C was

also placed in a surface and volume region with a high content of α-precipitates. They,

however, exhibit different orientations, and the distances between the α-particles are

smaller than for indent B. A high dislocation density is observed between two adjacent

precipitates, oriented nearly transversely to the indentation direction. d) Indent D was

performed in a region with a lower α-phase content than seen for indent “C”. The

precipitates are oriented parallel to each other, and one precipitate is included in the

indented area (inset). Regions of higher dislocation density are observed surrounding

the precipitates.

5.3 Nano-scale investigation: comparison between “as-built”, STA and

BASCA conditions

5.3.1 Quasi-static nanomechanical properties

Fig. 36a shows the mean load-depth curves and their scatter bands together

with SPM images of typical indents for the “as-built”, STA, and BASCA conditions. The

scatter bands are the minimum and maximum depth for each load value (i.e. from

different indents). The load-depth curves for the three conditions exhibit distinct

characteristics. The “as-built” condition displays a curve that is notably shifted to the

right, indicating a lower resistance to indentation. In contrast, the BASCA condition

exhibits a curve positioned in the middle, reflecting intermediate plastic deformation.

Finally, the curve of the STA condition is placed to the left, suggesting higher resistance

to indentation. It becomes clear that the “as-built” specimens exhibit the highest

plasticity, followed by BASCA and STA. The values of reduced modulus (Er), the

nanohardness (H), and stiffness (S) are displayed in Fig. 36b, Fig. 36c, and Fig. 36d,

respectively. In the BASCA condition, the median values for Er and S are the highest,

measuring 166 GPa and 134 µN/nm, respectively. However, the highest median for H

is observed in the STA condition (7 GPa). Conversely, the "as-built" condition shows

the lowest median values for Er, H, and S. The BASCA condition exhibits more

variability in Er and H, as indicated by a larger interquartile range (IQR) and extended

whiskers, indicating a wider range of data beyond the IQR. This is followed by the STA

and "as-built" conditions. Interestingly, the whiskers for S remain relatively consistent

across all three conditions (Fig. 36c). Outliers, which are individual data points

significantly deviating from the majority, are represented beyond the whiskers.

Specifically, the BASCA condition shows more outliers for Er, while both BASCA and

"as-built" conditions have more outliers for H.

Fig. 36: Results from quasi-static nanoindentation tests on the “as-built”, STA, and

BASCA conditions: a) mean load-depth curves with their scatter bands (minimum to

maximum values) and exemplary gradient SPM images of indents; b) Box plots

illustrating the mechanical properties, with the box representing the interquartile range

(IQR) containing the middle 50% of the data. The line inside the box indicates the

median. 'Whiskers' extend from the box to show the data range beyond the IQR, while

outliers, significantly deviating data points, are represented as individual points beyond

the 'whiskers'.

Fig. 37 shows 3D surface plots and SPM topography images (upper right corner

insets) of typical indents for each condition. The indent of the “as-built” condition is

much bigger than the indent of the STA condition, which is smaller than the indent of

the BASCA condition. This difference is reflected by the values of Ap. The “as-built”

indent also exhibits the greatest maximum indentation depth (Dmax) (Fig. 37a). The

smallest values of Ap, Ap-u, and Vp-u are found for the STA specimen (Fig. 37b), while

the indent of the BASCA condition exhibits the greatest Ap-u and Vp-u, with an

intermediate Ap (Fig. 37c).

Fig. 37: 3D surface plots and SPM topography images (upper right corner insets) of

examples of quasi-static indents for each condition: a) “as-built”, b) STA, and c)

BASCA. (Note the different scale of the z-axis (depth of the indent) as compared to the

lateral scale, chosen for better visibility of the indent in the 3D surface plots).

5.3.2 Cyclic deformation behavior

The progressions of the mean ratios of Dmin to Dmax and of ΔDmin, together with

their respective scatter bands, over the number of cycles, were analyzed to evaluate

the plastic deformation behavior. The Dmin/Dmax over N curves for all three conditions

exhibit very similar progressions with a nearly steady increase, and they differ only in

the extent of the increase (Fig. 38). The curve of BASCA exhibits higher values, while

the curve of the “as-built” and STA conditions increases at intermediary and at lower

values, respectively.

Fig. 38: Mean Dmin/Dmax over N curves together with their respective scatter bands for

the “as-built”, STA, and BASCA conditions.

Fig. 39 presents the mean ΔDmin-norm over N curves for the “as-built” and heat-

treated conditions. Because the number of cycles between the “measurement cycles”

is not constant, and because small differences in the minimum load are not avoidable,

ΔDmin was normalized by the number of cycles over which the change occurred and by

the minimum load, which gives us ΔDmin-norm. All conditions exhibit a decrease in ΔDmin-

norm over the course of loading, with no significant differences between the curves. For

cycle numbers above 104, the curve of the “as-built” condition progresses at slightly

higher values as compared to those of STA and BASCA.

Fig. 39: Mean ΔDmin-norm over a) 1 ≤ N ≤ 103, b) 102 ≤ N ≤ 104 and c) 104 ≤ N ≤ 105

curves together with their respective scatter bands for the “as-built” and heat-treated

conditions.

Dap versus N representing the cyclic deformation behavior is displayed in Fig.

40. For N ≤ 10 (Fig. 40a), all three conditions present a decrease in Dap. While very

similar values are observed for the heat-treated conditions, the curve for the “as-built”

condition exhibits lower values. The values of all three conditions increase for 10 ≤ N

≤ 102, with both heat-treated conditions increasing at the same extend, while for the

“as-built” condition the values progress at lower values. Cyclic hardening takes place

between 102 and 103 cycles for STA and BASCA conditions with very similar values

(Fig. 40b), and saturation is seen for the “as-built” condition. For 103 ≤ N ≤ 104 (Fig.

40c), a cyclic softening is observed for “as-built”, whereas STA and BASCA conditions

experience ongoing cyclic hardening. Subsequently, after 104 cycles (Fig. 40d), the

“as-built” condition exhibits cyclic hardening, and its curve progresses at visible higher

values as compared to the heat-treated states. In the heat-treated conditions,

saturation is observed, with the curve for BASCA progressing at slightly lower values

than STA.

Fig. 40: Mean curves with their respective scatter bands for Da,p versus N of “as-built”

and heat-treated conditions: a) 1 ≤ N ≤ 10, b) 10 ≤ N ≤ 103, c) 103 ≤ N ≤ 104, and d) 104

≤ N ≤ 105.

The analysis of the individual behavior of the indents (Fig. 41) reveals that

macroscopic stages of hardening, saturation, and softening are evident across all three

conditions. Fluctuations in Da,p are observed within each of these stages, with curves

differing from one another in terms of their courses for a given condition. Note that their

trajectories within the scatter band do not consistently follow the same pattern in terms

of hardening, softening, and saturation. Within a condition, an indent may exhibit a

softening stage, while another presents hardening or saturation stages. A greater

divergence in the curves and more pronounced fluctuations are observed in the “as-

built” condition, followed by the STA condition, with the BASCA condition showing

curves that deviate less from each other in terms of their courses.

Fig. 41: Cyclic deformation curves, Da,p over N, for all indents of a) “as-build”, b) STA

and c) BASCA conditions.

SPM images of selected cyclic indents were analyzed quantitatively in order to

evaluate their pile-up size and morphology. Ap, Ap-u, Vp-u, and hp-u,max, together with

Dmax at the maximum number of cycles (N = 105) and the arithmetic mean roughness

(Ra) of the surrounding surface are shown in Table 10. The indents made on the “as-

built” specimen exhibit the highest values for all parameters, followed by BASCA and

STA. Indents on BASCA surfaces are about 31% smaller and exhibit about 47%

smaller Vp-u than the indents made on the “as-built” specimen surface. The indents

made on the STA surface are even about 40% smaller and have the lowest values for

Vp-u (61% smaller).

Table 10: Quantitative evaluation of pile-up size and morphology of selected cyclic

indents from SPM images of each condition: projected indent area (Ap), pile-up area

(Ap-u), pile-up volume (Vp-u), maximum pile-up height (hp-u,max), maximum indentation

depth after fatigue loading (Dmax (N=105)), and arithmetic mean roughness (Ra) of the

surrounding surface.

Condition

Indent

Surface

Indent

Pile-up

Dmax (N=105)

Ap-u

hp-u,max

Vp-u

(µm)

(10-3 µm2)

(µm)

(10-3 µm2)

(µm)

(10-5 µm3)

as-built

0.171 ± 0.002

9.108

0,245

13.716

0.227

29.837

0.160 ± 0.001

7.002

0,236

12.896

0.180

19.491

0.155 ± 0.001

7.648

0,237

9.006

0.178

13.735

0.151 ± 0.001

7.333

0,204

7.922

0.173

11.491

0.157 ± 0.001

7.176

0,206

13.243

0.181

17.499

STA

0.132 ± 0.001

5.125

0,192

5.433

0.168

10.345

0.100 ± 0.002

4.484

0,165

5.311

0.121

7.927

0.090 ± 0.001

4.840

0,178

3.411

0.124

5.738

0.101 ± 0.001

4.740

0,177

4.364

0.119

6.567

0.106 ± 0.001

3.925

0,164

3.810

0.138

5.078

BASCA

0.141 ± 0.004

5.085

0,169

4.876

0.168

9.270

0.138 ± 0.003

5.952

0,181

4.742

0.200

12.460

0.139 ± 0.003

4.824

0,171

4.873

0.162

7.382

0.119 ± 0.004

4.991

0,179

4.250

0.164

8.553

0.113 ± 0.001

5.370

0,176

5.874

0.168

11.138

Fig. 42 displays SPM topography images (insets in the upper right corner) and

3D surface plots of typical cyclic nanoindents from each condition. The indent of the

“as-built” specimen (indent 15, Fig. 42a) exhibits the highest values in terms of Ap, Ap-

u, and Vp-u. This indent also presents the highest indentation depth after fatigue loading.

The indents of STA and BASCA specimens, indent 16 (Fig. 42b) and indent 10 (Fig.

42c), respectively, were performed on surfaces consisting of β-phase with α-phase

precipitates. Although both indents show similar Ap-u values, the morphology and

content of α-phase on their surrounding microstructures are quite different. While small

α-phase precipitates are seen for STA, the indent on the BASCA specimen was made

on a surface portion with a high number of precipitates, which are also much bigger.

The indent on the STA specimen also has roughly smaller values of Ap-u and Vp-u. in

comparison with those for BASCA.

Fig. 42: 3D surface plots and SPM topography images (upper right corner insets) of

examples of cyclic indents for each condition: a) “as-built” (indent 15), b) STA (indent

16), and c) BASCA (indent 10). (Note the different scale of the z-axis (depth of the

indent) as compared to the lateral scale, chosen for better visibility of the indent in the

3D surface plots).

5.4 Macro-scale investigation: STA condition

5.4.1 Estimation of the endurance limit in load increase tests

Fig. 43 shows εa,p, and σa mean value curves versus N with their scatter bands

(minimum to maximum values) from LIT. The portion where the slope of εa,p exhibits a

significant change (marked with an arrow) indicates an estimation value for the

endurance limit of about 165 MPa. Variability is observed in both the εa,p curves and

the endurance limits among the specimens.

Fig. 43: Progression of mean εa,p versus N with scatter band showing the minimum

and maximum values, together with typical σa versus N curve from load increase tests

and indication of endurance limit estimation.

5.4.2 Cyclic deformation behavior in simulated physiological media

The cyclic deformation behavior, measured in CAT at σa of 210 MPa, is

displayed in Fig. 44. Two different responses were observed: initial cyclic hardening

up to 2.0 x 104 cycles followed by cyclic softening or only cyclic softening (Fig. 44a)

and initial softening or hardening followed by saturation and subsequent hardening

(Fig. 44b). Supp. Fig. 1 shows the evolution of the εa,p versus N of all tested

specimens.

Fig. 45 depicts a characteristic plot of Ucorr and icorr versus the number of cycles

from CAT. The Ucorr curve exhibits an ascending trend with distinct drops until it

culminates in a maximum just before the failure of the specimen. The high increase in

fresh, non-passivated surface during catastrophic failure leads to a sharp decrease in

Ucorr. The icorr curve maintains a steady level with distinct peaks, corresponding to the

observed drops in Ucorr. This relation between higher peaks in icorr and the more

prominent local drops in the Ucorr could be noticed in all tests (data not shown).

Fig. 44: Evolution of the εa,p versus N describing the cyclic deformation behavior

measured in axial stress-controlled constant-amplitude tests at σa = 210 MPa.

Fig. 45: Characteristic plot of the free corrosion potential and the corrosion current

versus N from axial stress-controlled constant-amplitude tests at σa = 210 MPa.

5.4.3 Characterization of fatigue-induced surface damage in CAT

SEM was used to analyze surface damage caused by fatigue and correlate it

with the observed corrosion current and potential signals. The tests were interrupted

three times when a significant increase in icorr and a significant decrease in Ucorr were

observed. Fig. 46 shows the sections (marked by numbers) of the measured Ucorr and

icorr curves versus N.

Fig. 46: Three sections (marked by numbers) of the measured potential and current

curves versus the number of cycles of one typical specimen. The CAT tests

(σa = 210 MPa) was interrupted when a significant increase in current and decrease in

potential were detected.

The corresponding surface state of the specimen after each interruption is

shown in the SEM micrographs in Fig. 47. After the first significant increase in icorr with

a simultaneous drop of Ucorr at N1 ~ 500,000 cycles (row 1), four cracks approximately

perpendicular to the loading direction are observed in three areas (blue arrows A-D).

After the further increase in icorr and Ucorr decrease at N2 ~ 510,000 cycles (row 2),

crack A has propagated for more about 50 μm (identified as A.1). Crack B also

propagates to the left and right sides, marked as B.1 and B.2. Cracks C and D merge

and propagate further (marked as C.D) on the left side, and a new crack (E) is formed.

All cracks in row 2 present openings up to 2 μm. In row 3, after additional ~200 cycles

(N3 = 509,196), all cracks experience further opening displacement of about 7 μm,

without an extension, indicating that the crack is propagating inwards.

Fig. 47: SEM micrographs of the specimen surface after each interruption in the

corrosion current and potential measurements during CAT tests (σa = 210 MPa).

Fig. 48 illustrates the fracture surface of the specimen whose electrochemical

behavior has been shown in Fig. 46. An overview (Fig. 48a) reveals three distinct

fracture areas. Area 1 corresponds to the initiation of the fatigue crack, and details are

shown in Fig. 48b-d. Notably, the crack initiated internally (Fig. 48b), marked by radial

spreading lines (Fig. 48c). Other internal cracks are seen in this area, as exemplified

in Fig. 48d. Area 2 exhibits river marks that indicate the progression direction of the

fatigue crack. Fatigue striations, which reflect the crack growth per cycle, are visible

within this region (Fig. 48e). The final fracture surface is located in area 3, shown in

greater detail in Fig. 48f. A small final fracture surface reflects rapid and unstable crack

propagation. This area is predominantly ductile, as indicated by its rough appearance,

dimples, and plastically deformed portions (blue arrows in Fig. 48f). This ductile final

fracture surface is characteristic of β-metastable Ti-alloys produced via LPBF-M in the

solution treated and aged condition [155]. The observed dimples are associated with

the presence of αp, attributed to the interfacial cohesion strength between the α and β-

phases, potentially leading to void formation at the α/β interface. Additionally, the

dimples may result from the rupture of micropores and the concurrent destruction of

the surrounding material during plastic deformation [156–158].

The features described above were typically observed on most fracture surfaces

of the tested specimens (see other example in Supp. Fig. 2).

Fig. 48: Fracture surface of the specimen from axial stress-controlled constant-

amplitude tests (σa = 210 MPa), which was subjected to the fatigue-induced surface

damage investigation (compare Fig. 46 and 47). a) Overview showing three different

fracture areas marked by numbers in blue: (1) fatigue crack initiation, (2) fatigue

fracture, and (3) final fracture surface. b-d) Area 1, e) Area 2, and f) Area 3 in more

detail. The blue arrows show internal fatigue crack (d) and plastically deformed portions

(f).

6 DISCUSSION

6.1 Influence of α-precipitate orientation and distribution on the deformation

behavior of Ti-5553 in the ST condition

The nanoindentation results of the ST condition show a strong correlation

between the cyclic deformation data, the microstructure in the volume beneath the

nanoindents, and the fatigue-induced surface morphology and dislocation structure.

The orientation and distribution of αp play a critical role in the cyclic deformation

mechanisms of this β-metastable Ti-alloy, as discussed in detail in the following

subsections.

6.1.1 Microstructure

Through ST heat treatment below the β-transus temperature, an (ɑ+β)-

microstructure was achieved, with small ɑp-particles evenly distributed in the β-matrix,

as to be expected based on results reported for classically manufactured Ti-5553

[159,160]. The αp-phase fraction in the analyzed material is in the lower range of values

reached for the non-AM materials (12% versus 10%[161] to 26%[162]). Further, in

contrast to αp-chevrons[163], singular acicular particles are observed, which may be

due to the fundamentally different processing conditions before the heat treatment.

Over significant regions of each β-grain, the αp acicular precipitates are co-

aligned with their long axes parallel to each other; however, an overall preferred

orientation is neither observed within a single grain nor between grains. Further, the

EBSD measurements do not indicate an overall preferred crystal orientation and they

do not hint at an orientation relationship between the αp-phase and the surrounding β-

matrix. The latter result differs from the observation of others who reported that the αp-

phase has a Burgers orientation relationship with the β-matrix in (α+β)-solution

annealed Ti-5553 [164]. A likely explanation is differences in the processing of

conventional and AM Ti-5553. The conventional process involves thermo-mechanical

treatments before the actual heat treatment. Thus, a lamellar (α+β)-microstructure is

the starting point, while in the present case, it is purely β, which has neither been

plastically deformed nor recrystallized.

6.1.2 Cyclic deformation and creep behavior

Typical changes in hysteresis parameters (Dap, Dmin/Dmax, and ΔDmin-norm) over

the number of cycles are observed. The overall trend of the plastic depth amplitude

(Fig. 31) over the whole test reveals three subsequent stages of hardening, saturation,

and softening. Reports on the behavior of conventionally processed Ti-5553 on the

macro scale differ from this observation, and they are not consistent: only cyclic

softening, cyclic softening followed by saturation, or cyclic hardening followed by

softening were observed [165–167]. Microstructural and compositional differences are

one likely reason for the differences. All the reports refer to conventionally processed

metastable β-alloys, with different compositions and heat treatments, compared to

each other and the investigated alloy. Moreover, classical macro fatigue tests reveal

an average response over the microstructural constituents, whereas the local

interactions of phases with the deformation mechanisms (dislocation formation and

movement, twinning) were examined, without averaging. Thus, the influence of local

structural inhomogeneities on the cyclic deformation behavior is extracted, which also

explains the scatter between different indents (= regions), and the fluctuations seen in

the progression of some hysteresis parameters (see below).

The second parameter that was evaluated, ΔDmin-norm, decreases continuously

and significantly with increasing numbers of cycles, while Dmin/Dmax increases steadily.

Like Dap, ΔDmin-norm exhibits significant fluctuations at the beginning of the tests. Such

fluctuations are not seen for Dmin/Dmax. All curve progressions indicate a decrease in

plastic deformability over the course of loading. The ratio of the minimum displacement

reached in one cycle after unloading from the maximum load (which resulted in Dmax),

Dmin/Dmax, further hints at an overall more elastic unloading behavior with ongoing cyclic

deformation, which correlates with the saturation observed in the progression of Da,p.

The TEM investigations suggest that the interaction of dislocations with αp-

precipitates is the most important cyclic deformation mechanism influencing cyclic

hardening and softening and cyclic creep. The existence and orientation of the αp-

precipitates below and around the indents influence the formation of dislocation

structures. Most nanoindents exhibit alternating hardening and softening with an

overall trend for smaller plastic displacement amplitudes (cyclic hardening) over the

first ten cycles. Based on observations from macro fatigue tests [165,168], one may

hypothesize that the repeated indentation activates multiple slip systems, as well as

interactions of dislocations with each other and with nearby αp-precipitates, leading to

hardening. As for macro specimens, softening may arise from dislocation annihilation

due to mutual dislocation impingement. Such dislocation annihilation has been stated

to be the main reason for the predominance of cyclic softening in (α+β)-Ti-5553 and

(α+β)-Ti-1023 [165,168,169]. Most likely, these processes occur simultaneously under

the localized relatively high loads and the multiaxial stress and strain state in the

confined interaction volume below and around the indent. In the further course of

loading, for 10 ≤ N ≤ 104, only small further changes in the plastic deformation

amplitude are observed, with some indents showing overall hardening, and others

overall softening. In this “nearly-saturation” state, therefore, either the described

hardening or the softening mechanisms are dominant. Further, with ongoing loading,

more dislocations can be activated and interact with differently oriented αp-precipitates

in the indented volume, and the interaction volume expands - to a smaller or greater

extent, depending on the existence and orientation of αp particles nearby. Transversely

placed αp-precipitates block the motion of the dislocations more effectively than αp

whose long axis is oriented orthogonal to the surface, that is, parallel to the indentation

direction (compare Fig. 35c and Fig. 35d). Thus, the interaction volume can expand

more in the latter case, overcoming possible dislocation annihilation and strengthening

the β-matrix. αp orientation has also been reported to be an important factor influencing

the cyclic deformation response of conventionally manufactured Ti-5553 on the

macroscale, by influencing the prevalent micromechanisms [165]. Here, αp-precipitates

in one sample deformed to different strain levels depending on their orientation to the

loading direction.

For N ≥ 104, most nanoindents exhibit hardening, and only a few show softening.

Hardening may be explained by gradual activation and increasing interactions of

multiple slip systems in the αp-precipitates, together with the impingement of α-/β-

phase boundaries to dislocation movement [165]. Softening is likely due to new

dislocation arrangements occurring in αp and in the volume below the indents, thereby

facilitating plastic deformation [165,168]. Additionally, cyclic softening can be

intensified with dislocations relocating from regions of low dislocation density to regions

of high density [170].

Another mechanism is twin formation, shown by β-twins initiated at grain

boundaries during cyclic nanoindentation (see red arrow in Fig. 35a). With increasing

deformation, twin boundaries can progressively form and effectively block dislocation

movement.

Especially during the first ten cycles, relatively large fluctuations in the

progressions of Dap and ΔDmin-norm over N are observed. Such fluctuations are also

seen in the further course of loading for many of the indents, however with considerably

smaller amplitudes. It is especially noteworthy, that ΔDmin-norm occasionally even

acquires negative values. These indicate that the indenter is pushed up, instead of

being pushed to the same or a bigger depth than in the cycle before. A similar behavior

was observed during cyclic nanoindentation of an Al-Si alloy [144]. It may be explained

by the release of residual stresses due to the cyclic plastic deformation. Residual

stresses may result from thermally induced imbalances during LPBF-M, where fast

heating and cooling in nearby regions may lead to quickly changing states of thermal

expansion and contraction [171,172]. An additional cause can be modifications of the

local microstructure due to the β→α phase transformation during the heat treatment

[173]. All these processes can lead to complex residual stress/strain states. Further,

the LPBF-M process and the subsequent heat treatment promote the formation of high

dislocation densities primarily in the β-grain boundaries and in the αp-precipitates

(Supp. Fig. 3). Under repeated loading and unloading cycles, these dislocations can

be released and interact with each other and other dislocations. If this occurs stepwise,

to different amounts in different cycles, fluctuations, that is hardening and softening,

alternating from cycle to cycle or over tens to hundreds of cycles, are seen.

6.1.3 Fatigue-induced indent characteristics

The quantitative observations of indent and pile-up size and morphology

correlate well with the development of the cyclic deformation and creep response.

Larger projected areas (Ap) correlate with greater indent depths (Dmax), suggesting an

overall lower resistance to cyclic plastic deformation in the indented volume. This is

reflected by higher values of Dmin/Dmax and ΔDmin-norm curves and pronounced cyclic

softening for N ≥ 104 (see e.g. nanoindent 1). Correspondingly, the smallest Ap and

Dmax values correlate with cyclic hardening and the lowest values for ΔDmin-norm (see

e.g. nanoindent 16).

The extent of pile-up is strongly influenced by the microstructure surrounding

and below the indents, determining to what extent plastic deformation is hindered in

the volume below the indent. For the investigated material, the values thus depend on

how the precipitates influence the dislocation motion. For example, a high content of

precipitates with differing orientations will effectively restrict the movement of the

dislocations deeper into the material (see e.g. Fig. 35c). Consequently, they cause a

decrease in the amount of plastic deformation [174]. Hence, a strong pile-up along the

flanks of the indent appears as excessive material pushed out at the surface, as also

reported for an Al-Si alloy [144]. However, indents in pure β-regions present large pile-

up sizes (e.g. nanoindent A) as well. In this case, dislocations are more confined to the

surface due to the progressive formation of twins at β-grain (Fig. 35a) and large pile-

up happens above the twin boundary.

Cyclic indentations performed in regions with a relatively high content of widely

spaced αp-precipitates have intermediate sizes and large pile-up volumes (e.g.

nanoindent B, Fig. 35b). Here, precipitates may be located near/on the surface, in the

volume sideways of the indent, or in the volume below the indent. Such precipitates

favor the formation of dislocation structures because the growth of αp during the heat

treatment deforms the β-matrix, causing considerable stress and thus generating

dislocations. During the cyclic indentation, the high dislocation density is released from

the αp-precipitates that were directly encountered by the indenter. These dislocations

interact with each other and with the precipitates, offering resistance to dislocation

motion. However, a preferable placement of the precipitates still gives some space for

dislocation movement into deeper areas in the volume. This led to an intermediate

cyclic plastic deformation state, as exemplified by indent 3 in Fig. 33f.

In comparison, small nanoindents usually arise from cyclic indentation

performed directly on αp-precipitates, which offer a relatively high resistance to

deformation. When encountering one or more precipitates, the indenter cannot

penetrate further into the material. An underlying microstructure with fewer and more

preferably-oriented αp enables dislocation movement without significant restrictions, as

exemplified in Fig. 35d. A large volume of material is plastically deformed and the

dislocations move deeper below the indent, if only a low amount or no αp is present. In

this case, much less material is pushed out at the edges of the indent, resulting in a

very small pile-up size

6.1.4 Summarizing model mechanism

Summarizing the results, the interaction mechanism shown schematically in

Fig. 49 is proposed. Precipitates with a preferred orientation where their long axes are

parallel to the indentation direction are less effective in hindering dislocation slip (Fig.

49a) than precipitates that are not co-aligned (Fig. 49b). The latter arrangement

effectively hinders the expansion of the interaction volume, thus fostering hardening

due to the fast development of a high dislocation density in a confined volume.

Accordingly, a high number of precipitates is more effective than a lower number,

hindering dislocation movement faster, and a low or no content of αp results in more

space for the dislocation mobility, thus higher plastic deformation values and more

cyclic creep (Fig. 49c).

Fig. 49 [151]: Schematic representation of dislocation distribution and structures in the

volume beneath cyclic nanoindents: a) dislocation allocation affected by a low number

of αp-precipitates oriented in a way such that a direct path for the dislocations towards

deeper regions is available; b) high content of differently oriented αp, such that a high

dislocation density develops and dislocations are trapped between the transversely

positioned precipitates; c) indent surrounded only by the homogeneous β-matrix,

yielding regions of high dislocation densities along the grain boundaries. 1 = indent; 2

= grain boundary; 3 = low dislocation density; 4 = high dislocation density; 5 = very low

dislocation density.

6.2 Effect of STA and BASCA heat treatments on the microstructure and

nanomechanical properties of Ti-5553

This section discusses how the microstructure resulting from STA and BASCA

heat treatments affects the nanomechanical properties of AM Ti-5553. For this,

nanoindentation tests were performed to analyze the quasi-static and cyclic properties.

Following is a detailed discussion of the findings.

6.2.1 Microstructure

The “as-built” microstructure (Fig. 21a), composed of only β-phase, results from

the rapid cooling of the molten material above the β-transus temperature during the

LPBF-M process [69,175]. The retained β-phase is attributed to the low diffusion rate

due to the high content of β stabilizers in Ti-5553[176]. In conventionally manufactured,

β-annealed Ti-5553 [159,177–179], athermal ω-phase precipitates have been

observed. However, it was unable to prove these precipitates in the investigated

specimens. Beta-annealing is typically performed at 888 °C, approximately 50 °C

above the β-transus temperature for 150 minutes followed by quick cooling, such as

fan cooling [166]. Considering that the rapid cooling process in LPBF-M creates the

necessary conditions for the formation of these ω-precipitates [180], it is possible that

the microstructure in fact contains a very small amount of ω-phase. One may suspect

that their small size range of 10 to 20 nm [179] and their potential non-uniform

distribution within the material may have hindered their identification through XRD and

EBSD measurements on the analyzed cross-sections [181]. Multiple inverse Fourier

filtered transform (IFFT) or dark-field TEM images could be effective in detecting ω-

phase precipitation [178–180].

Both STA and BASCA microstructures observed in the investigated specimens

are consistent with findings reported for conventionally manufactured Ti-5553 [74]. In

the STA-microstructure, globular αp-phase precipitates were formed and dispersed

within a β-phase matrix after solution annealing in the α+β temperature range below

the β-transus temperature. Although αp-precipitates are attributed to nucleate along

the boundaries of β-grains in conventionally manufactured Ti-5553 [185], any

relationship between these precipitates and the grain boundary in the LPBF-M

manufactured material is not seen. Interestingly, the β-phase along the grain

boundaries appears devoid of αp, possibly due to the accumulation of β-stabilizing

elements in these regions during αp-phase precipitation, a phenomenon described by

Jones et al. [179,178]. Notably, the β-phase contains numerous very fine needle-like

αs-precipitates (Fig. 21d), consistent with prior findings [177]. These αs-precipitates

likely result from the aging step [100]. Regarding the BASCA-microstructure, it may be

assumed that accumulations of the αp-phase initially form along the grain boundaries

of the β-phase during the slow cooling phase to the aging temperature, appearing as

partially discontinuous precipitates of varying width (see Fig. 21e,f). The gradual

cooling process following the stage above β-transus temperature facilitates the growth

of these αp-precipitates due to enhanced diffusivity with prolonged exposure to higher

temperatures. Consequently, the BASCA microstructure contains a higher proportion

of α-phase compared to the STA condition, as evidenced in Fig. 21, Fig. 22, and Fig.

24. The β-phase fringes observed along the α-phase boundary (Fig. 24g) likely

became enriched with β-stabilizing elements during the slow cooling phase [100].

Additionally, fine lamellar αs-laths, uniformly distributed within the β-grains (Fig. 21e,f

and Fig. 24g), may have precipitated at numerous, homogeneously distributed

intragranular nucleation sites during the aging in the α+β field after the annealing stage

above β-transus [177,74].

The elongated shape of the β-grains, aligned with the AM build direction in the

longitudinal view (Fig. 24b,e,h), observed for all three conditions, is attributed to a

phenomenon called columnar grain growth [100]. This happens because adjacent

solidified melt lanes act as starting points for epitaxial grain growth in the direction of

heat flow during solidification. The material undergoes repeated heating and cooling

cycles as each new layer is added, resulting in the merging and elongation of the grains

in the longitudinal direction, extending across multiple layers of the build. This

phenomenon has been discussed in relation to the microstructure of Ti-5553

processed by LPBF-M [175]. This grain growth is more pronounced in the longitudinal

direction because the material spends more time in the elevated temperature zones

during the AM process, leading to a different thermal history compared to the cross-

section.

In the cross-sections (Fig. 23), slight grain growth for STA condition is observed.

This can be explained by possible recrystallization in the temperature range below the

β-transus, as reported for conventionally processed Ti-5553 [186]. Conversely, in the

BASCA condition, a high amount of α-phase may serve as preferred nucleation sites

of new grains during aging and slow-cooling stages [187]. The abundant α-precipitates

promote the formation of numerous smaller β-grains instead of allowing a few larger

grains to develop, resulting in a more fine-grained microstructure.

6.2.2 Quasi-static mechanical properties

The load-depth curves derived from quasi-static nanoindentation tests (Fig.

36a) provide valuable insights into the plasticity of the "as-built," BASCA, and STA

conditions. The observed plasticity variations may be attributed to the distinct

microstructural characteristics and responses to plastic deformation within these

conditions. In the case of the "as-built" condition, it is evident that the microstructure

composed solely of the bcc β-phase enables higher plasticity. The bcc structure offers

slip systems for dislocation motion, facilitating plastic deformation [188,189].

Additionally, the lower stacking fault energy (SFE) in the β-phase promotes easier

dislocation movement through the crystal lattice, contributing to increased plasticity

[188,190]. This is corroborated by the 3D surface plots and SPM topography images

in Fig. 37. The indents on the "as-built" specimen are larger and exhibit greater depth

compared to BASCA and STA specimens, signifying enhanced plastic deformation

(compare Fig. 37a). In contrast, the presence of hcp α-precipitates may lead to

changes in the local atomic arrangement and affect the SFE of the β-phase. The α-

precipitates can alter the SFE by creating obstacles or barriers for the movement of

dislocations, making it more difficult for them to traverse the crystal lattice. This results

in lower plasticity in BASCA and STA conditions compared to the "as-built" state. The

specific impact on plasticity depends on factors such as the size, distribution, and

volume fraction of the α-precipitates. Notably, the STA condition exhibited the lowest

plasticity and highest nanohardness, attributed to the presence of globular αp and

extremely fine, needle-like αs-precipitates in the microstructure. These αs-precipitates

significantly affect dislocation motion and material ductility [191], resulting in decreased

plasticity. This is validated by the smallest values of Ap, Ap-u, and Vp-u found for the STA

specimens (Fig. 37b), against the intermediate Ap-u and Vp-u values of the indent on the

BASCA specimens (Fig. 37c). The observed crystallographic texture and α-phase

colonies with preferred orientations of STA condition (Fig. 24c) may also contribute to

plasticity reduction. Although the orientation of the indented β-grain alone does not

significantly influence the load vs. displacement nanoindentation curves in Ti-5553

[134], the crystallographic orientation of the β-grains has a major effect on slip initiation

in the α-phase by generating a heterogeneous stress field in relation with the β-grain

orientation [192]. Crystal plasticity finite element (CPFE) simulations have shown that

high strengths are obtained for loading direction parallel or nearly parallel to the c-axis

for Ti-alloys with bimodal microstructure [60,193–195]. Experimental and CPFE

simulated nanoindentation on the β-metastable Ti–7Mo–3Nb–3Cr–3Al (Ti-7333, wt.%)

alloy with the volume fraction of αp-precipitates about 20% and β matrix with finely

dispersed αs [193], very similar to the presented STA-microstructure, reveals that the

accumulation of shear strain during indentation process can be affected by αp

orientations. During indenting deformation, the percentage of basal slip decreased,

while the prismatic slip increased to be the dominant slip activity in αp. Based on that,

it may be assumed that the slip activity is difficult in the observed β-grains oriented with

a (001) direction aligned with the loading axis together with the prismatic slip of the αp

colonies with preferred orientations. No further information regarding the slip activity in

αs was found.

The mechanical properties (Fig. 36b,c,d) further emphasize the influence of the

microstructure on the response of the material to the applied force. The hardness

reflects the resistance to plastic deformation, which is clearly influenced by the volume

fraction and morphology of the α-precipitates in the β-matrix. The highest H exhibited

by the STA condition (Fig. 36c) may be attributed to the interaction between the fine

αs with the coarse αp, which offers more obstacles to plastic deformation. Previous

studies support this [196,197]. The increase in H parallels the macroscopic hardness

response observed in conventionally manufactured Ti-5553, where the growth of αp

and the formation of αs within the β-matrix during aging were identified as factors

contributing to a Vickers hardness increase of over 50% in Ti-5553 following STA

[159,198,199].

The increase in Er (Fig. 36b) and in S (Fig. 36d) after the heat treatments

corroborate the report on conventionally manufactured Ti-5553 evaluated through

nanoindentation: the highest E was determined for the (α+β)-microstructure, while the

pure β-phase exhibited the lowest E [200]. The highest values of E and S for the

BASCA condition in relation to STA condition may be explained by the coarser α-

precipitates, which have a less elastic mismatch with the β-phase matrix compared to

finer precipitates [201].

The BASCA condition shows increased data variability in Er and H, as

evidenced by a larger IQR and extended whiskers, as well as more outliers for Er,

compared to the “as-built” and STA conditions. This variability may be attributed to the

indents made across the microstructure with α-precipitates differing in their size and

thickness. The indenter met sometimes portions with only αp, portions, with only αs-

laths, and sometimes portions comprising both αp and αs-laths precipitates.

6.2.3 Nanofatigue behavior

The hysteresis parameters Dap, Dmin/Dmax, and ΔDmin-norm over the number of

cycles were evaluated to assess the cyclic deformation behavior of Ti-5553 with

different microstructures. The Dap (Fig. 40) reveals hardening, saturation, and

softening stages for “as-built” condition, while STA and BASCA conditions show an

overall cyclic hardening, both exhibiting very close values. These results differ from the

macroscale cyclic deformation behavior of Ti-5553. Cyclic softening followed by

saturation was reported for Ti-5553 with a microstructure consisting solely of β-phase

[166], and only cyclic softening behavior for Ti-5553 with microstructures similar to that

of the analyzed material, resulting from STA and BASCA heat treatments [165,168]. It

is important to note that the existing literature on the macroscale focuses on

conventionally processed Ti-5553, studied using traditional fatigue tests that provide

an overall assessment of the behavior of the material across its microstructural

constituents, while the NI approach examines the specific interactions of phases with

the deformation mechanisms. Through NI, the influence of local structural

inhomogeneities on the cyclic deformation behavior is extracted, which also explains

the scatter between different indents (Fig. 41). By probing the local interplay between

phases, it is elucidated how the phases and the resulting dislocation interactions

contribute to the cyclic deformation behavior of different microstructures. This

nanoscaled approach hence offers specific insights that might not be apparent in

traditional macroscale fatigue testing.

Based on observations from macrofatigue tests on Ti-5553 alloy [165,166,168],

it may be hypothesized that the hardening behavior observed in the STA and BASCA

conditions for N ≤ 10 and 102 ≤ N ≤ 104 is due to interactions between dislocations and

of dislocations with α-precipitates [165,168], leading to the obstruction of slip at α/β

interfaces in the indented volume. As for macro-specimens, the softening effect in the

STA condition (10 ≤ N ≤ 102 ) can be attributed to multiple slip activities within αp, along

with slip initiation in the β-phase, and eventually also in the transformed β-matrix with

αs [168]. Additionally, slip transfer from the β-phase to the αs can occur. For the BASCA

microstructure, the softening may be caused by slip activities that not only occur in the

β-phase, but also initiate in the αs-laths, and slip transfer from the β-matrix to the αs-

laths. The “nearly-saturation” state of STA and BASCA conditions for N ≥ 104 likely

arises from the balance between the counteracting effects of the softening and

hardening mechanisms. The curve progression at lower values for BASCA may be

explained by a greater restriction of dislocation movement in the β-matrix due to the

higher amount of α-precipitates (compare Fig. 40d). Hence, the interaction volume

during the cyclic nanoindentation can expand more in the BASCA case than in the

STA. For the “as-built” condition, instead of interactions of dislocations with α-

precipitates, the cyclic hardening may arise from repeated indentation that activates

multiple slip systems in the first ten cycles, as well as dislocations interactions with

twin-structures[166,202] (N ≥ 104). The progressive formation of twin boundaries with

increasing deformation has been shown to effectively block the dislocation movement

in Ti-5553 with only β-phase [166]. The cyclic softening (10 ≤ N ≤ 102; 103 ≤ N ≤ 104 )

may be attributed to the dislocation annihilation resulting in mutual dislocation

impingement [165,166,168,169]. Another softening mechanism may be detwinning

and twin boundary degradation [166]. Considering observations on the macroscale, it

may be assumed that saturation takes place for only β-phase microstructure due to the

flip-flop movement of dislocation dipoles [166].

Upon individually analyzing the behavior of indents (Fig. 41), it is surprising that

the presumably most uniform, homogeneous β-microstructure of the “as-built”

condition shows a greater divergence in the curves and more pronounced hardening-

softening fluctuations compared to the heat-treated conditions (Fig. 41a). A cause for

that could be the inherent heterogeneity introduced by varied β-grain orientations (Fig.

24a). The curves for STA exhibit similar courses, differing in their extent, without strong

fluctuations as observed in the “as-built” condition. Although the indents were made in

surface regions always with αp-phase within β+αs, the STA microstructure comprises

regions with preferred oriented αp-precipitates (Fig. 24d). Hence, depending on the αp

orientation in the indented volume (=differences in appearance of the exposed

sectioned plane), variability in the individual curves is expected. However, this

variability is likely reduced by the texture in STA (Fig. 24c). Grains with a preferred

orientation may have more active slip systems aligned with the indentation direction.

This alignment can result in a more uniform activation of slip systems during cyclic

deformation, leading to less variability in the Dap curves compared to “as-built”

condition (Fig. 41b). The BASCA condition, in turn, shows curves that deviate less

from each other in terms of their courses (see Fig. 41c). Despite this condition also

presenting regions with α-precipitates preferred oriented, the coarser precipitates in a

high amount may result in a consistent pattern of obstacles for dislocations, making

the material less sensitive to localized variations in microstructure.

The parameter ΔDmin-norm decreases continuously and significantly with

increasing numbers of cycles for all three conditions. Like Dap, ΔDmin-norm indicate a

decrease in plastic deformability over the course of loading. The ratio of the minimum

depth reached in one cycle after unloading from the maximum load (which resulted in

Dmax), Dmin/Dmax, increases steadily, revealing an overall more elastic unloading

behavior with ongoing cyclic deformation. Indents made on the untreated condition

(“as-built”) with only β-phase present high values of projected areas, pile-up sizes and

maximum indentation depths, after fatigue loading (N = 105) (compare Fig. 42a, Table

10), which also correlate with ΔDmin-norm curve course at the highest values. The

absence of harder α-precipitates allows for greater elasticity, resulting in a higher Dmax

during loading cycles. Consequently, the ratio Dmin/Dmax is lower, indicating that the

material experiences a more significant elastic recovery during unloading from the

maximum load. For the heat-treated conditions, one may suggest that the morphology

and amount of the α-precipitates in the microstructure strongly influence the dislocation

formation and movement, and therefore the plastic deformation. During heat treatment,

the growth of the α-phase may alter the shape of the β-matrix, leading to significant

stress and triggering the generation of dislocations. When cyclic indentation occurs,

the direct contact of the indenter with the α-precipitates may result in the release of a

large number of dislocations. These dislocations then interact with each other and with

the precipitates, creating resistance to dislocation motion, suggesting an overall higher

resistance to cyclic plastic deformation in the indented volume. This is reflected by the

lowest values for ΔDmin-norm for STA and BASCA conditions (see Fig. 39). The lowest

values of ΔDmin-norm for the BASCA condition may be attributed to the high amount of

αp-precipitates (Fig. 23), which introduces constraints and obstacles to dislocation

movement, leading to a reduced deformation capacity and increased resistance to

plasticity. As a result, the BASCA condition experiences a lower Dmax during loading

cycles compared to “as-built” condition. Hence, the ratio Dmin/Dmax in BASCA condition

is higher, reflecting a relatively less elastic recovery during unloading due to the high

amount of precipitates that hinder dislocation motion. In the BASCA microstructure,

cyclic indentations revealed slightly bigger pile-up volumes compared to STA (Fig.

42c). This difference can be attributed to the relatively higher content of α-precipitates

in the BASCA condition. The presence of many α/β interfaces provides sites where

dislocations can be trapped on the surface, resulting in more material being pushed

out at the edges of the indent than that seen in STA. However, the larger spacing

between the αs-precipitates in BASCA still allows some dislocation movement into

deeper areas in the volume, leading to a slightly higher Dmax. During cyclic indentations

performed on the STA microstructure, the indenter is hindered to move further into the

material when it encounters both αp and closely spaced, fine dispersed αs (Fig. 21d).

The plastic deformation may be restricted to the αp-phase, with the hard β+αs staying

in its elastic domain [203]. Hence, it can be assumed that the movement of dislocations

within the β+αs is minimal because of the abundance of α/β interfaces that impede

dislocation gliding. In this case, the lowest Dmax is observed (Table 10). Dmin/Dmax

values for STA condition are even smaller than that of “as-built” condition. The

combination of different precipitate morphologies and sizes (acicular αp and very fine

αs) contributes to a more effective hindrance of dislocation motion, and dislocations

generated around the αs may lead to a reduction in elastic recovery.

6.3 Cyclic deformation behavior of Ti-5553 in quasi-physiological medium

Both LIT and CAT tests reveal different progressions of εa,p between the tests

(Fig. 43, Supp. Fig. 1). In LIT tests, the curves show scatter, presenting different levels

of εa,p. Both hardening and softening occur in the CAT tests, however in a different

order. These different behaviors may be attributed to the simultaneous occurrence of

plastic deformation and microcrack formation/growth, and their interplay determining

whether plastic deformation or crack formation and growth prevail. When flaws are

abundant, microcracks form early and spread swiftly, preventing sufficient time for

plastic deformation. Conversely, in cases with fewer flaws, microcracks emerge later,

allowing the plastic deformation to determine the cyclic behavior. The reports on the

cyclic deformation behavior of conventionally manufactured Ti-5553 alloy with a

microstructure very similar to that of the investigated material, featuring dispersed

small αp within the grains of the retained β-matrix with numerous very fine αs, also

show distinct εa,p responses [165,168]. It is important to note that the tests in these

reports were conducted in strain-controlled mode, while load-controlled tests were

performed within the scope of the present work. Strain and load-controlled tests are

comparable as long as only plastic deformation occurs; as soon as a crack forms, the

behavior diverges: in strain-controlled tests, the specimen then experiences lower

loads, while the real stress increases in load-controlled tests [204]. Although the

similarities in microstructure suggest that εa,p responses may be comparable between

conventionally and additively manufactured specimens, the unique characteristics

introduced by the AM process may result in differences. It is crucial to account for the

potential impact of layer interfaces, porosity, and other defects inherent to AM when

interpreting cyclic deformation behavior in both load-controlled and strain-controlled

modes.

The initial and final cyclic hardening observed in the tests may happen due to

dislocation accumulation, the interaction between dislocations, and the interaction

between dislocations and the α and β phases. Very fine αs-precipitates placed

transversely to each other (Fig. 41c) significantly affect the dislocation motion [205],

contributing to cyclic hardening. Moreover, simultaneous microcrack formation and

coalescence, as evidenced in Ti-alloys[206], can limit plastic deformation during cyclic

loading, resulting in hardening. In turn, cyclic softening is influenced by various slip

activities within αp, followed by slip in the β-phase, and subsequent slip in the

transformed β-matrix containing αs. Transfer of slip from the β-phase to αs is also

possible [165]. Furthermore, the flaws resulting from the LPBF-M process (Fig. 15,

Supp. Fig. 2d,e) introduce material discontinuities that cause localized stress

amplification. This facilitates dislocation movement by providing a source or a sink for

dislocations, or by reducing the resistance of the surrounding microstructure to slip

[207]. As a result, plastic deformation increases, leading to cyclic softening. The

saturation behavior emerges from the balance between the counteracting effects of

the softening and hardening mechanisms. Based on these observations, it may be

suggested that crack propagation is accelerated in specimens with a high amount of

flaws, resulting in initial cyclic hardening followed by cyclic softening or only cyclic

softening (Fig. 44a), while specimens containing few flaws present initial softening or

hardening followed by saturation and subsequent hardening (Fig. 44b).

The plastic strain amplitude progresses at very small values over the number of

cycles. This can be attributed to the high strength of the alloy as a result of the (α+β)-

microstructure, as the precipitation of αp and αs within the β-matrix is the main

hardening mechanism in the Ti-5553 alloy [162,208]. The α-precipitates contribute to

less plastic deformation during the cyclic loading process[209]. Furthermore, it is

important to consider that εa,p is also affected by the crack closure effect [210] (Fig. 47,

row 2). As the crack propagates, it creates a cyclic opening and closing of the crack

faces during each loading cycle. This crack closure phenomenon happens because

the material surrounding the crack resists further opening due to the residual

compressive stress state developing at the crack faces and ahead of the crack tip,

effectively reducing the stress concentration at the crack tip and leading to a reduction

in plastic strain amplitude. Another aspect to be considered is that the plastic strain

amplitude evolution may be influenced by the shot-peening process. The residual

compressive stresses introduced by shot peening may increase the yield strength of

the material, which can improve its resistance to plastic deformation [211,212]. The

strengthening of Ti-alloys may also be attributed to uneven distribution of alloying

elements. Certain elements, such as Fe, Mo, V, Ni, Cr, Nb, and Co are known to

enhance the mechanical strength of Ti-alloys [213]. Notably, elevated levels of Mo and

Fe were observed (compare Table 6 and Fig. 42) in the material under analysis. Mo

and Fe, as a β-stabilizer, have been previously associated with increased quasi-static

and cyclic strength in β-metastable Ti-alloys [214–221]. Additionally, the high strength

of (α+β)-Ti-alloys resulted in very small plastic strain amplitudes during CAT tests

[121,222]. Based on these observations, it may be suggested that the increased levels

of Mo and Fe contribute to the increase in strength in the investigated material,

consequently resulting in the observed minimal plastic strain amplitude.

6.3.1 Fatigue-induced surface damage

The identification of surface modifications caused by fatigue was achieved using

an open circuit setup that enables the simultaneous detection of the free potential and

corrosion current. The Ucorr potential represents the electrical potential difference

between the WE surface and RE. A more negative (lower) Ucorr indicates a loss of local

passivity of the oxide layer on the surface of the specimen, as it suggests that the WE

is more prone to losing electrons and undergoing oxidation. As Ti-alloys rapidly restore

this protective oxide film on the surface after it has been damaged, the Ucorr curve

increases again after each drop. Simultaneously, the icorr quantifies electron flow

between the WE and CE surfaces. Peaks in icorr indicate active corrosion, where the

oxide layer deteriorates and momentary oxidation occurs until repassivation ensues,

reducing the ion flux into the solution and restoring the icorr to a stable level. Hence, the

changes in overall potential and current over time depend on both activation and

repassivation rates.

Contrary to Leinenbach et al. [130], the SEM analysis could not distinctly detect

early-stage surface modifications near the suspected fatigue crack origin. The

investigated specimens have a rough surface from shot-peening (see Supp. Fig. 4),

complicating crack origin analysis compared to polished surfaces. The distinction

between cracks and edges or boundaries of the shot-peening depressions is

challenging due to the contrasting appearance in SEM images. A similar issue with

rough surfaces was also encountered in Ti-64 [125]. In the present study, crack

identification and propagation became clear only after 105 cycles. Before that,

identifying the crack leading to the final failure was difficult, as the electrochemical

measurement technique detected multiple crack initiation sites. Thus, a

comprehensive surface investigation was needed to identify the specific crack

requiring monitoring of its progression.

By comparing the findings of the microscopic examination with the results of the

electrochemical measurements, the fatigue crack growth, once the crack reaches the

surface, can be characterized into 3 stages. In the first stage, a specified length of the

crack is created after a considerable number of fatigue cycles (see crack A, row 1, in

Fig. 47). The crack propagation velocity is relatively low during this phase, as

evidenced by the small icorr peaks observed in Fig. 46 (marking 1), indicating that only

a small amount of the surface of the specimen is activated during each cycle. The

second stage includes cracks and microcracks (or incipient cracks) also opening and

closing, indicating a roughness-induced crack closing effect (see row 2 in Fig. 47). This

phenomenon is expected given that the surface roughness of the specimens under

analysis [197] is of the same order as the crack opening displacement [210,223]. The

crack closure effect derives from the misfit of the rough fracture surfaces of the upper

and lower parts of the crack. These misaligned edges engage during fatigue loading,

retarding crack propagation. In the axial corrosion-fatigue testing, however, the quasi-

physiological environment influences the contact pressure and friction between crack

faces. Consequently, the crack closure decreases, and crack growth accelerates[224–

227], which is evident by a sudden rise in the icorr peak and a rapid decrease in the

Ucorr seen in Fig. 46 (marking 2). The third and final stage consists of further cracks

opening (row 3 in Fig. 47, marking 3 in Fig. 46). In this stage, an even faster crack

propagation takes place, ultimately leading to a final fracture. The final failure is hence

marked by the significant increase of the activated surface during the last cycle, leading

to a pronounced peak in the icorr and a corresponding decrease in the Ucorr (see Fig.

45).

The typical fracture surface features shown in Fig. 48 illustrate the fatigue crack

development that led to the stages discussed above. The crack nucleation happens in

Area 1 (Fig. 48a-d). Internal cracks and microcracks form in stress concentration sites,

such as in spherical pores with concentric ridges observed in this area and in area 3

in most specimens (see Supp. Fig. 3). This defect type, previously observed in Ti-64

[228], occurs when incomplete re-melting happens on a specific area of a prior surface

of the layer due to laser scanning. This incomplete re-melting process leads to the

formation of ridges, which indicate solidification boundaries and create voids that the

molten metal cannot fill. Importantly, the presence of pores with ridges was noted in

LPBF-M-manufactured specimens even before fatigue testing[229,230]. This indicates

that these pores are intrinsic to the LPBF-M process itself and are not influenced by

the fatigue testing. As cracks become larger in the second stage, their propagation

becomes more pronounced, and the appearance of fatigue striations is observed in

area 2 (Fig. 48a,e). These striations are generated by microscopic plasticity resulting

from blunting and re-sharpening of the crack tip during load cycles. As the fatigue

testing progresses, the microcracks continue to propagate, and merge with other

microcracks, eventually reaching the surface and becoming visible as macroscopic

cracks. The rough appearance of the final fracture zone from area 3 (Fig. 48f) is given

by the rapid crack propagation, yet with a certain time for plastic deformation because

of the low-frequency loading. This ductile final fracture surface is characteristic of β-

metastable Ti-alloys produced via LPBF-M in the solution treated and aged condition

[155]. The observed dimples are associated with the presence of αp, attributed to the

interfacial cohesion strength between the α and β-phases, potentially leading to void

formation at the α/β interface. Additionally, the dimples may result from the rupture of

micropores and the concurrent destruction of the surrounding material during plastic

deformation [156–158].

7 SUMMARY AND ADDRESSED OPEN QUESTIONS

As reviewed in Chapter 2, the considerable difference in the deformation

capability of the different phases in the metastable β-Ti alloy Ti-5553 may induce very

heterogeneous deformation under cyclic loading conditions on the macroscale. This

leads to complex relationships between phase morphology and distribution and cyclic

deformation behavior. To elucidate such influences, cyclic nanoindentation tests were

performed on this alloy in the ST state up to a maximum cycle number of 105, yielding

information on the cyclic deformation behavior in loading regimes that are relevant to

many applications. By combining cyclic nanoindentation and high-resolution electron

microscopy, the influence of different α-phase orientations and distributions within the

β-grains on the deformation mechanisms under cyclic loading was unraveled. The αp

plays an important role in the local fatigue behavior: i) Dislocation-based deformation

mechanisms are the main origin of cyclic softening, cyclic hardening, and creep

processes during cyclic nanoindentation. ii) A strong correlation between the

microstructure in the volume beneath the nanoindents and the dislocation reaction was

observed: αp-phase orientation and distribution within the β-grains significantly

contribute to the effectiveness of the precipitates as barriers to dislocation motion. High

density, gathering, and trapping of dislocations were observed at α/β-interfaces. iii)

Pile-up occurrence and size are determined by the local plastic deformability, which in

turn is significantly influenced by the presence and orientation of αp-precipitates.

Although STA and BASCA heat treatments are commonly employed to improve

the strength of AM Ti-5553, no research on the local cyclic deformation response of Ti-

5553 across different microstructures resulting from heat treatments has been

reported. Hence, quasi-static and cyclic nanoindentation tests were performed to

comparatively investigate the influence of different microstructures on the mechanical

properties of AM-Ti-5553 in the STA and BASCA conditions. The results clearly show

that microstructural variations notably influence hardness, reduced elastic modulus,

and plasticity of Ti-5553. The plasticity depends on the interplay of size, distribution,

and volume fraction of the α-precipitates. The "as-built" condition displays increased

plasticity owing to favorable slip systems and a low SFE of its solely bcc β-phase.

Conversely, the presence of hcp α-precipitates may lead to changes in the local atomic

arrangement and affect the SFE of the β-phase, creating obstacles or barriers for the

movement of dislocations. This explains the observed lower plasticity in BASCA and

STA conditions compared to the "as-built" state. The STA condition exhibits the lowest

plasticity and highest nanohardness likely due to the dislocation motion hindrance

caused by its microstructural features. Furthermore, the increase in reduced elastic

modulus and nanohardness for the BASCA condition, in comparison to STA, may be

explained by the presence of coarser α-precipitates with less elastic mismatch. The

cyclic behavior analysis reveals hardening, saturation, and softening stages for “as-

built” condition, while STA and BASCA conditions show an overall cyclic hardening,

both exhibiting very close values. A decrease in plastic deformability over the course

of loading (ΔDmin-norm), and an overall more elastic unloading behavior with ongoing

cyclic deformation (=increase in Dmin/Dmax) were observed for all three conditions, with

their extent varying depending on the microstructure.

On the macro-scale perspective, most of the reviewed studies focus on the

application of Ti-5553 in the aerospace industry and explore its ease processability

through AM. However, a few studies have shown that the Ti-5553 alloy has potentially

suitable mechanical properties for orthopedic applications. Although clinical data on

conventionally or additively manufactured Ti-5553 as implant material was not found,

limited literature suggests that AM-Ti-5553 exhibits favorable osseointegration due to

its biocompatibility and surface properties. Additionally, this alloy exhibits high strength

and easily controlled microstructure achieved by heat treatments. Heat treatments may

increase its strength to values comparable to (α+β)-Ti-alloys without sacrificing the low

stiffness achieved by β-stabilizers alloying, which can also be beneficial for implant

applications. No research was reported on the corrosion-fatigue behavior of Ti-5553,

neither conventionally nor additively manufactured. Considering this, this thesis

investigates the cyclic deformation behavior of AM Ti-5553 in HBSS. Distinct numbers

of cycles to failure observed among specimens subjected to the same stress amplitude

were attributed to unevenly distributed flaws introduced during the LPBF-M process.

The cyclic deformation behavior of Ti-5553 in a quasi-physiological environment

exhibits two different responses: initial cyclic hardening followed by cyclic softening or

only cyclic softening, and initial cyclic softening or cyclic hardening followed by

saturation and subsequent cyclic hardening. The cyclic deformation was influenced by

the (α+β)-microstructure and crack propagation driven by the amount of flaws. Both

plastic deformation and microcrack formation/growth happen simultaneously, with their

interplay determining whether there is more plastic deformation or a stronger cracking

effect, resulting in varying levels of εa,p across the tests. Additionally, the shot-peening

post-processing and the high content of Mo and Fe promoted the increase in the

strength of the material, resulting in very small values of plastic strain amplitude during

cyclic loading. Finally, previous studies have demonstrated that surface damage can

be detected in Ti-alloys undergoing fatigue loading in physiological media by

measuring changes in the free corrosion potential and/or current. However, all the

reported measurements were on Ti-alloys with polished surfaces. In this context, the

electrochemical measurements performed within the context of this work demonstrated

high sensitivity in detecting damage caused by fatigue on a rough surface. This is

particularly noteworthy as rough surfaces are known to enhance effective

osseointegration, making them advantageous for implants. The microscopic

examination, the results of the electrochemical measurements, and fracture surface

features illustrate the fatigue crack development.

100

8 CONCLUSION AND OUTLOOK

This thesis highlights the potential of cyclic nanoindentation in unraveling the

influence of local microstructures on the cyclic deformation mechanisms within two-

phase alloys, such as the novel implant alloy investigated. Notably, the nano-scale

results diverge from those obtained through macro-fatigue tests on (α+β)-Ti-5553 alloy.

The testing methodology considers the influence of local structural inhomogeneities on

cyclic deformation. Therefore, the nano-scale approach provides unique insights by

considering the impact of these local structural variations during cyclic deformation. By

probing the local interplay between phases, it was possible to elucidate how the phases

and the resulting dislocation interactions contribute to the cyclic deformation behavior

of different microstructures. At the macro-scale, the cyclic deformation in simulated

physiological media offers insights into the response of AM Ti-5553 alloy to mechanical

stress and corrosion, offering insights into its performance as implant material. Further,

the correlation of fatigue-induced surface damage with corrosion potential and current

measurements proved to be also sensitive to indicate surface damage in situ for

implant materials with a rough surface, which generally promotes better

osseointegration.

In light of the findings presented in this study, some promising avenues for future

research emerge. The study of plasticity in relation to stress distribution under the

indents in β-metastable Ti-alloys such as Ti-5553, where the β grain size is in the

millimeter range and αs is a few tens of nanometers thick, is a significant challenge.

Indeed, these different length scales between phases have to be carefully considered

in further research. Additionally, comparative studies between AM and conventionally

manufactured Ti-5553 under similar corrosion-fatigue conditions can provide insights

into their macroscopic divergent behavior and guide process selection for specific

applications. This different behavior is also expected to be seen in cyclic

nanoindentation on Ti-5553 due to the distinctive microstructural characteristics

introduced by each manufacturing process. The AM Ti-5553 is expected to exhibit

unique features arising from layer-by-layer deposition during the manufacturing

process, leading to specific grain orientations, defects, and microstructural variations.

On the other hand, conventionally manufactured Ti-5553 might possess different grain

structures and defect distributions. These inherent differences in microstructure due to

different processing approaches are likely to result in varied responses during cyclic

101

nanoindentation. Furthermore, the development and validation of multiscale models

that integrate microstructure, flaws, and corrosion effects would enhance predictive

capabilities for Ti-5553 performance under various loading and environmental

conditions.

In conclusion, the gained deeper understanding of the mechanical performance

across different microstructures and length scales of this promising β-metastable

Ti-alloy helps establish a foundation for predicting and optimizing its fatigue

characteristics and performance. This study not only provides a pathway for advancing

fundamental knowledge about AM-Ti-5553 but also offers insights into its practical

applications, particularly in the context of implant materials.

102

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10 SUPPORTING INFORMATION

Supp. Fig. 1: Evolution of the εa,p versus N describing the cyclic deformation behavior

of all tested specimens measured in constant-amplitude tests performed at f = 5 Hz

with a load ratio of R = -1 at σa of 210 MPa.

122

Supp. Fig. 2: Typical fracture surface from axial stress-controlled constant-amplitude

tests. a) Overview showing three different fracture areas marked by numbers in blue:

(1) fatigue crack initiation, (2) fatigue fracture, and (3) final fracture surface. b) Area 1

and c) area 3 in more detail, the latter with plastically deformed portions marked by

blue arrows. d) Internal fatigue microcrack (indicated by blue arrows) originated from

a pore with ridges (see yellow arrow), which was found inside all areas. e) Pore with

ridges in more detail. f) Fatigue striations inside area 2.

123

Supp. Fig. 3: Bright-field TEM micrograph of a non-indented region with high

dislocation densities preferentially formed at the β-grain boundaries and in the αp-

precipitates during the LPBF-M and the heat treatment processes.

Supp. Fig. 4: Specimen surface exhibiting visible roughness resulting from the shot-

peening process after the LPBF-M.